JP3625821B6 - Nanocomposite magnet and method for manufacturing the same - Google Patents

Description

The present invention relates to a nanocomposite magnet suitably used for various motors and actuators and a method for producing the same, and in particular, a hard magnetic phase of a compound having an R ₂ Fe ₁₄ B type crystal structure and a soft magnetic phase such as α-Fe. Including nanocomposite magnets.

In recent years, further improvement in performance and reduction in size and weight have been required for home appliances, OA devices, electrical components, and the like. Therefore, for the permanent magnets used in such equipment, it is required to maximize the performance-to-weight ratio of the entire magnetic circuit, for example remanence B _r is 0.5 T (tesla) or more permanent It is required to use a magnet. However, by conventional relatively inexpensive hard ferrite magnets can not be the remanence B _r than 0.5 T.

Currently, the permanent magnet having a more high residual magnetic flux density B _r 0.5T, Sm-Co based magnet which is produced by the powder metallurgy method is known. Sm-Co based outside magnet, and Nd-Fe-B sintered magnets made by powder metallurgy, exert Nd-Fe-B based rapidly solidified magnet is high residual magnetic flux density B _r which is produced by a liquid quenching method can do. The former Nd—Fe—B based sintered magnet is disclosed in, for example, Patent Document 1, and the latter Nd—Fe—B based quenched magnet is disclosed in, for example, Patent Document 2.

  However, Sm-Co magnets have the disadvantage of high magnet prices because both Sm and Co as raw materials are expensive.

  Nd-Fe-B magnets contain inexpensive Fe as a main component (about 60% to 70% by weight of the total), and are therefore less expensive than Sm-Co magnets. There is a problem that the cost required is high. One of the reasons for the high manufacturing process cost is that a large-scale facility and a large number of processes are required for the separation and purification of Nd, in which the content is about 10 atomic% to 15 atomic%, and the reduction reaction. In addition, when the powder metallurgy method is used, the number of manufacturing steps is inevitably increased.

  On the other hand, an Nd—Fe—B type quenching magnet manufactured by a liquid quenching method can be obtained by a relatively simple process such as a melting process → a liquid cooling process → a heat treatment process. There is an advantage that the process cost is lower than that of the system magnet. However, in the case of the liquid quenching method, in order to obtain a bulk permanent magnet, it is necessary to mix a magnet powder produced from a quenching alloy with a resin to form a bonded magnet. The filling rate (volume ratio) is at most about 80%. Moreover, the quenched alloy produced by the liquid quenching method is magnetically isotropic.

For the above reasons, the Nd—Fe—B type quenching magnet manufactured by using the liquid quenching method has a lower _Br than the anisotropic Nd—Fe—B type sintered magnet manufactured by the powder metallurgy method. Have a problem.

As a technique for improving the characteristics of the Nd—Fe—B type quenching magnet, as described in Patent Document 3, at least one selected from the group consisting of Zr, Nb, Mo, Hf, Ta, and W is used. It is effective to add these elements in combination with at least one element selected from the group consisting of Ti, V and Cr. The addition of such elements, is improved and the coercivity H _cJ and corrosion resistance, an effective way to improve the remanence B _r is not known other than to increase the density of the bonded magnet. In addition, when the Nd-Fe-B quenching magnet contains a rare earth element of 6 atomic% or more, according to many prior arts, in order to increase the quenching rate of the melt, the melt is injected to the cooling roll through the nozzle. A melt spinning method is used.

In the case of an Nd-Fe-B quenching magnet, the composition of the rare earth element is relatively low, that is, a composition in the vicinity of Nd _3.8 Fe _77.2 B ₁₉ (atomic%), and an Fe ₃ B type compound is the main phase. Magnet materials have been proposed (Non-Patent Document 1). This permanent magnet material is obtained by subjecting an amorphous alloy produced by a liquid quenching method to a crystallization heat treatment, whereby a soft crystal Fe ₃ B phase and a hard magnetism Nd ₂ Fe ₁₄ B phase coexist. It has a metastable structure formed from a body and is called a “nanocomposite magnet”. Such nanocomposite magnet has been reported to have more high residual magnetic flux density B _r 1T, the coercive force H _cJ is relatively low and 160kA / m~240kA / m. Therefore, the use of this permanent magnet material is limited to applications where the operating point of the magnet is 1 or more.

  In addition, attempts have been made to improve the magnetic properties by adding various metal elements to the raw material alloy of the nanocomposite magnet (Patent Document 4, Patent Document 5, Patent Document 6, and Patent Document 7). "Characteristic value per cost" is not obtained. This is because a coercive force of a size that can be practically used in a nanocomposite magnet has not been obtained, and sufficient magnetic properties cannot be expressed in actual use.

Further, La having excellent amorphous forming ability is added to a raw material alloy, and a rapidly solidified alloy having an amorphous phase as a main phase is prepared by rapidly cooling a molten metal of the raw material alloy, and then Nd ₂ Fe ₁₄ B is obtained by crystallization heat treatment. There has been reported a technique in which both a phase and an α-Fe phase are precipitated and grown, and both phases are as fine as about several tens of nanometers (Non-patent Document 2). This paper shows that the addition of a small amount of refractory metal elements such as Ti (2 at%) improves the magnet characteristics and increases the composition ratio of the rare earth element Nd from 9.5 at% to 11.0 at%. Teaches that it is preferable to refine both the Nd ₂ Fe ₁₄ B phase and the α-Fe phase. The addition of the refractory metal suppresses the formation of borides (R ₂ Fe ₂₃ B ₃ and Fe ₃ B), and produces a magnet composed of only two phases of Nd ₂ Fe ₁₄ B phase and α-Fe phase. Has been done. The quenched alloy for the nanocomposite magnet is produced by a melt spinning method in which a molten alloy is injected onto the surface of a cooling roll that rotates at high speed using a nozzle. In the case of the melt spinning method, an extremely fast cooling rate can be obtained, which is suitable for producing an amorphous quenched alloy.

In order to solve the above problem, in the composition range in which the rare earth element concentration is less than 10 atomic% and the boron concentration exceeds 10 atomic%, the addition of Ti suppresses the precipitation of α-Fe during rapid cooling of the molten alloy, As a result, a nanocomposite magnet having an improved volume ratio of the compound having the R ₂ Fe ₁₄ B type crystal structure has been developed and disclosed in Patent Document 8 by the present applicant.

