JP5071409B2 - Iron-based rare earth nanocomposite magnet and manufacturing method thereof - Google Patents

Description

  The present invention relates to an iron-based rare earth nanocomposite magnet and a method for producing the same. The present invention also relates to a rapidly solidified alloy for an iron-based rare earth nanocomposite magnet, an iron-based rare earth nanocomposite magnet powder, and a bond magnet containing the powder.

At present, a hard magnetic phase such as R ₂ Fe ₁₄ B phase (Nd ₂ Fe ₁₄ B type crystal phase, R is one or more rare earth elements) and a soft magnetic phase such as iron-based boride and α-Fe are magnetic. Nanocomposite magnets have been developed that have mechanically coupled tissue structures. Part of Fe may be substituted with Co and / or Ni, and part of B in the R ₂ Fe ₁₄ B phase may be substituted with C (carbon).

By adding Ti to an alloy for nanocomposite magnets having a specific composition, the applicant suppresses the precipitation and growth of α-Fe phase during the cooling process of the molten alloy, and crystal growth of R ₂ Fe ₁₄ B phase. Has been found to proceed with priority. The applicant of the present invention is a nanocomposite magnet having a metal structure in which an R ₂ Fe ₁₄ B phase is preferentially precipitated as an effect of added Ti and an iron-based boride phase is precipitated in the form of a film in the vicinity of a grain boundary. A configuration and a manufacturing method are disclosed in Patent Document 1.

  The Ti-containing nanocomposite magnet described in Patent Document 1 is a nanocomposite magnet mainly composed of an iron-based boride phase (Fe-B phase) as a soft magnetic phase (hereinafter referred to as "Fe-B nanocomposite magnet"). And a coercive force of about 500 to 1000 kA / m, which is very high as a quenching magnet and excellent in heat resistance. In addition, a rapidly solidified alloy can be manufactured by strip casting and the thickness of the slab can be increased, so the aspect ratio of the magnet powder shape (ratio of the size of the minor axis to the major axis size of the powder particles) is close to 1. It becomes possible to do. When the aspect ratio of the powder particles is close to 1, the fluidity of the compound mixed with the resin is improved, so that the magnetic powder is suitable for bond magnet molding (particularly injection molding and extrusion molding).

  Further, the present applicant is a nanocomposite magnet mainly composed of an α-Fe phase as a soft magnetic phase in Patent Documents 2 and 3 (hereinafter sometimes referred to as “α-Fe nanocomposite magnet”). Discloses a Ti-containing nanocomposite magnet having a high residual magnetic flux density.

Japanese Patent No. 3583116 WO2006 / 064794 International Publication Pamphlet WO2006 / 101117 International Publication Pamphlet

  In general, a nanocomposite magnet hardly exhibits an intrinsic coercive force, and a magnetizing magnetic field that is about four times the intrinsic coercive force is required to secure sufficient magnetic properties. Therefore, it cannot be said that the magnetizing characteristics of the nanocomposite magnet are good. The Fe-B-based nanocomposite magnet described in Patent Document 1 is an excellent magnet having a high intrinsic coercive force, but because of its high coercive force, the hard-to-adhere magnetism, which is a weak point of the nanocomposite magnet, becomes obvious and has a small diameter. -When applied to small motors for multi-information home appliances, etc., it has a technical problem of improving the magnetization characteristics.

  The Fe-B nanocomposite magnet described in Patent Document 1 has an aspect ratio close to 1 and excellent compound fluidity, but the iron-based boride phase is generally hard and difficult to grind. Since it is a crystalline phase, when it is pulverized for a bond magnet, the pulverizer tends to wear, and when the bonded magnet is formed, the mold tends to wear (hereinafter referred to as “metal mold”). "Wearing properties such as mold") may occur.

  To improve the magnetization characteristics, homogenization of the metal structure and refinement of the soft magnetic phase are considered effective. If the metal structure is inhomogeneous and an intrinsic coercive force distribution resulting therefrom is generated, the magnetization characteristics tend to be deteriorated. If the soft magnetic phase is not fine, the exchange coupling acting between the soft magnetic phase and the hard magnetic phase is weak, and the magnetic moment of the soft magnetic phase is likely to cause springback particularly in a low magnetization field. In order to prevent this, it is necessary to cause irreversible reversal of the magnetic moment of the hard magnetic phase by increasing the magnetizing magnetic field, but by reducing the crystal grain size of the soft magnetic phase, it works between the soft magnetic phase and the hard magnetic phase. It is considered that the exchange coupling force can be increased, the magnetization reversal of the magnetic moment of the soft magnetic phase in a low magnetization field can be suppressed, and the magnetization characteristics can be improved.

