JP4329318B2 - Rare earth sintered magnet and manufacturing method thereof - Google Patents

Description

[0001]
BACKGROUND OF THE INVENTION
The present invention relates to an R-Fe-B rare earth magnet, an alloy powder for the magnet, and a method for producing the same. Here, R is a rare earth element, Fe is iron, and B is boron. In this specification, “rare earth” refers to a group obtained by adding a lanthanoid to scandium (Sc) and yttrium (Y) in the group 3A of the periodic table. This is a general term for elements, and the rare earth element R is at least one element contained in the rare earth element.
[0002]
[Prior art]
The rare earth sintered magnet is manufactured through a sintering step and an aging heat treatment step after press-molding an alloy powder formed by pulverizing a rare earth magnet alloy (raw material alloy). Currently, two types of sintered rare earth magnets are widely used in each field: samarium / cobalt magnets and rare earth / iron / boron magnets. Among them, rare earth / iron / boron magnets (hereinafter referred to as “R—Fe—B rare earth magnets”) have the highest magnetic energy product among various permanent magnets and are relatively inexpensive. Actively adopted in various electronic devices. R-Fe-B rare earth magnets are R ₂ Fe ₁₄ It has a main phase composed of a compound having a B-type crystal structure and a grain boundary phase located at the grain boundary portion of this main phase. Part of Fe may be substituted with a transition metal element such as Co, and part of B may be substituted with C (carbon). For simplification, a magnet in which such element substitution is partially performed is also referred to as an “R—Fe—B rare earth magnet” in this specification.
[0003]
The raw material alloy powder for R-Fe-B rare earth magnets may be produced by a method including a first pulverization step for coarsely pulverizing the raw material alloy and a second pulverization step for finely pulverizing the raw material alloy. . In that case, in the first pulverization step, the hydrogen storage phenomenon is used to embrittle the raw material alloy, for example, coarsely pulverized to a size of several hundred μm or less, and then in the second pulverization step, the coarsely pulverized alloy (coarse pulverization). The powder is finely pulverized to a size having an average particle diameter of about several μm by a jet mill apparatus or the like. Such a grinding method is disclosed in Patent Document 1.
[0004]
There are roughly two types of methods for producing the raw material alloy itself. The first method is an ingot casting method in which a molten raw material alloy is placed in a mold and cooled relatively slowly. The second method is a strip casting method in which a molten alloy is rapidly cooled by bringing it into contact with a single roll, a twin roll, a rotating disk, or a rotating cylindrical mold, and a solidified alloy thinner than an ingot alloy is produced from the molten alloy. And a rapid cooling method represented by centrifugal casting. The strip casting method is described in Patent Document 2, for example.
[0005]
In the case of such a rapid cooling method, the average cooling rate of the molten alloy is 10 ² ℃ / second or more 2 × 10 ^Four It is in the range of ° C / second or less. Therefore, in this specification, the molten alloy is 10 ² ℃ / second or more 2 × 10 ^Four An alloy prepared by quenching at an average cooling rate of 0 ° C./second or less is referred to as a “quenched alloy”. The thickness of the quenched alloy produced by a quenching method such as a strip cast method is in the range of 0.03 mm to 10 mm. The molten alloy solidifies from the contact surface (roll contact surface) of the cooling roll, and crystals grow in a columnar shape from the roll contact surface in the thickness direction. As a result, the quenched alloy has a short axis size of 0.1 μm to 100 μm and a long axis size of 5 μm to 500 μm. ₂ Fe ₁₄ Crystal phase having a B-type crystal structure (hereinafter referred to as “R ₂ Fe ₁₄ May be abbreviated as “B-type crystal phase”), R ₂ Fe ₁₄ It has a fine crystal structure containing an R-rich phase (a phase in which the concentration of the rare earth element R is relatively high) present in a dispersed manner in the grain boundary portion of the B crystal phase. The R-rich phase is a nonmagnetic phase in which the concentration of the rare earth element R is relatively high, and its thickness (corresponding to the width of the grain boundary phase in an arbitrary cross section) is 10 μm or less.
[0006]
The quenched alloy is cooled in a relatively short time compared to an alloy (ingot alloy) produced by a conventional ingot casting method (die casting method). Is small. Further, since the crystal grains are finely dispersed and the area of the grain boundary is wide, and the R-rich phase spreads thinly in the grain boundary, the R-rich phase is excellent in dispersibility.
[0007]
After such a quenched alloy is pulverized by the above-described method, the powder is compression-molded by a press device to produce a powder compact. By sintering this compact, an R—Fe—B rare earth magnet can be obtained.
[0008]
The powder compact of the R—Fe—B based quenched alloy is sintered at a temperature of 1000 to 1100 ° C., for example. The sintering process is promoted by changing the R-rich phase having a relatively low melting point contained in the powder into a liquid phase. After the sintered body is formed by the sintering in a high temperature region of 1000 to 1100 ° C., the sintered body is subsequently cooled. At this time, in order to improve the characteristics of the sintered magnet, the grain boundary phase that has been in the liquid phase is made amorphous by increasing the average cooling rate to 100 ° C./min or more. For example, Patent Documents 3 to 13 describe that the cooling rate is adjusted and additional heat treatment is performed when the sintered body temperature is lowered from the sintering temperature to room temperature.
[0009]
Since the coercive force expression mechanism of the R—Fe—B rare earth sintered magnet is a nucleation type, the state of the grain boundary phase strongly affects the coercive force. Conventionally, it is known that if the average cooling rate immediately after sintering is too slow, the grain boundary phase is completely crystallized and the coercive force is lowered. That is, in order to mass-produce sintered magnets with high coercive force, it is necessary to completely amorphize the grain boundary phase, and the average cooling rate when cooling the sintered body to the room temperature level is 100 ° C. / More than a minute is done.
[0010]
In particular, when Co is added to the raw material alloy, if the average cooling rate after sintering is low, Nd (Fe _1-x Co _x ) ₂ Since a ferromagnetic compound such as x (0.5 to 0.7) is precipitated in the grain boundary phase, the coercive force may be lowered. The ferromagnetic compound thus produced may be demagnetized by applying a high-temperature heat treatment exceeding 900 ° C. However, if such a high-temperature heat treatment is performed, the grain growth of the sintered body proceeds. As a result, the coercive force is lowered.
[0011]
On the other hand, conventionally, heat treatment (aging treatment) at about 400 to 500 ° C. has been performed on the sintered body for the purpose of improving the magnet characteristics. Fine irregularities exist in the grain boundary of the sintered body immediately after sintering, which hinders the improvement of the coercive force, but if such a sintered body is subjected to heat treatment at about 400 to 500 ° C., Since the interface between the field phase and the main phase changes to a smooth one, the coercive force is improved.