Patent Documents 9 and 10 disclose a large number of elements (Al, Si, V, Cr, Mn, Ga, Zr, Mb, Mo, Ag, Hf, Ta, W, Pt, Au, and the like that can be added to the nanocomposite magnet. Pb).
JP 59-46008 A JP-A-60-9852 Japanese Unexamined Patent Publication No. 1-7502 JP-A-3-261104 Japanese Patent No. 2727505 Japanese Patent No. 2727506 International Publication WO003 / 03403 JP 2002-175908 A Japanese Patent No. 2002-285301 Japanese Patent No. 3297676 R. Coehoorn et al., J. de Phys, C8, 1988, pages 669-670 WCChan, et.al. "THE EFFECTS OF REFRACTORY METALS ON THE MAGNETIC PROPERTIES OF α-Fe / R2Fe14B-TYPE NANOCOMPOSITES", IEEE, Trans. Magn. No. 5, INTERMAG. 99, Kyongiu, Korea pp.3265-3267, 1999

  The nanocomposite magnet disclosed in Patent Document 8 realizes a novel configuration in which fine soft magnetic phases are dispersed at the grain boundaries of the hard magnetic phase by the action of Ti. However, in a composition region where the rare earth concentration is lower than that of this type of nanocomposite magnet, a nanocomposite magnet exhibiting excellent magnetic properties has not been obtained unless the boron concentration is less than 10 atomic%.

  On the other hand, when the composition of both the rare earth element concentration and the boron concentration is less than 10 atomic%, there are a problem that the viscosity of the molten alloy increases and a problem that it is difficult to make the quenched alloy structure fine. The molten alloy cooling method (liquid quenching method) includes a melt spinning method in which the cooling roll is rotated at a relatively high speed and a strip casting method in which the cooling roll is rotated at a relatively low speed. However, it is said that this is a liquid quenching method suitable for mass production because a thick quenching alloy ribbon can be obtained relatively slowly.

  However, in order to mass-produce nanocomposite magnets with a rare earth element concentration reduced to 7 atomic% or less by a liquid quenching method with a relatively low cooling rate such as a strip casting method, a composition with a boron concentration exceeding 10 atomic% It is necessary to. Nevertheless, in the composition range in which the rare earth element is 4 to 7 atomic% and the boron concentration is 10 to 15 atomic%, when there is no addition of “Ti” or “Ti + Nb”, it is practical for use in a motor or the like. Not only the necessary coercive force HcJ of 400 kA / m or more could not be obtained, but also only a nanocomposite magnet having a poor demagnetization curve squareness was obtained.

  The present invention has been made in view of the above circumstances, and an object of the present invention is to provide a nanocomposite magnet that exhibits excellent magnet characteristics in a composition having a relatively large amount of rare earth elements and a relatively large amount of boron. It is to provide.

Nanocomposite magnet of the present invention, composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n (T is selected from the group consisting of Co and Ni One or more selected elements, R is a rare earth metal element, M is selected from the group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W One or more elements), and the composition ratios (atomic ratios) x, y, z, w, n, m, and p are 10 <x ≦ 15 atomic%, 4 ≦ y <7 atomic%, 0.5 ≦ z ≦ 8 atomic%, 0.01 ≦ w ≦ 6 atomic%, 0 ≦ n ≦ 10 atomic%, 0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5 are satisfied, And a hard magnetic phase having an R ₂ Fe ₁₄ B type crystal structure and a soft magnetic phase, and V is added to at least one of the coercive force and the maximum magnetic energy product. It is improved by 1% or more compared to the state where it is not.

In a preferred embodiment, the volume ratio of the hard magnetic phase having the R ₂ Fe ₁₄ B type crystal structure by addition of Ti and V is 40% or more.

In a preferred embodiment, the hard magnetic phase having the R ₂ Fe ₁₄ B type crystal structure has an average particle size of 10 nm to 200 nm, and the soft magnetic phase has an average particle size of 1 nm to 100 nm.

  In a preferred embodiment, the soft magnetic phase includes α-Fe and a ferromagnetic iron-based boride.

Method for producing a nanocomposite magnet rapidly solidified alloys of the present invention, composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n (T is Co And one or more elements selected from the group consisting of Ni and R, R is a rare earth metal element, M is Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W One or more elements selected from the group consisting of: x, y, z, w, n, m, and p are 10 <x ≦ 15 atomic%, 4 ≦ y <7 atomic%, 0.5 ≦ z ≦ 8 atomic%, 0.01 ≦ w ≦ 6 atomic%, 0 ≦ n ≦ 10 atomic%, 0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ A step of preparing a melt of an alloy satisfying 0.5 and a step of preparing a rapidly solidified alloy by quenching the melt.

  In a preferred embodiment, the molten metal is rapidly cooled using a strip casting method.

Method for producing a nanocomposite magnet powder of the present invention, composition formula _{(Fe 1-m T m)} -xyzwn (B 1-p C p) x R y Ti z V w M n (T consists of Co and Ni One or more elements selected from the group, R is a rare earth metal element, M is a group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W Selected one or more elements), and the composition ratio (atomic ratio) x, y, z, w, n, m, and p is 10 <x ≦ 15 atomic%, 4 ≦ y <7, respectively. Atomic%, 0.5 ≦ z ≦ 8 atomic%, 0.01 ≦ w ≦ 6 atomic%, 0 ≦ n ≦ 10 atomic%, 0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5 A step of preparing a satisfactory rapidly solidified alloy, a heat treatment of the rapidly solidified alloy, a hard magnetic phase having an R ₂ Fe ₁₄ B type crystal structure, and a soft magnetism A step of producing a nanocomposite magnet alloy including a phase, and a step of pulverizing the nanocomposite magnet alloy.

Method for producing a nanocomposite magnet of the present invention, composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n (T consists of Co and Ni One or more elements selected from the group, R is a rare earth metal element, M is a group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W Selected one or more elements), and the composition ratio (atomic ratio) x, y, z, w, n, m, and p is 10 <x ≦ 15 atomic%, 4 ≦ y <7, respectively. Atomic%, 0.5 ≦ z ≦ 8 atomic%, 0.01 ≦ w ≦ 6 atomic%, 0 ≦ n ≦ 10 atomic%, 0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5 A hard magnetic phase satisfying and having an R ₂ Fe ₁₄ B type crystal structure and a soft magnetic phase, wherein at least one of the coercive force and the maximum magnetic energy product is V Including a step of preparing a nanocomposite magnet powder that is improved by 1% or more compared to a state in which no is added, and a step of forming the nanocomposite magnet by molding the powder.