  In order to improve the above-mentioned magnetization characteristics while maintaining excellent compound fluidity, which is a feature of the nanocomposite magnet described in Patent Document 1, quenching with a device having a low alloy cooling rate such as a strip casting method is performed. Solidified alloy production is a prerequisite. This is to increase the thickness of the rapidly solidified alloy slab and bring the aspect ratio of the magnet powder shape after pulverization close to 1. However, the rapidly solidified alloy slab produced by the strip cast method has a large standard deviation in thickness, and thus it is difficult to obtain a homogeneous metal structure. Further, in the magnet described in Patent Document 1, since the iron-based boride phase, which is a soft magnetic phase, is deposited in the form of a film near the crystal grain boundary, there is a limit to reducing the crystal grain size of the soft magnetic phase.

  The problem of wearability of molds and the like is very difficult to improve as long as the main soft magnetic phase is an iron-based boride phase that is hard and difficult to grind.

  As described above, the Fe-B nanocomposite magnet has high intrinsic coercive force and excellent compound fluidity, but it has been difficult to improve the magnetizing characteristics and the wear properties of the mold.

  On the other hand, the α-Fe-based nanocomposite magnets described in Patent Documents 2 and 3 mainly have an α-Fe phase having a high saturation magnetization as a soft magnetic phase, and thus have a high residual magnetic flux density, but a hard magnetic phase. When the ratio of the soft magnetic phase is the same as that of the soft magnetic phase, the coercive force is lower than that of the Fe-B nanocomposite magnet due to the higher saturation magnetization, so there are few problems with the magnetization characteristics. Further, since the soft magnetic phase as a main component is not an iron-based boride phase that is hard and difficult to grind, there are few problems of wear such as molds. However, it is difficult for the α-Fe-based nanocomposite magnet to increase the coercive force, for example, it is difficult to have a practical high intrinsic coercive force of 640 kA / m or more.

  Further, in order to make the main soft magnetic phase an α-Fe phase, it is necessary to lower the B concentration of the alloy composition. However, if the B concentration is low, the amorphous forming ability of the alloy is significantly reduced. For this reason, in order to produce a magnet having good magnetic properties, it is an essential condition to produce a rapidly solidified alloy using an apparatus having a high molten metal cooling rate such as a melt spinning method. When a rapidly solidified alloy was produced with an apparatus having a low cooling rate of the molten alloy, a desired metal structure was not obtained, and good magnetic properties were not obtained. When a rapidly solidified alloy is produced using an apparatus having a high molten metal cooling rate such as a melt spinning method, the thickness of the rapidly solidified alloy is reduced, and the magnet powder after pulverization has a flat shape, that is, the above aspect ratio is small, so that it is high. It is difficult to obtain compound fluidity.

  As described above, the α-Fe-based nanocomposite magnet has fewer problems of wear characteristics such as magnetization characteristics and molds than the Fe-B-based nanocomposite magnet, but has the Fe-B-based nanocomposite magnet. It was difficult to have a high intrinsic coercivity and excellent compound fluidity.

That is, the nanocomposite magnet is required to solve the following four technical problems, but the conventional nanocomposite magnet cannot realize all of them simultaneously.
(1) High intrinsic coercive force of 640 kA / m or more (2) Excellent compound fluidity (3) Excellent magnetizing properties (4) Improvement of wear properties such as molds

  The present invention has been made to solve the above-mentioned problems, and the object of the present invention is the characteristics of a conventional Ti-containing nanocomposite magnet, (1) high intrinsic coercivity, and (2) excellent compound. An object of the present invention is to provide an iron-based rare earth nanocomposite magnet in which (3) magnetizing characteristics and (4) wear resistance such as a mold are improved while maintaining fluidity.

  Another object of the present invention is to provide a rapidly solidified alloy for the iron-based rare earth nanocomposite magnet, a powder of the iron-based rare earth nanocomposite magnet, a bonded magnet containing the powder, and the like.