[0012]
[Patent Document 1]
Japanese Patent Publication No. 6-6728
[Patent Document 2]
US Pat. No. 5,383,978
[Patent Document 3]
JP-A-60-77961
[Patent Document 4]
JP 60-182107 A
[Patent Document 5]
JP-A-60-184626
[Patent Document 6]
JP-A-61-87825
[Patent Document 7]
JP 61-81605 A
[Patent Document 8]
Japanese Patent Application Laid-Open No. 61-252604
[Patent Document 9]
JP-A 61-264133
[Patent Document 10]
JP-A-1-232704
[Patent Document 11]
Japanese Patent Publication No. 2-16368
[Patent Document 12]
Japanese Patent No. 272040
[Patent Document 13]
Japanese Patent No. 2948223
[0013]
[Problems to be solved by the invention]
However, when a powder compact made from a quenched alloy is sintered, the effect of improving the coercive force is not sufficient even if a heat treatment of about 400 to 500 ° C. is additionally performed, so that a higher coercive force is achieved. Requires detailed analysis and modification of the grain boundary phase.
[0014]
The present invention has been made in view of such various points, and its main object is to provide a rare earth sintered magnet having improved coercive force while using a quenched alloy in which a ferromagnetic compound of a transition metal is likely to precipitate, and a method for producing the same. There is to do.
[0015]
[Means for Solving the Problems]
The rare earth sintered magnet of the present invention is R ₂ Fe ₁₄ A rare earth sintered magnet having a main phase composed of a compound having a B-type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase, wherein the grain boundary phase is It includes an amorphous layer portion that covers the surface of the main phase thinly, and a nonmagnetic crystal layer portion that exists in a region sandwiched by the amorphous layer portions.
[0016]
In a preferred embodiment, the ratio of the area occupied by the nonmagnetic crystal layer portion in the grain boundary phase in an arbitrary cross section is 20% or more.
[0017]
In a preferred embodiment, the main phase has an average particle size of 1.5 μm or more and 20 μm or less.
[0018]
In a preferred embodiment, the nonmagnetic crystal layer portion contains a transition metal element.
[0019]
In a preferred embodiment, the concentration of Mn and Si is adjusted to 1% by mass or less.
[0020]
In a preferred embodiment, the average thickness of the grain boundary phase is 1 nm or more and 100 nm or less.
[0021]
In a preferred embodiment, the raw material alloy is a quenched alloy.
[0022]
The method for producing a rare earth sintered magnet according to the present invention is described as R ₂ Fe ₁₄ A method for producing a rare earth sintered magnet having a main phase composed of a compound having a B-type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase, wherein R-Fe- A step of preparing a B-based quenched alloy powder, a step of forming a sintered body by sintering the powder at a temperature T1 of 1000 ° C. or higher and 1100 ° C. or lower; The step of lowering the temperature of the sintered body to a temperature T2 of 400 ° C. or higher and 900 ° C. or lower by cooling at an average cooling rate of less than or equal to minutes, and the heat treatment at the temperature T2 for 0.01 hours for the sintered body And a step of performing for 100 hours or less.
[0023]
Another method of manufacturing a rare earth sintered magnet according to the present invention is R ₂ Fe ₁₄ A method for producing a rare earth sintered magnet having a main phase composed of a compound having a B-type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase, wherein R-Fe- A step of preparing a B-based quenched alloy powder, a step of forming a sintered body by sintering the powder at a temperature T1 of 1000 ° C. to 1100 ° C., and cooling the sintered body from the temperature T1. The step of lowering the temperature of the sintered body to a temperature lower than 400 ° C. and heating the sintered body to a temperature T 3 higher than the temperature T 2 of 400 ° C. to 900 ° C. A step of raising the temperature, a step of lowering the temperature of the sintered body to a temperature T2 by cooling the sintered body at an average cooling rate of 50 ° C./min or less from the temperature T3, and the sintered body Heat treatment at temperature T2 And a step of performing during the following 100 hours or more .01 hours.
[0024]
Still another method of manufacturing a rare earth sintered magnet according to the present invention is R ₂ Fe ₁₄ A method for producing a rare earth sintered magnet having a main phase composed of a compound having a B-type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase, ₂ Fe ₁₄ Preparing a sintered body having a main phase composed of a compound having a B-type crystal structure and an amorphous grain boundary phase located at a grain boundary portion of the main phase; and After heating to a temperature T3 higher than a temperature T2 of 900 ° C. or lower, and cooling at an average cooling rate of 50 ° C./min or lower, the temperature of the sintered body is reduced to a temperature T2 of 400 ° C. or higher and 900 ° C. or lower. And performing the heat treatment at that temperature for 0.01 hour or more and 100 hours or less.
[0025]
The step of preparing the R-Fe-B quenching alloy powder is performed by using a molten raw material alloy 10 ² ℃ / second or more 2 × 10 ^Four It includes a step of producing a quenched alloy by cooling at an average cooling rate of not more than ° C./second, and a step of pulverizing the quenched alloy.
[0026]
In a preferred embodiment, the molten material alloy is cooled by a strip casting method.
[0027]
In a preferred embodiment, the step of crushing the quenched alloy includes a step of embrittlement of the quenched alloy by hydrogen treatment.
[0028]
DETAILED DESCRIPTION OF THE INVENTION
The present inventor forms a grain boundary phase having a novel structure that has not been conventionally obtained by controlling the process of cooling the sintered body after sintering the compact of the quenched alloy powder (controlled cooling), As a result, it has been found that the coercive force can be improved, and the present invention has been conceived.
[0029]
First, basic features of the present invention will be described with reference to FIGS.
[0030]
FIG. 1 (a) shows a typical temperature profile in a conventional sintering process. In the illustrated example, the sintering temperature T1 is 1080 ° C., and the holding time at the temperature T1 is 6 hours. By the heat treatment for 6 hours at temperature T1, the powder compact is sintered and a sintered body is formed. The sintered body temperature at this time is T1, and the grain boundary part is in a liquid phase. Thereafter, the sintered body temperature is rapidly lowered from T1 by rapid cooling. The average cooling rate at this time is usually set to 100 ° C./min or more, typically about 150 ° C./min, and the grain boundary phase of the sintered body becomes amorphous. FIG. 2A is a cross-sectional view schematically showing a structure of a conventional rare earth sintered magnet having an amorphized grain boundary phase.
[0031]
FIG. 1 (b) shows an example of a temperature profile for sintering and controlled cooling employed in the present invention. In the present invention, after sintering for a predetermined time at a temperature T1 of 1000 ° C. or more and 1100 ° C. or less, it is gradually cooled from the temperature T1 at an average cooling rate of 50 ° C./min or less (preferably 30 ° C./min or less), The temperature of the sintered body is lowered to a temperature T2 of 400 ° C. or higher and 900 ° C. or lower. And after performing the heat processing at the temperature T2 for a predetermined time, it cools to room temperature level. The average cooling rate after the heat treatment at the temperature T2 may be rapidly performed at a rate exceeding 100 ° C./min.