  According to the present invention, it is possible to produce a practical and mass-produceable nanocomposite magnet in a composition range in which a nanocomposite magnet exhibiting sufficiently excellent magnetic properties could not be produced by simultaneous addition of Ti and V. It becomes possible.

In the present invention, even if the amount of boron is relatively large, the formation of Ti and boron compounds can be suppressed by adding V. Therefore, an increase in the viscosity of the molten alloy is avoided, and α-Fe / R by strip casting is used. _This will pave the way for mass production of ₂ Fe ₁₄ B nanocomposite magnets.

  The present inventor has made the squareness of the demagnetization curve by simultaneously adding Ti and V to an alloy having a composition region containing 4 to 7 atomic% rare earth element and 10 to 15 atomic% boron and carbon. The inventors have found that an excellent nanocomposite magnet can be produced, and have come up with the present invention. The aforementioned Patent Documents 9 and 10 describe a large number of additive elements, but do not describe any unexpected effects due to simultaneous addition of Ti and V.

The addition of Ti exhibits the effect of suppressing the precipitation and growth of α-Fe in the process of rapidly cooling the molten alloy, as disclosed in Patent Document 8 by the present applicant. As a result of such addition of Ti, a structure in which the volume ratio of the hard magnetic phase having the R ₂ Fe ₁₄ B type crystal structure exceeds 50% and the soft magnetic phase is dispersed or thinly present at the grain boundaries of the hard magnetic phase is realized. To do.

Thus, although the addition of Ti has an excellent effect, it has been found that reducing the amount of rare earth element to further increase the magnetization of the nanocomposite magnet disclosed in Patent Document 8 deteriorates the magnet characteristics. Note that a hard magnetic phase having an R ₂ Fe ₁₄ B type crystal structure and a soft magnetic phase composed of α-Fe or iron-based boride are mixed in the same metal structure, and these constituent phases are magnetized by exchange interaction. In a nanocomposite magnet that is bonded to each other, it is considered effective to decrease the concentration of rare earth elements and thereby increase the volume ratio of α-Fe in order to increase the magnetization. This is because the saturation magnetization of α-Fe is higher than the saturation magnetization of the hard magnetic phase having the R ₂ Fe ₁₄ B type crystal structure.

However, when the rare earth element concentration is reduced to 7 atomic% or less in a composition system to which Ti is added, a coercive force H _cJ of 400 kA / m or more is not achieved unless the amount of added Ti is increased. It turned out that the squareness of a curve is also bad and the problem that a favorable magnetic characteristic cannot be obtained arises. On the other hand, if the amount of Ti is simply increased to solve such a problem, a large amount of nonmagnetic compounds in which Ti and boron are bonded will be precipitated, and the magnet characteristics will deteriorate. Occurred.

  Therefore, the present inventor conducted various experiments in which Ti and other metal elements were simultaneously added in a composition range in which the amount of rare earth elements was reduced to 7 atomic% or less and the boron concentration was relatively increased. As a result, by simultaneously adding Ti and V, the present inventors succeeded in obtaining a nanocomposite magnet that can be manufactured by strip casting while increasing the volume ratio of iron-based boride and α-Fe having high magnetization.

Nanocomposite magnet of the present invention is represented by the composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n, R 2 Fe 14 B type crystals It includes a hard magnetic phase having a structure and a soft magnetic phase. Here, T is one or more elements selected from the group consisting of Co and Ni, R is a rare earth metal element, M is Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, One or more elements selected from the group consisting of Ag, Ta, and W. The composition ratios (atomic ratios) x, y, z, w, n, m, and p each satisfy the following relational expression.

10 <x ≦ 15 atomic%,
4 ≦ y <7 atomic%,
0.5 ≦ z ≦ 8 atomic%,
0.01 ≦ w ≦ 6 atomic%,
0 ≦ n ≦ 10 atomic%,
0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5.

According to the present invention, by simultaneous addition of Ti and V, at least one of the coercive force and the maximum magnetic energy product is improved by 1% or more compared to the state without V addition. Moreover, the addition of Ti and V increases the volume ratio of the hard magnetic phase having the R ₂ Fe ₁₄ B type crystal structure to 40% or more. In the finally obtained structure, the average particle size of the hard magnetic phase having an R ₂ Fe ₁₄ B type crystal structure is 10 nm or more and 200 nm or less, and the average particle size of the soft magnetic phase is 1 nm or more and 100 nm or less. is there.

In a preferred embodiment of the present invention, an iron group containing Fe, B, C, R (one or more rare earth metal elements including Y), Ti and V as essential elements by a liquid quenching method such as a strip casting method. The molten alloy is preferably cooled in a reduced pressure atmosphere, thereby producing a quenched alloy containing a fine R ₂ Fe ₁₄ B type compound phase. Then, if necessary, a heat treatment is performed on the quenched alloy to crystallize the amorphous material remaining in the quenched alloy.

  The strip casting method is a method for producing a ribbon of a quenched alloy by bringing a molten alloy into contact with the surface of a cooling roll and cooling the molten alloy. In the present invention, the molten alloy is rapidly cooled and solidified by a cooling roll rotating at a higher speed than the conventional strip casting method. Compared with the melt spinning method, which uses a nozzle orifice to inject molten alloy onto the surface of the cooling roll, the strip casting method has a lower cooling rate but can produce a wide and relatively thick quenched alloy ribbon. Excellent in properties. The nanocomposite magnet of the present invention can also be produced by a melt spinning method in which a molten metal is injected onto a cooling roll using a nozzle.

According to the present invention, by adding Ti having a function of suppressing the precipitation and growth of α-Fe, the rare earth element concentration is set low and V is added, so that α- A nanocomposite magnet structure in which an appropriate amount of Fe or iron-based boride is present can be obtained. Because of the action of Ti and V, the R ₂ Fe ₁₄ B-type compound phase and α-Fe phase can be prevented from coarsening even if the cooling rate of the molten alloy is slowed down. Therefore, even after heat treatment, R ₂ Fe ₁₄ A high-performance nanocomposite magnet in which a soft magnetic phase such as an α-Fe phase having an average particle diameter of the B-type compound phase of 10 nm to 200 nm and an average particle diameter of 1 nm to 100 nm is dispersed can be obtained.