Iron-based rare earth nanocomposite magnet of the present invention, the composition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T is selected from the group consisting of Fe, or, Co and Ni Transition metal elements including one or more elements and Fe, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo, Ag, Hf , Ta, W, Pt, Au, and Pb), and the composition ratios x, y, z, t, m, and p are respectively 7 ≦ x ≦ 9 atomic%, 6.5 ≦ y ≦ 9 atomic%, 2 ≦ z ≦ 5 atomic%, 0.5 ≦ t ≦ 3 atomic%, 4 ≦ z + t ≦ 7 atomic%, 0.5 ≦ z / (z + t) ≦ 0.95, 0 ≦ m ≦ 5 atomic%, and 0 ≦ p ≦ 0.5 are satisfied.

In a preferred embodiment, the soft magnetic phase includes an α-Fe phase having an average crystal grain size of 1 nm to 50 nm, and the hard magnetic phase includes a metal structure including an R ₂ Fe ₁₄ B phase having an average crystal grain size of 5 nm to 100 nm. Yes.

In a preferred embodiment, it has permanent magnet characteristics of residual magnetic flux density B _r ≧ 0.75 T, maximum energy product (BH) _max ≧ 100 kJ / m ³ , and intrinsic coercive force H _cJ ≧ 640 kA / m.

  The iron-based rare earth nanocomposite magnet powder of the present invention is a powder of any of the above iron-based rare earth nanocomposite magnets, having an average particle size of 10 μm or more and 300 μm or less, and the length of the powder particles The ratio of the minor axis size to the axial size is not less than 0.3 and not more than 1.

  The bonded magnet of the present invention includes the iron-based rare earth nanocomposite magnet powder and a binder that binds the iron-based rare earth nanocomposite magnet powder.

  In a preferred embodiment, the magnetization magnetic field required for 90% magnetization when compression-molded and the magnetization field of 4.8 MA / m is 100% is 1.5 MA / m or less.

Iron-based rare-earth rapidly solidified alloy for a nanocomposite magnet of the present invention, the composition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T consists of Fe or, Co and Ni A transition metal element including one or more elements selected from the group and Fe, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo And one or more metal elements selected from the group consisting of Ag, Hf, Ta, W, Pt, Au, and Pb), and the composition ratios x, y, z, t, m, and p are respectively 7 ≦ x ≦ 9 atomic%, 6.5 ≦ y ≦ 9 atomic%, 2 ≦ z ≦ 5 atomic%, 0.5 ≦ t ≦ 3 atomic%, 4 ≦ z + t ≦ 7 atomic%, 0.5 ≦ z /(Z+t)≦0.95, 0 ≦ m ≦ 5 atomic%, 0 ≦ p ≦ 0.5, and the average thickness is 60 μm or more and 300 μm or less .

Method of manufacturing an iron-based rare earth nanocomposite magnet of the present invention, the composition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T is Fe or a group consisting of Co and Ni A transition metal element including one or more elements selected from Fe and Fe, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo, One or more metal elements selected from the group consisting of Ag, Hf, Ta, W, Pt, Au, and Pb), and the composition ratios x, y, z, t, m, and p are respectively 7 ≦ x ≦ 9 atomic%, 6.5 ≦ y ≦ 9 atomic%, 2 ≦ z ≦ 5 atomic%, 0.5 ≦ t ≦ 3 atomic%, 4 ≦ z + t ≦ 7 atomic%, 0.5 ≦ z / (Z + t) ≦ 0.95, 0 ≦ m ≦ 5 atomic%, forming a molten alloy satisfying 0 ≦ p ≦ 0.5, and quenching the molten metal And a cooling step of forming a rapidly solidified alloy having a thickness of 60 μm to 300 μm, and a step of crystallizing the rapidly solidified alloy by heat treatment to produce an alloy having permanent magnet characteristics.

  In a preferred embodiment, the quenching method is a strip casting method.