[0032]
In the example shown in FIG. 1B, the temperature T1 is 1080 ° C., the temperature T2 is 630 ° C., and the cooling from the temperature T1 to the temperature T2 takes 6 hours. The average cooling rate at this time is 1.25 ° C./min. The heat treatment at temperature T2 is performed for 12 hours.
[0033]
After the rapidly cooled alloy powder is sintered, the grain boundary portion of the sintered body has a structure as shown in FIG. That is, R ₂ Fe ₁₄ An amorphous layer portion (thickness: for example, 0.1 nm to 10 nm) that thinly covers the surface of the main phase made of a compound having a B-type crystal structure (R is a rare earth element), and a region sandwiched between the amorphous layer portions A grain boundary phase having a structure (layered structure) including a crystal layer portion (thickness: for example, 1 nm to 10 nm) existing in the film is formed. The amorphous layer portion is a region containing a relatively large amount of rare earth elements, and the crystalline layer portion is a nonmagnetic region containing a transition metal. The thickness of the amorphous layer portion and the crystal layer portion is a thickness in a region excluding the vicinity of the triple point of the crystal grain boundary, and the average thickness of the grain boundary phase is, for example, 1.2 nm or more and 30 nm or less. Become. The thickness in the vicinity of the triple point may reach 1000 nm. When the triple point is considered, the average thickness of the entire grain boundary phase is 1 nm or more and 100 nm or less.
[0034]
When quenching is performed after sintering, as shown in FIG. 2A, the grain boundary phase becomes amorphous and, for example, NdCo ₂ A transition metal ferromagnetic compound or a ferromagnetic element is deposited. In particular, when the powder is made of a quenched alloy as in the present invention, the transition metal ferromagnetic compound is likely to precipitate in the grain boundary phase. In the case of a quenched alloy manufactured by the strip casting method, etc., the time required for the molten alloy alloy to solidify is short, so solidification is completed before the transition metal distribution contained in the molten alloy reaches a complete equilibrium state. End up. As a result, a large amount of transition metal is present in the grain boundary phase, and a ferromagnetic compound is easily formed.
[0035]
On the other hand, according to the present invention, the transition metal present in the grain boundary phase is controlled by appropriately controlling the temperature profile of the cooling process, despite the fact that the sintered body is produced using the quenched alloy powder. It can concentrate as a nonmagnetic compound located in the center part of a grain boundary phase, and as a result, the coercive force of a magnet can be improved. In particular, in the present invention, since all of the grain boundary phase is not crystallized and the surface of the main phase is covered with a thin amorphous layer, the occurrence of reverse magnetic domains is suppressed, and the conventional grain boundary phase is provided. It is thought that coercive force higher than that of a sintered magnet is developed.
[0036]
In addition, instead of monotonously cooling the sintered body from the temperature T1 to the temperature T2, a resuperheat treatment as shown in FIG. 1C is performed between the sintering at the temperature T1 and the heat treatment at the temperature T2. Also good.
[0037]
FIG.1 (c) has shown the example of the heat processing profile of the other sintering process employ | adopted by this invention. In this example, the sintered body temperature is lowered to a low temperature (typically a room temperature level) once lower than 400 ° C. after sintering. The average cooling rate at this time may exceed 100 ° C./min as in the prior art. Next, the sintered body is reheated to a temperature T3 exceeding the temperature T2 (set in a range of 800 ° C. or higher and 1000 ° C. or lower), held for a predetermined time, and then sintered at an average cooling rate of 50 ° C./min or lower. The bonded body is cooled (controlled cooling), the sintered body temperature is lowered to T2, and the temperature T2 is maintained for a predetermined time. In the example of FIG. 1C, the temperature T3 is 750 ° C., the holding time at the temperature T3 is 8 hours, the temperature T2 is 630 ° C., and the holding time at the temperature T2 is 12 hours.
[0038]
Of the temperature profile shown in FIG. 1 (c), the portion related to sintering before performing re-superheat treatment is the same as the temperature profile shown in FIG. 1 (a), so a low temperature below 400 ° C. after sintering. When the temperature of the sintered body is lowered to (typically the room temperature level), the entire grain boundary phase of the sintered body is amorphized, and is in the same state as a conventional sintered body. In the example of FIG. 1 (c), the structure of the grain boundary phase is changed to the above-described structure (amorphous layer portion) by continuously performing reheating treatment and controlled cooling treatment on such a sintered body. In a region surrounded by a non-magnetic crystal portion). In other words, it is possible to obtain the sintered magnet of the present invention by changing the structure of the grain boundary phase of the sintered magnet by performing such reheating treatment on the sintered magnet on the market. Is possible.
[0039]
During the reheating treatment at 800 ° C. to 1000 ° C., the grain boundary portion of the sintered body becomes a liquid phase and atoms diffuse. In the subsequent slow cooling process, a non-magnetic crystal part grows in the center of the grain boundary phase. At this time, since the transition metal present in the grain boundary phase is taken into the nonmagnetic crystal at the center of the grain boundary phase, a ferromagnetic compound that causes a reduction in coercive force is not formed. Further, it is considered that the nonmagnetic property of the ferromagnetic compound further proceeds and the nonmagnetic compound is stabilized by the subsequent heat treatment at the temperature T2.
[0040]
It is considered that a phenomenon similar to the phenomenon occurring in the process of such reheating treatment and control cooling occurs also in the heat treatment (control cooling process) shown in FIG. The difference between the two lies in whether or not the grain boundary phase once amorphized is converted into a liquid phase.
[0041]
Note that the interface between the grain boundary phase and the main phase shown in FIG. 2B is not smooth, but further, by performing a known aging treatment in the range of 430 ° C. to 600 ° C., the main phase and the grain boundary The shape of the interface with the phase can be made smooth. FIG. 2 (c) schematically shows the grain boundary phase and its vicinity after the aging treatment. When the interface is matched by such an aging treatment, the coercive force is further improved.
[0042]
Hereinafter, preferred embodiments of the present invention will be described with reference to the drawings.
[0043]
[Raw material alloy]
First, using a single roll strip caster (quenching device) shown in FIG. 3, a raw material alloy of an R—Fe—B magnet alloy having a desired composition is prepared. The quenching device of FIG. 3 has a quenching chamber 1 in which the inside can be brought into a vacuum state or a reduced pressure state in an inert gas atmosphere. A melting furnace 3 for forming the molten metal 2, a cooling roll 5 for rapidly cooling and solidifying the molten alloy 2 supplied from the melting furnace 3, and a chute for guiding the molten alloy 3 from the melting furnace 3 to the cooling roll 5 (tan 4) and a recovery container 8 for recovering the ribbon-shaped alloy 7 which has solidified and has been peeled off from the cooling roll 5.