Conventionally, a nanocomposite magnet of a type composed of an R ₂ Fe ₁₄ B type compound phase and α-Fe exhibits excellent magnetization because the volume ratio of α-Fe having a high saturation magnetization is as high as 5 to 50%. However, when the boron concentration is lower than the range of the present invention and the cooling rate is slow as in the strip casting method, the crystal grain size becomes coarse and the nanocomposite magnet characteristics tend to deteriorate. However, according to the present invention, since the boron concentration has a high value exceeding 10 atomic%, it is possible to mass-produce excellent nanocomposite magnet characteristics by the strip casting method.

In the present invention, the volume ratio of the hard magnetic phase having the R ₂ Fe ₁₄ B type crystal structure is increased as much as possible by adding Ti and V while reducing the concentration of rare earth elements, while soft magnetism having high magnetization The squareness of the demagnetization curve is improved by increasing the volume ratio to some extent while suppressing the coarsening of the phase. The increase in magnetization is caused by the action of Ti and V to form a boride phase such as a ferromagnetic iron-based boride from the boron-rich nonmagnetic amorphous phase present in the rapidly solidified alloy, and the nonmagnetic remaining after the crystallization heat treatment It is thought that it contributes also to reducing the volume ratio of an amorphous phase.

In the present invention, iron having a saturation magnetization equivalent to or higher than the saturation magnetization of the R ₂ Fe ₁₄ B type compound phase is adjusted by adjusting the production conditions such as the alloy composition, the cooling rate of the alloy, and the heat treatment temperature. It becomes possible to produce a base boride and α-Fe. The generated soft magnetic phase is α-Fe (saturation magnetization 2.1T), Fe ₂₃ B ₆ (saturation magnetization 1.6T), or the like. Here, the saturation magnetization of R ₂ Fe ₁₄ B is about 1.6 T when R is Nd.

  The “amorphous phase” in this specification is not only a phase constituted only by a part in which the atomic arrangement is completely disordered, but also a crystallization precursor, a microcrystal (size: several nm or less), or an atom It also includes phases that partially contain clusters. Specifically, a phase in which the crystal structure cannot be clearly identified by X-ray diffraction or transmission electron microscope observation will be widely referred to as an “amorphous phase”. A structure that can clearly identify the crystal structure by X-ray diffraction or transmission electron microscope observation is referred to as a “crystal phase”.

  According to the inventor's experiment, only when Ti and V are added at the same time, the magnetization does not decrease compared to the case where Ti and Cr or Ti and Zr are added, and the squareness of the demagnetization curve is particularly good. It became a thing. From these facts, it is considered that the addition of Ti and V plays an important role in suppressing the formation of borides having low magnetization.

  When Ti and Nb were added simultaneously, the magnetization did not decrease, but there was a problem that the coercive force decreased.

  Thus, only when Ti and V are added at the same time, it is possible to obtain excellent nanocomposite magnet characteristics in a composition range where the rare earth element is 4 to 7 atomic% and the boron concentration is 10 to 15 atomic%.

In the present invention, not only a sufficient amount of boron is contained, but also carbon is contained as an essential element. For this reason, the kinematic viscosity of the molten metal is 5 × 10 ⁻⁶ m ² / sec or less, the molten metal flow is smoothed, and the adhesion between the molten metal and the cooling roll is improved. For this reason, the heat removal effect by the cooling roll is improved, and a good nanocomposite magnet can be produced even when the rotation speed of the cooling roll is low.

  For this reason, in a preferred embodiment of the present invention, it is possible to use a strip casting method in which the molten metal is poured directly from the chute (guide means) onto the cooling roll without controlling the flow rate of the molten metal by the nozzle orifice. This is higher in productivity and lower in manufacturing cost than in the case of the melt spinning method using a nozzle orifice. Thus, in order to make a molten R-Fe-B rare earth alloy amorphous in a cooling rate range that can be achieved even by a strip casting method, it is usually necessary to add 10 atomic% or more of B (boron). When a large amount of B is added in this way, a non-magnetic amorphous phase with a high B concentration remains in the metal structure even after the crystallization heat treatment of the quenched alloy, and a homogeneous fine crystal structure is obtained. Absent. As a result, the volume ratio of the ferromagnetic phase is reduced, leading to a decrease in magnetization. However, when Ti and V are added simultaneously as in the present invention, the phenomenon described above is observed, so that a highly magnetized iron-based boride is generated and the magnetization is improved.

[Reason for limiting composition]
When the total composition ratio x of B and C is 10 atomic% or less, when the cooling rate at the time of quenching is relatively slow, such as 10 ² ° C / second to 10 ⁵ ° C / second, the R ₂ Fe ₁₄ B type crystal phase It becomes difficult to produce a quenched alloy in which an amorphous phase is mixed, and a high coercive force cannot be obtained even if heat treatment is performed thereafter. Further, when the composition ratio x is 10 atomic% or less, iron-based borides exhibiting high magnetization are not generated. For this reason, x needs to exceed 10 atomic%. On the other hand, when the composition ratio x exceeds 15 atomic%, the volume ratio of the amorphous phase remaining after the crystallization heat treatment increases, and at the same time, the abundance ratio of α-Fe having the highest saturation magnetization in the constituent phases decreases. As a result, the residual magnetic flux density _Br decreases. From the above, the composition ratio x is preferably set so as to exceed 10 atomic% and not more than 15 atomic%. A more preferable range of the composition ratio x is more than 11 atomic% and not more than 14 atomic%.

  The ratio p of C to the whole of B and C is preferably in the range of 0.01 to 0.5 in terms of atomic ratio. In order to obtain the effect of C addition, if the ratio p of C is less than 0.01, the effect of C addition is hardly obtained. On the other hand, if p is too larger than 0.5, the amount of coarse α-Fe phase produced increases, resulting in a problem that the magnetic properties deteriorate. The lower limit of the ratio p is preferably 0.02, and the upper limit of p is preferably 0.25 or less. The ratio p is more preferably 0.05 or more and 0.15 or less.