Method of manufacturing an iron-based rare earth nanocomposite magnet powder of the present invention, the composition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T consists of Fe or, Co and Ni A transition metal element including one or more elements selected from the group and Fe, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo And one or more metal elements selected from the group consisting of Ag, Hf, Ta, W, Pt, Au, and Pb), and the composition ratios x, y, z, t, m, and p are respectively 7 ≦ x ≦ 9 atomic%, 6.5 ≦ y ≦ 9 atomic%, 2 ≦ z ≦ 5 atomic%, 0.5 ≦ t ≦ 3 atomic%, 4 ≦ z + t ≦ 7 atomic%, 0.5 ≦ z /(Z+t)≦0.95, 0 ≦ m ≦ 5 atomic%, 0 ≦ p ≦ 0.5, forming a molten alloy, and Cooling by a cooling method, thereby forming a rapidly solidified alloy having a thickness of 60 μm or more and 300 μm or less, a step of crystallizing the rapidly solidified alloy by heat treatment to produce an alloy having permanent magnet characteristics, Pulverizing to form a powder having an average particle size of 10 μm or more and 300 μm or less and a ratio of the minor axis size to the major axis size of the powder particles of 0.3 or more and 1 or less.

  In a preferred embodiment, the quenching method is a strip casting method.

  According to the present invention, the compound has a homogeneous metal structure in which the α-Fe phase is finely dispersed and has a high intrinsic coercive force of 640 kA / m or more, while having excellent compound fluidity, mold wear resistance, An iron-based rare earth nanocomposite magnet having magnetic properties can be provided.

It is a figure which shows the structural example of the strip casting apparatus used suitably by embodiment of this invention. 10 is a graph showing the magnetization characteristics of Example 5 and Comparative Example 9. It is the graph which expanded a part of FIG.

  The present inventors have prepared a rapidly solidified alloy under conditions where the molten metal quenching rate is low, such as a strip casting method, but if the alloy composition including the additive element is appropriately controlled, α having a homogeneous metal structure can be obtained. -Fe-based nanocomposite magnets, that is, nanocomposite magnets having a high intrinsic coercive force and excellent magnetization characteristics could be produced, and studies were repeated. It should be noted that producing a rapidly solidified alloy under a condition where the alloy melt quenching speed is low means that a thick rapidly solidified alloy is produced. When such a thick rapidly solidified alloy is pulverized, a magnet powder having an excellent compound fluidity with an aspect ratio close to 1 can be produced. As a result, even in an alloy composition with a low B concentration, if Zr is added at a relatively high composition ratio and a small amount of Ti is added, a homogeneous metal structure in which the α-Fe phase is finely dispersed can be obtained, and 640 kA / The present inventors have found that an iron-based rare earth nanocomposite magnet having a high intrinsic coercive force of m or more and excellent magnetization characteristics can be realized.

In the present invention, first, to produce a rapidly solidified alloy by cooling the alloy melt expressed by a composition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m. Here, T is a transition metal element containing Fe and one or more elements selected from the group consisting of Fe or Co and Ni, and R is one or more rare earth elements, and M is Al, Si, V, One or more metal elements selected from the group consisting of Cr, Mn, Cu, Zn, Ga, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb. The composition ratios x, y, z, t, m, and p in the above composition formula are respectively 7 ≦ x ≦ 9 atomic%, 6.5 ≦ y ≦ 9 atomic%, 2 ≦ z ≦ 5 atomic%, 0 5 ≦ t ≦ 3 atomic%, 4 ≦ z + t ≦ 7 atomic%, 0.5 ≦ z / (z + t) ≦ 0.95, 0 ≦ m ≦ 5 atomic%, 0 ≦ p ≦ 0.5 Yes.

  In the present invention, in order to finally obtain an α-Fe-based nanocomposite magnet, the total composition ratio x of B and C is 7 atomic% or more and 9 atomic% or less, and the composition ratio y of the rare earth element R is 6.5 atoms. % To 9 atomic%. By using an α-Fe-based nanocomposite magnet, the magnet is mainly composed of an α-Fe phase rather than an iron-based boride phase that is hard and hard to pulverize as a soft magnetic phase. Improved.

In the present invention, in a composition range where the composition ratio x of B and C is low so that an α-Fe-based nanocomposite magnet can be obtained, a relatively thick rapid solidification is performed by a method with a low alloy cooling rate such as a strip cast method. Even when the alloy is produced, by adding Zr to increase the amorphous forming ability, the rapidly solidified alloy can be made amorphous or a structure in which a fine R ₂ Fe ₁₄ B phase exists in the amorphous. Powder particles produced by pulverizing a thick rapidly solidified alloy have an advantage that the aspect ratio of the shape is close to 1, that is, a shape close to a sphere, and the powder flowability is improved. Therefore, (2) excellent compound fluidity can be obtained.