[0044]
The melting furnace 3 can supply the molten metal 2 produced by melting the alloy raw material to the chute 4 at a substantially constant supply amount. This supply amount can be arbitrarily adjusted by controlling the operation of tilting the melting furnace 3.
[0045]
The outer peripheral surface of the cooling roll 5 is made of a material having good thermal conductivity such as copper, and has a diameter of 30 cm to 100 cm and a width of 15 cm to 100 cm, for example. The cooling roll 5 is cooled by passing water through the roll. The cooling roll 5 can be rotated at a predetermined rotational speed by a driving device (not shown). By controlling this rotational speed, the peripheral speed of the cooling roll 5 can be arbitrarily adjusted. The average cooling rate by the quenching device is, for example, 10 by selecting the rotation speed of the cooling roll 5 or the like. ² ℃ / sec ～ 2 × 10 ^Four It can be controlled in the range of ° C / second.
[0046]
The tip of the chute 4 is disposed at a position having a certain angle θ with respect to a line connecting the topmost part of the cooling roll 5 and the center of the roll. The molten alloy 2 supplied onto the chute 4 is supplied from the tip of the chute 4 to the cooling roll 5.
[0047]
The chute 4 is made of ceramics or the like, delays the flow rate so as to temporarily store the molten alloy 2 continuously supplied from the melting furnace 3 at a predetermined flow rate, and rectifies the flow of the molten alloy 2. Can do.
[0048]
By using the chute 4, the molten alloy 2 can be supplied in a state where the cooling roll 5 is expanded to a substantially uniform thickness over a certain width in the trunk length direction (axial direction). In addition to the above function, the chute 4 has a function of adjusting the temperature of the molten alloy 2 immediately before reaching the cooling roll 5. The temperature of the molten alloy 2 on the chute 4 is preferably 100 ° C. or higher than the liquidus temperature. This is because if the temperature of the molten alloy 2 is too low, primary crystals that adversely affect the alloy characteristics after rapid cooling are locally nucleated and may remain after solidification. The molten metal residence temperature on the chute 4 can be controlled by adjusting the molten metal temperature at the time of pouring from the melting furnace 3 into the chute 4 and the heat capacity of the chute 3 itself. ) May be provided.
[0049]
Using the above quenching apparatus, for example, Nd: 33% by mass, Dy: 1.0% by mass, Co: 1.0% by mass, B: 1.0% by mass, the balance Fe and inevitable impurities (Mn and Si are 0 .1% by mass or less) is melted to form a molten alloy. After this molten alloy was kept at 1350 ° C., it was brought into contact with the surface of a cooling roll, and the molten alloy was quenched to obtain a flaky alloy ingot having a thickness of about 0.1 to 5 mm. The rapid cooling conditions at this time are a roll peripheral speed of about 1 to 3 m / sec and an average cooling speed of 10 ² ~ 2x10 ^Four C./second.
[0050]
The quenched alloy slab thus produced is pulverized into flakes having a size of 1 to 10 mm before the next hydrogen pulverization. In addition, the manufacturing method of the raw material alloy by the strip casting method is also disclosed in, for example, US Pat. No. 5,383,978.
[0051]
[First grinding step]
The raw material alloy slabs roughly crushed into flakes are filled into a plurality of raw material packs (for example, made of stainless steel) and mounted on a rack. Thereafter, the rack on which the raw material pack is mounted is inserted into the hydrogen furnace. Next, the lid of the hydrogen furnace is closed, and a hydrogen embrittlement process (hereinafter sometimes referred to as “hydrogen pulverization process”) is started. The hydrogen crushing process is executed according to a temperature profile shown in FIG. 4, for example. In the example of FIG. 4, first, the vacuuming process I is performed for 0.5 hours, and then the hydrogen storage process II is performed for 2.5 hours. In the hydrogen storage process II, hydrogen gas is supplied into the furnace and the furnace is filled with a hydrogen atmosphere. The hydrogen pressure at that time is preferably about 200 to 400 kPa.
[0052]
Subsequently, after the dehydrogenation process III is performed for 5.0 hours under a reduced pressure of about 0 to 3 Pa, the raw material alloy cooling process IV is performed for 5.0 hours while supplying argon gas into the furnace.
[0053]
In the cooling process IV, when the atmospheric temperature in the furnace is relatively high (for example, when the temperature exceeds 100 ° C.), normal temperature inert gas is supplied into the hydrogen furnace and cooled. Thereafter, when the raw material alloy temperature is lowered to a relatively low level (for example, when the temperature is 100 ° C. or lower), the inert gas cooled to a temperature lower than normal temperature (for example, about room temperature minus 10 ° C.) is introduced into the hydrogen furnace 10. It is preferable to supply from the viewpoint of cooling efficiency. The supply amount of argon gas is 10 to 100 Nm ^Three / Min.
[0054]
When the temperature of the raw material alloy is lowered to about 20 to 25 ° C., an inert gas at a substantially normal temperature (a temperature lower than room temperature, but a temperature within a range of 5 ° C. or less) is blown into the hydrogen furnace, It is preferable to wait until the temperature reaches the normal temperature level. By doing so, it is possible to avoid a situation in which condensation occurs inside the furnace when the lid of the hydrogen furnace is opened. If moisture is present inside the furnace due to condensation, the moisture is frozen and vaporized in the evacuation process, which makes it difficult to increase the degree of vacuum and the time required for the evacuation process I is not preferable.
[0055]
When the coarsely pulverized alloy powder after hydrogen pulverization is taken out from the hydrogen furnace, it is preferable to perform the take-out operation in an inert atmosphere so that the coarsely pulverized powder does not come into contact with the atmosphere. This is because the coarsely pulverized powder is prevented from oxidizing and generating heat, and the magnetic properties of the magnet are improved. Next, the coarsely pulverized raw material alloy is filled into a plurality of raw material packs and mounted on a rack.
[0056]
By the hydrogen treatment, the rare earth alloy is pulverized to a size of about 0.1 mm to several mm, for example, and the average particle size becomes 500 μm or less. After the hydrogen pulverization, it is preferable that the embrittled raw material alloy is further crushed and cooled by a cooling device such as a rotary cooler. When the raw material is taken out in a relatively high temperature state, the time for the cooling process by a rotary cooler or the like may be made relatively long.
[0057]
A large amount of rare earth elements such as Nd are exposed on the surface of the coarsely pulverized powder produced by hydrogen pulverization, and it is very easily oxidized. Before the next pulverization step, zinc stearate is added as a pulverization aid in an amount of about 0.04% by mass.