R is one or more elements selected from the group of rare earth elements (including Y). If La or Ce is present, R (typically Nd) in the R ₂ Fe ₁₄ B phase is replaced with La or Ce, and the coercive force and squareness deteriorate, so La and Ce are substantially not included. Is preferred. However, when a very small amount of La or Ce (0.5 atomic% or less) exists as an unavoidable impurity, there is no problem in terms of magnetic characteristics. Therefore, it can be said that La and Ce are not substantially contained when 0.5 or less atomic percent La and Ce are contained.

More specifically, R preferably contains Pr or Nd as an essential element, and part of the essential element may be substituted with Dy and / or Tb. When the R composition ratio y is less than 4 atomic% of the total, the compound phase having the R ₂ Fe ₁₄ B type crystal structure necessary for the expression of the coercive force is not sufficiently precipitated, and a high coercive force H _cJ can be obtained. become unable. Further, when the composition ratio y of R is 7 atomic% or more, the amorphous forming ability is lowered, and the amount of soft magnetic phases such as α-Fe phase and Fe-B phase is reduced, resulting in a decrease in magnetization. Therefore, the composition ratio y of the rare earth element R is preferably adjusted to a range of 4 atomic% or more and less than 7 atomic%. A more preferable range of R is 5 atom% or more and less than 6 atom%.

The addition of Ti is adapted to exert an effect of precipitating and growing the hard magnetic phase in the rapidly solidified alloy molten earlier than the soft magnetic phases, improved and demagnetization curve of coercive force H _cJ and remanence B _r It contributes to the improvement of the squareness and increases the maximum energy product (BH) _max .

When the Ti composition ratio z is less than 0.5 atomic% of the whole, the effect of Ti addition is not sufficiently exhibited. On the other hand, the mole fraction z of Ti exceeds 8 atomic% of the total, the volume ratio of the amorphous phase that remains after the crystallization heat treatment is increased, tends to lead to reduction of the remanence B _r. From the above, the Ti composition ratio z is preferably in the range of 0.5 atomic% to 8 atomic%. The lower limit of the more preferable z range is 1.0 atomic%, and the upper limit of the more preferable z range is 6 atomic%. Furthermore, the upper limit of the preferable range of z is 5 atomic%.

Further, the higher the composition ratio x of Q composed of C and / or B, the easier it is to form an amorphous phase containing excessive Q (for example, boron). Therefore, it is preferable to increase the Ti composition ratio z. Ti has a strong affinity for B and is concentrated at the grain boundaries of the hard magnetic phase. If the ratio of Ti to B is too high, Ti may not enter the grain boundary but enter the R ₂ Fe ₁₄ B compound, which may reduce the magnetization. On the other hand, if the ratio of Ti to B is too low, many magnetic B-rich amorphous phases will be generated. According to experiments, it is preferable to adjust the composition ratio so as to satisfy 0.05 ≦ z / x ≦ 0.4, and it is more preferable to satisfy 0.1 ≦ z / x ≦ 0.35. More preferably, 0.13 ≦ z / x ≦ 0.3.

  V exhibits the effect of reducing the amount of Ti added to obtain excellent magnet characteristics. For this reason, since the production | generation of the compound of Ti and a boron can be suppressed, while improving magnetization, the effect of reducing a molten metal viscosity is also acquired. As a result, the rapid cooling process by the strip cast method is facilitated.

  When the composition ratio w of V is less than 0.01 atomic%, such V addition effect cannot be obtained, and when it exceeds 6 atomic%, the V—Fe—B compound is precipitated, so that the magnet characteristics deteriorate. Problem arises. For this reason, the composition ratio w is preferably set to 0.01 atomic% or more and 6 atomic% or less. A more preferable range of the composition ratio w is 0.1 atomic percent or more and 4 atomic percent or less, and a further preferable range is 0.5 atomic percent or more and 2 atomic percent or less.

  In order to obtain various effects, the metal element M may be added. M is one or more elements selected from the group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W. According to the experiment, the effect obtained by adding Ti and V was not greatly inhibited even when the metal element M was added as long as the amount did not exceed 10 atomic%. .

Fe occupies the remainder of the above-mentioned elements, but desired hard magnetic properties can be obtained even if a part of Fe is replaced with one or two transition metal elements (T) of Co and Ni. When substituted atom ratio of the number m of T exceeds 0.5 for Fe, it can not be obtained more high residual magnetic flux density B _r 0.7 T. For this reason, it is preferable that the substitution atom number ratio m is limited to a range of 0 to 0.5. By replacing part of Fe with Co, the squareness of the demagnetization curve is improved and the Curie temperature of the R ₂ Fe ₁₄ B phase is increased, so that the heat resistance is improved. A preferable range of the substitution atom number ratio m is 0.05 or more and 0.4 or less.

  Next, preferred embodiments of the present invention will be described in detail with reference to the drawings.

(Embodiment 1)
First, a first embodiment of the present invention will be described.

  In this embodiment, a rapidly solidified alloy is produced using the strip casting apparatus shown in FIG. In order to prevent the oxidation of the raw material alloy containing rare earth elements R and Fe which are easily oxidized, the quenching alloy is manufactured in an inert gas atmosphere. As the inert gas, a rare gas such as helium or argon or nitrogen can be used. Note that since nitrogen is relatively easy to react with the rare earth element R, it is preferable to use a rare gas such as helium or argon.

  The strip casting apparatus of FIG. 1 is disposed in a chamber (not shown) that can be reduced in pressure in an inert gas atmosphere. This strip casting apparatus includes a melting furnace 11 for melting alloy raw materials, a cooling roll 13 for rapidly cooling and solidifying molten alloy 12 supplied from the melting furnace 11, and a molten metal 12 from the melting furnace 11 to the cooling roll 13. And a chute (guide means) 14 for guiding.

  The melting furnace 11 can supply the molten metal 12 produced by melting the alloy raw material to the chute 14 with a substantially constant supply amount. This supply amount can be arbitrarily adjusted by controlling the operation of tilting the melting furnace 11.

The outer peripheral surface of the cooling roll 13 is formed 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 13 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 13 can be arbitrarily adjusted. The cooling rate by the strip casting apparatus can be controlled in the range of about 10 ³ ° C / second to about 10 ⁵ ° C / second by selecting the rotation speed of the cooling roll 13 or the like.