However, the addition of Zr, when produced a thick rapidly solidified alloy, fine R ₂ Fe ₁₄ B phase less to precipitate during rapid solidification, further, since Zr is also serves to suppress the grain growth by heat treatment, R ₂ When heat treatment is performed at a temperature at which the Fe ₁₄ B phase does not become coarse or non-uniform, a relatively fine R ₂ Fe ₁₄ B phase is surrounded by an amorphous phase, resulting in a low residual magnetic flux density. Become. When the heat treatment temperature is further increased to increase the residual magnetic flux density by crystallization of this amorphous phase, the precipitated R ₂ Fe ₁₄ B phase grows at once, and the metal structure becomes coarse and non-uniform. Magnet characteristics such as coercive force and magnetizing characteristics deteriorate. Since this grain growth occurs abruptly at a certain temperature, it is very difficult to control the magnet structure after the heat treatment by adjusting the crystallization heat treatment temperature. Also, if the amount of Zr added is adjusted to increase the fine R ₂ Fe ₁₄ B phase in the rapidly solidified alloy in order to avoid the above-mentioned grain growth, this time, the R ₂ Fe ₁₄ B phase in the rapidly solidified alloy is increased. The α-Fe phase is likely to precipitate, α-Fe coarsening is likely to occur after heat treatment, and the metal structure becomes non-uniform, resulting in deterioration of magnet characteristics such as coercive force and magnetization characteristics.

  To solve this problem, a small amount of Ti was added. Surprisingly, the microstructure of the metal structure was homogenized, and (1) a high intrinsic coercive force of 640 kA / m or more (3) excellent magnetization characteristics. In other words, a nanocomposite magnet in which all the problems (1) to (4) were solved could be obtained. It has also been found that such excellent characteristics cannot be realized if the added amount of Zr and the added amount of Ti deviate from the optimum ranges.

Since Ti has a high affinity with B, when Ti is added, a region with a relatively low B concentration and a large amount of Fe is formed in the alloy in the process of cooling the molten alloy. In this region, the effective B amount contributing to the amorphous formation is substantially reduced, and a fine R ₂ Fe ₁₄ B phase is likely to precipitate in the rapidly solidified alloy. In this alloy system, when the amount of effective B that contributes to amorphous formation decreases, the α-Fe phase preferentially precipitates in the rapidly solidified alloy over the R ₂ Fe ₁₄ B phase, but the α-Fe phase precipitates. Is suppressed by the effect of Ti addition, and the rapidly solidified alloy has a fine structure in which a fine R ₂ Fe ₁₄ B phase and an amorphous phase are uniformly mixed, or a fine R ₂ Fe ₁₄ B phase, a fine α-Fe phase, And it becomes a fine structure in which the amorphous phase is uniformly mixed. In this rapidly solidified alloy, a sufficient amount of fine R ₂ Fe ₁₄ B phase is present, and precipitation and growth of α-Fe phase are suppressed by the Ti addition effect, so that when it is completely crystallized by heat treatment, ideal magnet tissues relatively small alpha-Fe phase is present than R ₂ Fe ₁₄ B phase at the grain boundary triple point of appropriate size of the R ₂ Fe ₁₄ B phase is formed.

  Hereinafter, the effects of adding Zr and Ti and the respective composition ranges will be described in more detail.

  As described above, an alloy composition having a low total composition ratio x of B and C as in the present invention generally has a low ability to form an amorphous material of a rapidly solidified alloy, and thus a production of a rapid molten metal quenching rate such as a melt spinning method. It is difficult to obtain good magnetic characteristics unless an apparatus is used. In the present invention, by adding Zr, the amorphous forming ability of a rapidly solidified alloy is improved, and a rapidly solidified alloy having a desired metal structure can be produced even by a manufacturing method with a slow alloy quenching rate such as a strip cast method. It is said.