[0058]
[Second grinding step]
Next, the coarsely pulverized powder produced in the first pulverization step is finely pulverized using a jet mill device. In the present embodiment, a cyclone classifier suitable for removing ultrafine powder is connected to the pulverizer.
[0059]
Hereinafter, the fine pulverization step (second pulverization step) performed using the jet mill apparatus will be described in detail with reference to FIG.
[0060]
The illustrated jet mill apparatus 10 includes a raw material charging machine 12 for supplying a rare earth alloy (a material to be pulverized) coarsely pulverized in a first pulverization process, and a pulverizer for pulverizing the material to be crushed from the raw material charging machine 12. 14, a cyclone classifier 16 that classifies the powder obtained by pulverizing the material to be crushed by the pulverizer 14, and a collection tank 18 that collects the powder having a predetermined particle size distribution classified by the cyclone classifier 16. I have.
[0061]
The raw material charging machine 12 includes a raw material tank 20 that accommodates the material to be crushed, a motor 22 that controls the supply amount of the material to be crushed from the raw material tank 20, and a spiral feeder connected to the motor 22 (screw feeder). 24.
[0062]
The pulverizer 14 has a vertically long and substantially cylindrical pulverizer body 26, and a plurality of nozzles for attaching nozzles for injecting inert gas (for example, nitrogen) at high speed to the lower portion of the pulverizer body 26. A mouth 28 is provided. A raw material input pipe 30 for supplying a material to be crushed into the pulverizer body 26 is connected to a side portion of the pulverizer body 26.
[0063]
The raw material input pipe 30 is provided with a valve 32 for temporarily holding the material to be supplied and confining the pressure inside the pulverizer 14, and the valve 32 has a pair of an upper valve 32a and a lower valve 32b. is doing. The feeder 24 and the raw material input pipe 30 are connected by a flexible pipe 34.
[0064]
The pulverizer 14 includes a classification rotor 36 provided above the pulverizer body 26, a motor 38 provided above the pulverizer body 26, and a connection pipe 40 provided above the pulverizer body 26. Have. The motor 38 drives the classification rotor 36, and the connection pipe 40 discharges the powder classified by the classification rotor 36 to the outside of the pulverizer 14.
[0065]
The pulverizer 14 includes a plurality of leg portions 42 serving as support portions. A base 44 is disposed in the vicinity of the outer periphery of the pulverizer 14, and the pulverizer 14 is placed on the base 44 by the legs 42. In the present embodiment, a weight detector 46 such as a load cell is provided between the leg portion 42 and the base 44 of the pulverizer 14. Based on the output from the weight detector 46, the control unit 48 can control the rotation speed of the motor 22 and thereby control the input amount of the object to be crushed.
[0066]
The cyclone classifier 16 has a classifier body 64, and an exhaust pipe 66 is inserted into the classifier body 64 from above. An introduction port 68 for introducing the powder classified by the classifying rotor 36 is provided on the side of the classifier body 64, and the introduction port 68 is connected to the connection pipe 40 by a flexible pipe 70. An outlet 72 is provided in the lower part of the classifier main body 64, and a desired finely pulverized powder recovery tank 18 is connected to the outlet 72.
[0067]
The flexible pipes 34 and 70 are preferably made of resin, rubber, or the like, or made of a highly rigid material in a bellows shape or a coil shape so as to have flexibility. When such flexible pipes 34 and 70 are used, changes in the weight of the raw material tank 20, the feeder 24, the classifier main body 64, and the collection tank 18 are not transmitted to the legs 42 of the pulverizer 14. Therefore, if the weight is detected by the weight detector 46 provided in the leg portion 42, the weight of the pulverized material staying in the pulverizer 14 and its change amount can be accurately detected, and the pulverized material supplied into the pulverizer 14 is supplied. The amount of objects can be controlled accurately.
[0068]
Next, a grinding method using the jet mill apparatus 10 will be described.
[0069]
First, the material to be crushed is charged into the raw material tank 20. The material to be crushed in the raw material tank 20 is supplied to the pulverizer 14 by the supply machine 24. At this time, the supply amount of the object to be crushed can be adjusted by controlling the rotation speed of the motor 22. The material to be crushed supplied from the feeder 24 is temporarily dammed up by the valve 32. Here, the pair of upper valve 32a and lower valve 32b alternately perform opening and closing operations. That is, when the upper valve 32a is opened, the lower valve 32b is closed, and when the upper valve 32a is closed, the lower valve 32b is opened. Thus, by alternately opening and closing the pair of valves 32a and 32b, the pressure in the pulverizer 14 can be prevented from leaking to the raw material input device 12 side. As a result, the material to be crushed is supplied between the pair of the upper valve 32a and the lower valve 32b when the upper valve 32a is opened. Then, when the lower valve 32 b is next opened, it is guided to the raw material input pipe 30 and introduced into the pulverizer 14. The valve 32 is driven at a high speed by a sequence circuit (not shown) different from the control circuit 48, and the material to be crushed is continuously supplied into the pulverizer 14.
[0070]
The object to be crushed introduced into the pulverizer 14 is wound up into the pulverizer 14 by high-speed injection of an inert gas from the nozzle port 28 and swirls with the high-speed air current in the apparatus. And it pulverizes finely by the mutual collision of the objects to be crushed.
[0071]
The finely pulverized powder particles are guided to the classifying rotor 36 in an ascending air current, classified by the classifying rotor 36, and the coarse powder is pulverized again. On the other hand, the powder pulverized to a predetermined particle size or less is introduced into the classifier body 64 of the cyclone classifier 16 from the introduction port 68 via the connection pipe 40 and the flexible pipe 70. By using the classification rotor 36, powder particles having a particle size larger than the particle size showing the peak of the particle size distribution can be efficiently removed. If there are many coarse powder particles having a particle size of more than 10 μm in the finally obtained powder, the coercive force of the sintered magnet will be reduced. Therefore, powder particles having a particle size of more than 10 μm using the classification rotor 36. Is preferably reduced. In the present embodiment, the final powder is adjusted to 10% or less of the total particle volume of particles having a particle size exceeding 10 μm.
[0072]
In the classifier body 16, relatively large powder particles having a predetermined particle size or more are deposited in a recovery tank 18 installed at the lower part, but the ultrafine powder is discharged to the outside through the exhaust pipe 66 together with the inert gas flow. . In the present embodiment, the ultrafine powder is removed through the exhaust pipe 66, thereby reducing the volume ratio of the ultrafine powder (particle size: 1.0 μm or less) to the powder collected in the recovery tank 18. In a preferred embodiment, the volume ratio of the ultrafine powder (particle size: 1.0 μm or less) is adjusted to 10% or less.