  The surface of the chute 14 for guiding the molten metal is inclined at an angle (inclination angle) α with respect to the horizontal direction, and the distance between the tip of the chute 14 and the surface of the cooling roll is kept at several mm or less. And the chute | shoot 14 is arrange | positioned so that the line which connects the front-end | tip part and the center of the cooling roll 13 may form the angle (alpha) (0 degrees <= alpha <= 90 degrees) with respect to the perpendicular direction. The angle α preferably satisfies the relationship of 10 ° ≦ α ≦ 55 °.

  The inclination angle β of the chute 14 is preferably 1 ° ≦ β ≦ 80 °, and more preferably satisfies the relationship of 5 ° ≦ β ≦ 60 °.

  The molten metal 12 supplied onto the chute 14 is supplied from the tip of the chute 14 to the surface of the cooling roll 13, and forms the molten metal paddle 6 on the surface of the cooling roll 13. The chute 14 can rectify the flow of the molten metal 12 by delaying the flow rate so as to temporarily store the molten metal 12 continuously supplied from the melting furnace 11 at a predetermined flow rate. If a damming plate that can selectively dam the flow of the molten metal surface portion of the molten metal 12 supplied to the chute 14 is provided, the rectifying effect can be further improved. By using the chute 14, the molten metal 12 can be supplied in a state where the cooling roll 13 is widened to a substantially uniform thickness over a certain width in the trunk length direction (axial direction: perpendicular to the paper surface). By adjusting the inclination angle β of the molten metal guide surface of the chute 14, the molten metal supply speed can be finely adjusted. The molten metal flows along the inclined guide surface of the chute 14 by its own weight, and has a momentum component parallel to the horizontal direction. As the inclination angle β of the chute 14 increases, the flow rate of the molten metal increases and the momentum also increases.

In addition to the above function, the chute 14 also has a function of adjusting the temperature of the molten metal 12 immediately before reaching the cooling roll 13. The temperature of the molten metal 12 on the chute 14 is desirably a temperature that is 100 ° C. or more higher than the liquidus temperature. This is because if the temperature of the molten metal 12 is too low, primary crystals such as TiB ₂ that adversely affect the alloy characteristics after quenching are locally nucleated and may remain after solidification. On the other hand, if the molten metal temperature is too low, the viscosity of the molten metal increases and splash is likely to occur. The molten metal temperature on the chute 14 can be controlled by adjusting the molten metal temperature at the time of pouring from the melting furnace 11 into the chute 14, the heat capacity of the chute 14 itself, etc. 1 (not shown) may be provided.

  The chute 14 in the present embodiment has a plurality of discharge portions provided at predetermined intervals along the axial direction of the cooling roll at the end arranged to face the outer peripheral surface of the cooling roll 13. ing. The width of the discharge part (width of one flow of the molten metal) is preferably set to 0.5 cm to 10.0 cm, and more preferably 0.7 cm to 4.0 cm. In this embodiment, the width of each molten metal flow in the discharge part is set to 1 cm. In addition, the width of the flow of the molten metal tends to spread in the lateral direction as the distance from the discharge portion increases. When a plurality of discharge portions are provided on the chute 14 to form a plurality of molten metal flows, it is preferable that adjacent molten metal flows do not contact each other.

  The molten metal 12 supplied onto the chute 14 contacts the cooling roll 13 having substantially the same width as each discharge portion along the axial direction of the cooling roll 13. Thereafter, the molten metal 12 that has come into contact with the cooling roll 13 with a predetermined discharge width moves on the circumferential surface of the roll as the cooling roll 13 rotates (as if it is pulled up to the cooling roll 13), and is cooled in this moving process. The In order to prevent molten metal leakage, the distance between the tip of the chute 14 and the cooling roll 13 is preferably set to 3 mm or less (particularly in the range of 0.4 to 0.7 mm).

  The gap between adjacent discharge portions is preferably set to 1 cm to 10 cm. Thus, if the molten metal contact part (molten cooling part) in the outer peripheral surface of the cooling roll 13 is isolate | separated into a several location, the molten metal discharged | emitted from each discharge part can be cooled effectively. As a result, a desired cooling rate can be achieved even when the amount of molten metal supplied to the chute 14 is increased.

  In addition, the form of the chute 14 is not limited to the above form, and may have a single discharge part, or the tapping width may be set larger. Further, a tubular opening may be provided at the tip (bottom) of the chute 14 and the molten metal 12 may be supplied to the surface of the cooling roll 13 through the tubular opening.

  The molten alloy 12 solidified on the outer peripheral surface of the rotating cooling roll 13 becomes a strip-shaped solidified alloy 15 and peels from the cooling roll 13. In the case of this embodiment, the molten metal flowing out from each of the plurality of discharge portions becomes a band having a predetermined width and solidifies. The peeled solidified alloy 15 is crushed and collected by a collection device (not shown).

  In this way, the strip casting method does not use a nozzle like the melt spinning method, and there is no problem of injection speed limitation due to the nozzle diameter and clogging of the molten metal due to solidification at the nozzle part, so it is suitable for mass production. Yes. Moreover, since the heating equipment of the nozzle part and the pressure control mechanism for controlling the molten metal head pressure are not required, the initial equipment investment and the running cost can be reduced.

  In the melt spinning method, the nozzle part cannot be reused, so the nozzle with high processing cost had to be made disposable, but in the strip casting method, the chute can be used repeatedly, so the running cost Is cheap.

  Furthermore, according to the strip casting method, the cooling roll can be rotated at a slower speed than the melt spinning method, and the amount of the molten alloy hot water can be increased, so that the quenched alloy ribbon can be thickened.

  By appropriately selecting the shape of the chute 14, the width and number of the molten metal discharge section, the molten metal supply speed, etc., the thickness (average value) and width of the obtained ribbon-like quenched alloy can be adjusted within an appropriate range. The width of the ribbon-like quenched alloy is preferably in the range of 15 mm to 80 mm. Moreover, if the thickness of the ribbon-like alloy is too thin, the bulk density becomes low, which makes it difficult to recover. If it is too thick, the cooling rate differs between the roll contact surface of the molten metal and the free surface (molten surface). Is not preferable because sufficient cannot be obtained. For this reason, the thickness of the ribbon-shaped alloy is preferably 50 μm or more and 250 μm or less, and more preferably 60 μm or more and 200 μm or less. A more preferable range of the thickness of the quenched alloy is 70 μm or more and 90 μm or less. In consideration of the packing density of the bonded magnet, the thickness of the quenched alloy preferably exceeds 80 μm. The quenching atmosphere at this time is preferably set to a reduced pressure state of 10 to 101.2 kPa.