In the present invention, the R ₂ Fe ₁₄ B phase preferentially precipitates during rapid solidification and the α-Fe phase precipitates during crystallization heat treatment, but the crystal grains of the R ₂ Fe ₁₄ B phase during crystallization heat treatment by addition of Zr Since the growth is suppressed, the α-Fe phase can be finely precipitated, and further, the metal structure can be homogeneously refined, so that good magnetic characteristics and magnetization characteristics can be obtained. When the composition ratio z of Zr is less than 2 atomic%, the effect of adding Zr does not sufficiently appear. On the other hand, the mole fraction z of Zr is more than 5 atomic%, fine R ₂ Fe ₁₄ B phase not sufficiently precipitated during rapid solidification, R ₂ Fe ₁₄ B phase precipitated in the crystallization heat treatment to increase grain growth As a result, the α-Fe phase is difficult to be generated, and good magnetic properties cannot be obtained, and the viscosity of the molten alloy is increased, making it difficult to supply the molten alloy to the cooling roll. 2 ≦ z ≦ 4.8 atomic% is preferable, and 2.5 ≦ z ≦ 4.8 atomic% is more preferable.

Ti forms a nonmagnetic compound such as Ti-B, Ti-C, Ti-B-C, etc., and exists at the grain boundary of each crystal grain, so that the R ₂ Fe ₁₄ B phase and the α-Fe phase between the crystal grains And serves to reduce exchange interaction between crystal grains of R ₂ Fe ₁₄ B. Ti has an effect of suppressing the precipitation and growth of the α-Fe phase. The composition ratio y of the rare earth element R in the present invention is 9 atomic% or less, and this composition ratio y is smaller than the stoichiometric composition (R: 11.8 atomic%) of the R ₂ Fe ₁₄ B type compound phase. For this reason, unless Ti is added, α-Fe is likely to be formed as the primary crystal when the molten alloy is rapidly cooled. Although α-Fe can be grown by subsequent heat treatment, it cannot be made fine or extinguished. For this reason, in order to realize a finer structure, it is preferable not to generate as much as possible during rapid cooling, and even if it is generated, it is necessary to suppress the size to several nanometers. Ti is one of the elements indispensable for the present invention that needs to suppress the production of α-Fe. According to experiments, when Zr was added, the most preferable magnet characteristics were obtained when 0.5 ≦ t ≦ 3 atomic%.

  Since the composition ratio of Zr is z and the composition ratio of Ti is t, the total composition ratio of Zr and Ti is z + t. When z + t is less than 4 atomic%, the above effects are not sufficiently exhibited, and when it exceeds 7 atomic%, the residual magnetic flux density Br is significantly reduced. It is preferable that 4.5 ≦ z + t ≦ 6 atomic%, and it is more preferable that 5 ≦ z + t ≦ 5.5 atomic%.

  The effects of Zr and Ti are as described above, but in order to exert a synergistic effect, the atomic ratio of Zr to the whole of Zr and Ti, that is, the size of z / (z + t) is important. When z / (z + t) is low, the effect of Zr addition does not appear, and the amorphous forming ability of the rapidly solidified alloy is lowered, so that it is difficult to obtain good magnetic properties under a slow alloy cooling rate. Accordingly, the atomic ratio of Zr is 0.5 ≦ z / (z + t) ≦ 0.95, and preferably 0.5 ≦ z / (z + t) ≦ 0.82.

As described above, when Ti is added, a portion having a low B concentration distribution is generated in the alloy, and the substantial effective B amount in the portion is reduced, which contributes to the precipitation of a fine R ₂ Fe ₁₄ B phase. Since Ti is easily combined with B, Ti—B compounds such as TiB ₂ are likely to be precipitated in the early stage of quenching of the melt. Since the Ti-B compound is a non-magnetic phase, if it precipitates in the alloy, it causes a decrease in magnetization and non-uniformity of the metal structure. By substituting 1 to 50% of the B content with C, the amount of Ti-B compound precipitated in the molten metal can be suppressed, and the balance of reduction in the effective B amount can be adjusted. The C substitution amount p is preferably 1% or more in terms of the number ratio of atoms with respect to the whole of B and C. Substitution with C up to 50% of B does not affect the magnetic properties and metal structure, so the upper limit of the substitution amount p is 50%. The range of p is more preferably 0.02 ≦ p ≦ 0.5, and further preferably 0.04 ≦ p ≦ 0.3.