[0073]
When the R-rich ultrafine powder is removed in this way, the amount of the rare earth element R in the sintered magnet consumed for bonding with oxygen can be reduced, and the magnet characteristics can be improved.
[0074]
As described above, in the present embodiment, the cyclone classifier 16 with blow-up is used as a classifier connected to the subsequent stage of the jet mill apparatus (pulverizer 14). According to such a cyclone classifier 16, ultrafine powder having a predetermined particle size or less rises upside down without being collected in the collection tank 18, and is discharged from the pipe 66 to the outside of the apparatus.
[0075]
The particle size of the fine powder removed from the pipe 66 to the outside of the apparatus can be controlled by appropriately defining the parameters of each part of the cyclone and adjusting the pressure of the inert gas flow.
[0076]
According to this embodiment, the alloy powder has an average particle size (FSSS particle size) of, for example, 4.0 μm or less, and the volume of ultrafine powder having a particle size of 1.0 μm or less is 10% or less of the total powder volume. Can be obtained.
[0077]
In order to suppress oxidation in the pulverization process as much as possible, it is preferable to suppress the amount of oxygen in the high-speed gas stream (inert gas) used when pulverizing to within a range of, for example, about 1000 to 20000 volume ppm, It is more preferable to keep it low to about 5000 to 10000 volume ppm. A fine pulverization method for controlling the oxygen concentration in the high-speed gas stream is described in Japanese Patent Publication No. 6-6728.
[0078]
As described above, it is preferable to adjust the oxygen content of the alloy powder after pulverization to 6000 mass ppm or less as a whole by controlling the concentration of oxygen contained in the atmosphere during pulverization. If the oxygen content in the rare earth alloy powder exceeds 6000 ppm by mass, the proportion of non-magnetic oxide in the sintered magnet increases and the magnetic properties of the final sintered magnet deteriorate. Because it ends up.
[0079]
In the present embodiment, since the R-rich ultrafine powder is appropriately removed, the oxygen concentration of the powder is controlled to 6000 mass ppm or less by adjusting the oxygen concentration in the inert gas atmosphere at the time of fine pulverization. It is possible to reduce the oxygen concentration in the inert gas atmosphere if the volume ratio of the ultrafine powder exceeds 10% of the total without removing the R-rich ultrafine powder. However, the oxygen concentration in the finally obtained powder exceeds 6000 mass ppm. However, when the powder is molded in an air atmosphere, it is preferable to contain 3500 mass ppm or more of oxygen in the powder in order to suppress oxidation and heat generation of the molded body.
[0080]
In the present embodiment, the second pulverization step is performed using the jet mill apparatus 10 having the configuration shown in FIG. 5, but the present invention is not limited to this, and the jet mill apparatus having another configuration, or Other types of pulverizers (for example, an attritor or a ball mill pulverizer) may be used. Moreover, you may use classifiers other than a cyclone classifier as a classifier for removing ultrafine powder.
[0081]
[Addition of lubricant]
A liquid lubricant or binder mainly composed of a fatty acid ester or the like is added to the raw material alloy powder produced by the above method. For example, it is preferable to add and mix, for example, 0.15 to 5.0% by mass of a lubricant in an inert atmosphere using an apparatus such as a rocking mixer. Examples of fatty acid esters include methyl caproate, methyl caprylate, and methyl laurate. The important point is that the lubricant can be volatilized and removed in a later step. In addition, when the lubricant itself is a solid that is difficult to be uniformly mixed with the alloy powder, the lubricant may be diluted with a solvent. As the solvent, a petroleum solvent typified by isoparaffin, a naphthene solvent, or the like can be used. The timing of addition of the lubricant is arbitrary, and may be any of before pulverization, during pulverization, and after pulverization. The liquid lubricant coats the surface of the powder particles and exhibits an antioxidant effect on the particles. The liquid lubricant also has a function of making the density of the molded body uniform during pressing, reducing friction between powder particles, improving compressibility, and suppressing the disorder of orientation. When a solid lubricant such as zinc stearate is used, it can be added before pulverization and mixed during pulverization. This mixing may be performed with a rocking mixer after pulverization.
[0082]
[Press molding]
Next, the magnetic powder produced by the above-described method is molded in an orientation magnetic field using a known press machine. In the present embodiment, the press pressure is adjusted in the range of 5 to 100 MPa, preferably 15 to 40 MPa, in order to increase the orientation in the magnetic field. After the press molding is completed, the powder compact is pushed up by the lower punch and taken out of the press apparatus.
[0083]
Next, the molded body is placed on a sintering base plate made of, for example, a molybdenum material and mounted on the sintering case together with the base plate. The sintered case carrying the sintered body is transferred into a sintering furnace and a sintering process is performed in the furnace.
[0084]
The sintering process in this embodiment may be performed using either a batch furnace or a continuous furnace. In a preferred embodiment of the present invention, the heater and cooling mechanism of the sintering furnace are controlled so that the temperature of the molded body (sintered body) changes with a profile as shown in FIGS.
[0085]
The temperature profiles shown in FIGS. 1B and 1C are merely examples that can be adopted in the present invention. The time for maintaining the temperatures T1 and T3 is appropriately determined within the range of 0.5 to 100 hours and 0.01 to 100 hours, respectively. As described above, the holding time at the temperature T2 is set within the range of 0.01 hours to 100 hours, and preferably 0.5 hours to 50 hours.
[0086]
Further, the temperature profile until the temperature is raised to the sintering temperature T1 is not limited to that shown in FIGS. 1B and 1C, and the temperature profile is maintained for a predetermined time at the temperature before reaching the sintering temperature T1. A plurality of steps of heat treatment may be performed. The important point for the present invention is to form the grain boundary phase having the structure shown in FIG. 2B by controlling the cooling process of the liquid phase grain boundary phase. In order to suppress the grain growth of the main phase by shortening the holding time at the sintering temperature T1, dehydration is performed at a temperature around 800 ° C. for a predetermined time (for example, 0.5 to 6 hours) before reaching the sintering temperature T1. It is preferable to perform an elementary process.
[0087]
As the atmosphere in the furnace when performing the heat treatment for sintering, a gas (inert gas such as argon gas) that does not easily react with the powder of the compact is used. Further, the pressure of the atmospheric gas is preferably set lower than the atmospheric pressure, and is reduced to, for example, 17 Pa or less.
[0088]
After the sintering process of the present embodiment, the sintered body finally cooled to the room temperature level is appropriately subjected to processing such as cutting, polishing, and protective film deposition.
[0089]
[Alloy composition]
Specifically, as the rare earth element R, it is preferable to use at least one element of Pr, Nd, Sm, Gd, Tb, Dy, Ho, Er, Tm, and Lu. In order to obtain sufficient magnetization, 50 atomic% or more of the rare earth element R is preferably occupied by either or both of Pr and Nd.