  The nanocomposite magnet of the present invention can also be produced by a liquid quenching method (melt spinning method, atomizing method, etc.) below the strip casting method.

[Heat treatment]
In the present embodiment, the heat treatment is performed in an argon atmosphere. Preferably, the heating rate is 5 ° C./second to 20 ° C./second, and the temperature is maintained at 550 ° C. to 850 ° C. for 30 seconds to 20 minutes, and then cooled to room temperature. By this heat treatment, fine crystals of metastable phase precipitate and grow in the remaining amorphous phase, and a nanocomposite structure is formed.

When the heat treatment temperature is lower than 550 ° C., a large amount of amorphous phase remains even after the heat treatment, and the coercive force may not reach a sufficient level depending on the rapid cooling conditions. Further, when the heat treatment temperature exceeds 850 ° C., the grain growth of the respective constituent phases is significantly reduced residual magnetic flux density B _r, is deteriorated squareness of the demagnetization curve. For this reason, the heat treatment temperature is preferably 550 ° C. or more and 850 ° C. or less, but the more preferable heat treatment temperature range is 570 ° C. or more and 820 ° C. or less.

The heat treatment is preferably performed in an inert gas atmosphere such as Ar gas or N ₂ gas or in flowing air to prevent oxidation of the alloy. Moreover, you may heat-process in the vacuum of 0.1 kPa or less.

In the present invention, even if a soft magnetic phase such as α-Fe, Fe ₃ B, or Fe ₂₃ B ₆ is finally present, the average crystal grain size of the soft magnetic phase is equal to the average crystal grain size of the hard magnetic phase. Therefore, the soft magnetic phase and the hard magnetic phase are magnetically coupled by exchange interaction, and excellent magnetic properties are exhibited.

The average crystal grain size of the Nd ₂ Fe ₁₄ B type compound phase after the heat treatment needs to be 300 nm or less, which is a uniaxial crystal grain size, preferably 10 nm or more and 200 nm or less, and preferably 10 nm or more and 150 nm or less. Is more preferable. On the other hand, if the average crystal grain size of the ferromagnetic iron-based boride phase or α-Fe phase exceeds 100 nm, the exchange interaction acting between each constituent phase is weakened, and the squareness of the demagnetization curve is deteriorated. (BH) _max decreases. Usually, these phases do not become precipitates having a diameter smaller than 1 nm, but precipitates having a size of several nm. From the above, the average crystal grain size of the soft magnetic phase such as boride phase or α-Fe phase is preferably 1 nm or more and 100 nm or less, and more preferably 5 nm or more and 50 nm or less. More preferably, it is 5 nm or more and 30 nm or less. In order to exhibit excellent performance as an exchange spring magnet, the average crystal grain size of the Nd ₂ Fe ₁₄ B type compound phase is preferably larger than the average crystal grain size of the soft magnetic phase.

  Note that the quenched alloy ribbon may be roughly cut or pulverized before the heat treatment. If the obtained magnet is pulverized after heat treatment to produce magnet powder (magnetic powder), various bonded magnets can be produced from the magnetic powder by known processes. When producing a bonded magnet, the iron-based rare earth alloy magnetic powder is mixed with an epoxy resin or a nylon resin and formed into a desired shape. At this time, the nanocomposite magnetic powder may be mixed with other types of magnetic powder, for example, Sm—Fe—N magnetic powder or hard ferrite magnetic powder.

  Various rotary machines, such as a motor and an actuator, can be manufactured using the above-mentioned bond magnet.

  When the magnetic powder obtained by the method of the present invention is used for an injection-molded bonded magnet, it is preferably pulverized so that the average particle size is 200 μm or less, and more preferably the average particle size of the powder is 30 μm or more and 150 μm or less. . Moreover, when using for a compression-molded bond magnet, it is preferable to grind | pulverize so that a particle size may be 300 micrometers or less, and the average particle diameter of a more preferable powder is 30 micrometers or more and 250 micrometers or less. A more preferable range is 50 μm or more and 200 μm or less.

(Example 1)
B, C, Fe, Co, Ti, V, and Nd having a composition shown in Table 1 below and having a purity of 99.5% or more so that the total amount is 600 grams (g), and an alumina crucible It was thrown into. Thereafter, these alloy raw materials were melted in an argon (Ar) atmosphere at a pressure of 70 kPa by high-frequency heating to produce a molten alloy. After the molten metal temperature reached 1500 ° C., it was cast on a water-cooled copper mold to produce a flat alloy. Then, it weighed so that the total amount would be 15 grams, and put in a quartz crucible having an orifice having a diameter of 0.8 mm at the bottom. This was melted by high-frequency heating in an Ar gas atmosphere at a pressure of 1.33 to 47.92 kPa to form a molten metal at 1350 ° C. The molten metal surface was pressurized with Ar gas and sprayed with a cooling roll in an Ar atmosphere at room temperature to prepare a rapidly solidified alloy. The surface peripheral speed of the cooling roll was set to 10 m / sec, and the distance between the cooling roll surface and the nozzle selection position was 0.7 mm. As a result, a continuous rapidly solidified alloy ribbon (ribbon) having a width of 2 to 3 mm and a thickness of 20 to 50 μm was obtained.

ここで、Ｎｏ．１〜４の試料が本発明の実施例であり、Ｎｏ．５〜８の試料が比較例である。

Here, no. Samples 1 to 4 are examples of the present invention. Samples 5 to 8 are comparative examples.

  Next, this rapidly solidified alloy was kept in a temperature range of 600 to 800 ° C. for 6 to 8 minutes in an Ar atmosphere, and then cooled to room temperature. About the sample obtained in this way, the magnetic characteristic was measured using VSM. The measurement results are shown in Table 2.

表２からわかるように、実施例では、磁化が向上し、減磁曲線の角形性も良くなった結果、最大磁気エネルギー積（ＢＨ）_maxが比較例よりも向上した。特にＮｏ．１の試料とＮｏ．５の試料は、Ｖ添加の有無以外の点で等しい条件で作成されたものであるので、両者を比較すると、Ｎｏ．１の試料における残留磁束密度および最大磁気エネルギー積のいずれもがＮｏ．５の試料に比べて１％以上向上していることがわかる。

As can be seen from Table 2, in the example, the magnetization was improved and the squareness of the demagnetization curve was improved. As a result, the maximum magnetic energy product (BH) _max was improved as compared with the comparative example. In particular, no. No. 1 and No. 1 The sample No. 5 was prepared under the same conditions except for the presence or absence of V addition. Both the residual magnetic flux density and the maximum magnetic energy product in the sample No. 1 It can be seen that there is an improvement of 1% or more compared to the sample of 5.

FIG. 2 shows a graph of powder XRD after heat treatment (700 ° C., 6 minutes). The graph shows sample A (Nd _5.8 Fe _75.7 B ₁₂ C ₁ Ti ₄ V _1.5 ), sample B (Nd ₉ Fe ₇₃ B _12.6 C _1.4 Ti ₄ ), and sample C (Nd _4.5 Fe ₇₃ B _{18.5 in order from the top.} Data relating to Co ₂ Cr ₂ ) is shown.

Sample A is an example of the present invention, and it can be seen from the graph that the intensity peak of α-Fe is significantly observed as compared with other samples. Sample A is an “α-Fe / R ₂ Fe ₁₄ B nanocomposite magnet” containing Nd ₂ Fe ₁₄ B (hard magnetic phase) and α-Fe (soft magnetic phase) as main constituent phases. This phase also contains Fe ₂₃ B ₃ (soft magnetic phase).

On the other hand, Sample B is a nanocomposite magnet having a structure in which iron-based borides such as Fe ₃ B exist at the grain boundaries of Nd ₂ Fe ₁₄ B (hard magnetic phase). Further, sample C, the volume ratio of the Fe ₃ B is a major _{"Fe 3 B / Nd 2 Fe 14} B based nanocomposite magnet".

  The nanocomposite magnet according to the present invention is suitably used for various motors and actuators.

It is a figure which shows the structural example of the strip cast apparatus used suitably in this invention. It is a graph which shows the powder XRD regarding the Example and comparative example of this invention.

Explanation of symbols

11 Melting furnace 12 Molten alloy 14 Chute (Melt guiding means)
13 Cooling roll 15 Quenched alloy

Claims (6)

Composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n (T is at least one element selected from the group consisting of Co and Ni, R is a rare earth metal element, M is one or more elements selected from the group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W) And the composition ratio (atomic ratio) x, y, z, w, n, m, and p are respectively
10 <x ≦ 15 atomic%,
4 ≦ y <7 atomic%,
0.5 ≦ z ≦ 8 atomic%,
0.01 ≦ w ≦ 6 atomic%,
0 ≦ n ≦ 10 atomic%,
0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5
And a hard magnetic phase having an R ₂ Fe ₁₄ B type crystal structure, and a soft magnetic phase,
The hard magnetic phase has an average particle size of 10 nm to 200 nm, the soft magnetic phase has an average particle size of 1 nm to 100 nm,
The soft magnetic phase includes α-Fe and a ferromagnetic iron-based boride,
A nanocomposite magnet in which at least one of the coercive force and the maximum magnetic energy product is improved by 1% or more compared to a state where V is not added. _{2. The} nanocomposite magnet according to claim 1, wherein a volume ratio of the hard magnetic phase having the R ₂ Fe ₁₄ B type crystal structure is 40% or more. Composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n (T is at least one element selected from the group consisting of Co and Ni, R is a rare earth metal element, M is one or more elements selected from the group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W) And the composition ratio (atomic ratio) x, y, z, w, n, m, and p are respectively
10 <x ≦ 15 atomic%,
4 ≦ y <7 atomic%,
0.5 ≦ z ≦ 8 atomic%,
0.01 ≦ w ≦ 6 atomic%,
0 ≦ n ≦ 10 atomic%,
0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5
Preparing a molten alloy that satisfies
By rapidly cooling the molten metal to produce a rapidly solidified alloy;
A method for producing a rapidly solidified alloy for a nanocomposite magnet. The manufacturing method according to claim 3, wherein the rapid cooling of the molten metal is performed using a strip casting method. Composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n (T is at least one element selected from the group consisting of Co and Ni, R is a rare earth metal element, M is one or more elements selected from the group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W) And the composition ratio (atomic ratio) x, y, z, w, n, m, and p are respectively
10 <x ≦ 15 atomic%,
4 ≦ y <7 atomic%,
0.5 ≦ z ≦ 8 atomic%,
0.01 ≦ w ≦ 6 atomic%,
0 ≦ n ≦ 10 atomic%,
0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5
Preparing a rapidly solidified alloy satisfying
Heat treating the rapidly solidified alloy to produce a nanocomposite magnet alloy including a hard magnetic phase having an R ₂ Fe ₁₄ B type crystal structure and a soft magnetic phase;
Crushing the nanocomposite magnet alloy;
Only including,
The hard magnetic phase has an average particle size of 10 nm to 200 nm, the soft magnetic phase has an average particle size of 1 nm to 100 nm, and the soft magnetic phase includes α-Fe and ferromagnetic iron-based boron. A method for producing a nanocomposite magnet powder comprising forming a nanocomposite magnet powder containing a chemical . Composition formula _{(Fe 1-m T m)} 100-xyzwn (B 1-p C p) x R y Ti z V w M n (T is at least one element selected from the group consisting of Co and Ni, R is a rare earth metal element, M is one or more elements selected from the group consisting of Al, Si, Cr, Mn, Cu, Zn, Ga, Nb, Zr, Mo, Ag, Ta, and W) And the composition ratio (atomic ratio) x, y, z, w, n, m, and p are respectively
10 <x ≦ 15 atomic%,
4 ≦ y <7 atomic%,
0.5 ≦ z ≦ 8 atomic%,
0.01 ≦ w ≦ 6 atomic%,
0 ≦ n ≦ 10 atomic%,
0 ≦ m ≦ 0.5, and 0.01 ≦ p ≦ 0.5
And a hard magnetic phase having an R ₂ Fe ₁₄ B type crystal structure and a soft magnetic phase, the hard magnetic phase has an average particle size of 10 nm to 200 nm, and the soft magnetic phase The soft magnetic phase contains α-Fe and a ferromagnetic iron-based boride, and at least one of the coercive force and the maximum magnetic energy product is a state where V is not added. Preparing a nanocomposite magnet powder that is improved by 1% or more compared to
Forming a nanocomposite magnet by molding the powder;
A method for producing a nanocomposite magnet, comprising:

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