[0090]
If the rare earth element R is less than 8 atomic%, the coercive force may decrease due to precipitation of the α-Fe phase. When the rare earth element R exceeds 18 atomic%, the desired tetragonal R ₂ Fe ₁₄ In addition to the B-type compound, a large amount of the R-rich second phase may be precipitated and the magnetization may be reduced. For this reason, it is preferable that the rare earth element R exists in the range of 8-18 atomic% of the whole.
[0091]
In order to improve the magnet characteristics, a part of Fe may be replaced with Co. Co addition is likely to cause the precipitation of ferromagnetic compounds in the grain boundary phase, particularly when a quenched alloy is used, leading to a decrease in coercive force. However, according to the present invention, when 0.5 atomic% or more of Co is added However, almost no ferromagnetic Co compound is formed, and the coercive force is improved. In addition to Co, transition metal elements such as Ni, V, Cr, Mn, Cu, Zr, Mb, and Mo can be used as the transition metal element that replaces Fe. The proportion of Fe in the entire transition metal element is preferably 50 atomic% or more. When the proportion of Fe falls below 50 atomic%, Nd ₂ Fe ₁₄ This is because the saturation magnetization itself of the B-type compound is reduced.
[0092]
B and / or C is tetragonal Nd ₂ Fe ₁₄ It is essential for stably depositing the B-type crystal structure (particularly B is essential). When the addition amount of B and / or C is less than 3 atomic%, R ₂ T ₁₇ Since the phase precipitates, the coercive force decreases, and the squareness of the demagnetization curve is significantly impaired. On the other hand, when the addition amount of B and / or C exceeds 20 atomic%, a second phase with small magnetization is precipitated.
[0093]
In order to further increase the magnetic anisotropy of the powder, another additive element M may be added. As the additive element M, at least one element selected from the group consisting of Al, Ti, V, Cr, Ni, Ga, Zr, Nb, Mo, In, Sn, Hf, Ta, and W is preferably used. The Such an additive element M may not be added at all. When adding, it is preferable to make addition amount into 3 atomic% or less. This is because if the added amount exceeds 3 atomic%, the second phase precipitates, not ferromagnetism, and the magnetization decreases. In order to obtain magnetically isotropic magnetic powder, the additive element M is not necessary, but Al, Cu, Ga, or the like may be added to increase the intrinsic coercive force.
[0094]
[Examples and Comparative Examples]
In this example, an alloy melt containing 33 mass% Nd, 1.0 mass% Dy, 1.0 mass% Co, 1.0 mass% B, and Fe (residue) is strip cast. And solidified by cooling.
[0095]
The quenched alloy was pulverized by the pulverization method described above to produce finely pulverized powder having an average particle size (FSSS particle size) of about 2.8 to 4.0 μm. The particle size distribution was measured by using a particle size distribution measuring apparatus (product number HELOS Particle Size Analyzer) manufactured by Sympatec. This particle size distribution measuring device uses the fact that the amount of transmitted light is reduced when a laser beam scanned at high speed is blocked by particles, and directly calculates the particle size from the time required for the laser beam to pass through the particles. Can be requested.
[0096]
Next, 0.3% by mass of methyl caproate diluted with a petroleum solvent was added to these powders, mixed, and then press-molded with a mold press machine to obtain a width of 20 mm × A powder compact having a size of 20 mm in length and 10 mm in height was produced. The breath pressure is about 35 MPa and the molding density is 4.5 g / cm. ^Three It was. During pressing, an orientation magnetic field (1.0 MA / m) in a direction perpendicular to the uniaxial compression direction was applied.
[0097]
After pressing, the compact was sintered in an argon atmosphere (pressure: 2 Pa). The sintering conditions are as follows.
[0098]
Temperature increase rate from room temperature to T1 (= 1080 ° C .: 3 ° C./min)
Holding time at temperature T1: 6 hours
(Cooling rate from temperature T1 to temperature T2 (= 630 ° C.): 1.25 ° C./min)
Holding time at temperature T2: 12 hours
(Cooling rate from temperature T2 to room temperature: 80 ° C./min)
[0099]
The coercive force iHc of the sintered magnet obtained in this example was 1232 kA / m, and the residual magnetic flux density Br was 1.35T.
[0100]
Next, as a comparative example, a sintering process was performed under the following heat treatment conditions using a powder compact produced under the same conditions.
[0101]
Temperature increase rate from room temperature to T1 (= 1080 ° C .: 3 ° C./min)
Holding time at temperature T1: 6 hours
Cooling rate from temperature T1 to room temperature (= 20 ° C): 80 ° C / min)
Aging treatment: 12 hours at a temperature of 630 ° C
[0102]
As for the coercive force of the sintered magnet of the comparative example thus obtained, iHc was 1159 kA / m and the residual magnetic flux density Br was 1.35T.
[0103]
A comparison of the example and the comparative example shows that the coercive force is the same, but the coercive force is improved by about 6% in the example.
[0104]
The sintered magnets of Examples and Comparative Examples were observed with a transmission electron microscope (TEM). 6 and 7 are TEM photographs of the example and the comparative example, respectively. Sites 1 and 3 shown in the TEM photograph show the main phase region, and site 2 shows the grain boundary. As can be seen from the TEM photograph of FIG. 6, the grain boundary phase of the example has an amorphous layer portion (thickness: 0.1 nm to 10 nm) on the side in contact with the main phase, and the amorphous layer portion. There is a crystal layer portion (thickness: 1 nm to 10 nm) in the region sandwiched between (the center portion of the grain boundary phase). On the other hand, as can be seen from the TEM photograph of FIG. 7, the grain boundary phase of the comparative example is entirely amorphous. The thicknesses of the amorphous layer portion and the crystal layer portion are values when the triple point of the crystal grain boundary is not included, and may reach 1000 nm in the vicinity of the triple point.
[0105]
When the TEM photograph of this example was analyzed, in the arbitrary cross section of the rare earth sintered magnet according to the present invention, the proportion of the area occupied by the nonmagnetic crystal layer portion in the grain boundary phase was 20% or more, and the average of the main phase The particle size was 1.5 μm or more and 20 μm or less. The average thickness of the grain boundary phase was 1 nm or more and 100 nm or less.
[0106]
In addition, the average crystal grain size of the main phase in this specification indicates a value measured by a method defined in JIS H0501.
[0107]
【The invention's effect】
According to the present invention, by optimizing the heat treatment after sintering, the transition metal present in the grain boundary phase is crystallized at the center of the grain boundary phase as a nonmagnetic compound, and an amorphous layer is formed around it. Is done. The grain boundary phase having such a structure provides a rare earth magnet with improved coercive force in order to suppress the occurrence of reverse magnetic domains.
[Brief description of the drawings]
FIG. 1A is a graph showing a temperature profile of a sintering process in the prior art, and FIG. 1B is a graph showing a temperature profile of a sintering process in the first aspect of the present invention. c) is a graph which shows the temperature profile of the sintering process in the 2nd aspect of this invention. In each graph, the vertical axis represents temperature and the horizontal axis represents time.
2A is a cross-sectional view schematically showing a grain boundary phase in a conventional rare earth sintered magnet and its vicinity, and FIG. 2B is a grain boundary phase in the rare earth sintered magnet of the present invention and its vicinity. It is sectional drawing which shows the vicinity typically, (c) is sectional drawing which shows typically the grain boundary phase after the aging treatment of the sintered magnet of this invention, and its vicinity.
FIG. 3 is a view showing a configuration of a single roll type strip caster suitably used in the embodiment of the present invention.
FIG. 4 is a graph showing an example of a temperature profile of a hydrogen pulverization process performed in a coarse pulverization step according to the present invention.
FIG. 5 is a cross-sectional view showing a configuration of a jet mill apparatus suitably used for performing a pulverization step in the present invention.
FIG. 6 is a TEM photograph showing a cross-sectional structure of a sintered magnet in an example of the present invention.
FIG. 7 is a TEM photograph showing a cross-sectional structure of a sintered magnet in a comparative example.
[Explanation of symbols]
1 quenching room
2 Alloy melt
3 Melting furnace
4 Shoot (tundish)
5 Cooling roll
7 Rapidly solidified alloy
8 Collection container
10 Jet mill equipment
12 Raw material charging machine
14 Crusher
16 Cyclone classifier
18 Collection tank
20 Raw material tank
22 Motor
24 Feeder (screw feeder)
26 Crusher body
28 Nozzle port
30 Raw material input pipe
32 valves
32a Upper valve
32b Lower valve
34 Flexible pipe
36 classification rotor
38 motor
40 Connection pipe
42 legs
44 base
46 Weight detector
48 Control unit
64 classifier
66 Exhaust pipe
68 Inlet
70 Flexible pipe
72 Exit

Claims (10)

A rare earth sintered magnet having a main phase composed of a compound having an R ₂ Fe ₁₄ B type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase,
The grain boundary phase includes an amorphous layer portion that thinly covers the surface of the main phase, and a nonmagnetic crystal layer portion that exists in a region sandwiched between the amorphous layer portions,
The main phase has an average particle size of 1.5 μm or more and 20 μm or less,
The amorphous layer portion contains a relatively large amount of rare earth elements and has a thickness of 0.1 nm to 10 nm in a region excluding the vicinity of the triple point of the crystal grain boundary,
The nonmagnetic crystal layer portion includes a transition metal, and has a thickness of 1 nm to 10 nm in a region excluding the vicinity of the triple point of the grain boundary,
An R—Fe—B based rare earth sintered magnet in which, in an arbitrary cross section of the rare earth sintered magnet, the proportion of the area occupied by the nonmagnetic crystal layer portion in the grain boundary phase is 20% or more.   The R-Fe-B rare earth sintered magnet according to claim 1, wherein the content concentrations of Mn and Si are each adjusted to 1 mass% or less.   3. The R—Fe—B rare earth sintered magnet according to claim 1, wherein an average thickness of the grain boundary phase is 1 nm or more and 100 nm or less.   The R-Fe-B rare earth sintered magnet according to any one of claims 1 to 3, wherein the raw material alloy is a strip cast quenched alloy. An R—Fe—B rare earth sintered magnet having a main phase composed of a compound having an R ₂ Fe ₁₄ B type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase. A manufacturing method comprising:
Preparing a powder of R-Fe-B quenching alloy;
Forming a sintered body by sintering the powder at a temperature T1 of 1000 ° C. or higher and 1100 ° C. or lower;
Cooling the sintered body at a mean cooling rate of 30 ° C./min or less from the temperature T1 to lower the temperature of the sintered body to a temperature T2 of 400 ° C. or more and 900 ° C. or less;
Performing a heat treatment at a temperature T2 on the sintered body for a period of 0.01 hours to 100 hours;
Of manufacturing R-Fe-B rare earth sintered magnet containing An R—Fe—B rare earth sintered magnet having a main phase composed of a compound having an R ₂ Fe ₁₄ B type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase. A manufacturing method comprising:
Preparing a powder of R-Fe-B quenching alloy;
Forming a sintered body by sintering the powder at a temperature T1 of 1000 ° C. or higher and 1100 ° C. or lower;
Cooling the sintered body from temperature T1 to lower the temperature of the sintered body to a temperature lower than 400 ° C .;
Heating the sintered body to raise the temperature of the sintered body to a temperature T3 higher than a temperature T2 of 400 ° C. or higher and 900 ° C. or lower;
Lowering the temperature of the sintered body to a temperature T2 by cooling the sintered body at an average cooling rate of 30 ° C./min or less from the temperature T3;
Performing a heat treatment at a temperature T2 on the sintered body for 0.01 hours to 100 hours;
Of manufacturing R-Fe-B rare earth sintered magnet containing An R—Fe—B rare earth sintered magnet having a main phase composed of a compound having an R ₂ Fe ₁₄ B type crystal structure (R is a rare earth element) and a grain boundary phase located at a grain boundary portion of the main phase. A manufacturing method comprising:
Providing a sintered body having a main phase composed of a compound having an R ₂ Fe ₁₄ B-type crystal structure and an amorphous grain boundary phase located in a grain boundary portion of the main phase;
After heating the sintered body to a temperature T3 higher than the temperature T2 of 400 ° C. or more and 900 ° C. or less, the sintered body is cooled at an average cooling rate of 30 ° C./min or less, so that the temperature of the sintered body is 400 ° C. or more and 900 ° C. Lowering to a temperature T2 of not more than ° C., and performing a heat treatment at that temperature for 0.01 hours or more and 100 hours or less,
Of manufacturing R-Fe-B rare earth sintered magnet containing The step of preparing the R-Fe-B quenching alloy powder,
A step of producing a quenched alloy by cooling a molten raw material alloy at an average cooling rate of 10 ² ° C / second or more and 2 × 10 ⁴ ° C / second or less,
The method for producing an R—Fe—B rare earth sintered magnet according to claim 5, comprising a step of pulverizing the quenched alloy.   The method for producing an R—Fe—B rare earth sintered magnet according to claim 8, wherein the molten metal alloy is cooled by a strip casting method.   The method for producing an R—Fe—B rare earth sintered magnet according to claim 8 or 9, wherein the step of pulverizing the quenched alloy includes a step of embrittlement of the quenched alloy by hydrogen treatment.

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