  The total composition ratio x of B and C is 7 atom% or more and 9 atom% or less. When the composition ratio x exceeds 9 atomic%, not the α-Fe but the iron-based boride phase that is hard to be pulverized is precipitated as the soft magnetic phase, and the wear properties of the mold and the like are not improved. Further, when the composition ratio x is less than 7 atomic%, the amorphous forming ability is greatly reduced, so that a homogeneous fine metal structure is not obtained, and good magnetization characteristics and magnetic characteristics cannot be obtained. The composition ratio x preferably satisfies the relationship of 7.3 ≦ x ≦ 8.5 atomic%, and more preferably satisfies the relationship of 7.5 ≦ x ≦ 8.3 atomic%.

The rare earth element R is one or more rare earth metals. When the composition ratio y of the rare earth element R is less than 6.5 atomic%, H _cJ of 640 kA / m or more cannot be obtained and a practical permanent magnet cannot be obtained. If the composition ratio y exceeds 9 atomic%, α-Fe which is a soft phase does not precipitate in the rapidly solidified alloy, and it becomes difficult to obtain good magnetic properties. It is preferable that 7 ≦ y ≦ 8 atomic%. Note that R may include La and Ce, which are inevitable in manufacturing, but the maximum amount is within 5% of the R amount.

T, which is substantially Fe, occupies the residual content 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 of Co and Ni. If the substitution amount for Fe exceeds 50%, a residual magnetic flux density of 0.75 T or more cannot be obtained, so the substitution amount is limited to a range of 50% or less. Incidentally, the substitution with Co improves the squareness of the demagnetization curve and increases the Curie temperature of the R ₂ Fe ₁₄ B phase, thereby improving the heat resistance. Furthermore, stable quenching can be maintained because the viscosity of the molten alloy decreases during quenching by strip casting. Preferably, the substitution amount of Co is 0.5% or more and less than 15%.

  Further, by adding one or more additive elements M of Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb. In addition to improving the magnetic characteristics, the effect of expanding the optimum heat treatment temperature range can be obtained. If the composition ratio m of the element M exceeds 5 atomic%, the magnetization is lowered, so that it is limited to 0 ≦ m ≦ 5 atomic%. A preferable range of the composition ratio m is 0.1 ≦ m ≦ 4 atomic%. Nb is a particularly effective element in magnetic powder production because high magnetic properties can be obtained in a wide cooling rate range.

  Hereinafter, a preferred embodiment of a method for producing an iron-based rare earth nanocomposite magnet according to the present invention will be described.

  In a preferred embodiment of the present invention, a rapidly solidified alloy is produced by cooling a molten metal having the above composition by a strip casting method. Then, if necessary, heat treatment is performed on the rapidly solidified alloy to crystallize the amorphous material remaining in the rapidly solidified alloy.

  The strip casting method is a method for producing a rapidly solidified alloy ribbon by bringing a molten alloy into contact with the surface of a cooling roll and cooling the molten alloy. The strip casting method has a lower cooling rate because the amount of molten alloy supplied to the surface of the cooling roll is larger than that of the melt spinning method, but it is excellent in mass productivity and finally has a particle size with an aspect ratio close to 1. Easy to get.

  First, the configuration of the rapid cooling apparatus used in the present embodiment will be described with reference to FIG. In this embodiment, a rapidly solidified alloy is produced using the strip casting apparatus shown in FIG. In order to prevent oxidation of a raw material alloy containing rare earth elements R and Fe that are easily oxidized, a rapidly solidified alloy is produced 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 evacuated in an inert gas atmosphere. The strip casting apparatus includes a melting furnace 1 for melting an alloy raw material, a molten alloy 2 supplied from the melting furnace 1, a cooling roll 3 for rapidly cooling and solidifying the molten alloy 2, and cooling from the melting furnace 1. A chute (guide means) 4 for guiding the molten alloy 2 to the roll 3 is provided. In addition to the shape shown in the figure, the chute 4 may be shaped to store the molten alloy in a portion close to the cooling roll 3 for the purpose of controlling the flow rate of the molten alloy and rectifying the flow of the molten metal. The rectification effect is enhanced by providing a weir before storage. Furthermore, by providing a hole in the lower portion of the storage portion, it is possible to make the amount of the molten alloy 2 supplied to the cooling roll 3 constant while having the above-described effects, and the thickness and width of the rapidly solidified alloy ribbon can be increased. It becomes easy to control constant. The above object can also be achieved by attaching a groove on the slit to the chute 4 and supplying the molten alloy 2 to the cooling roll 3 along the groove. The number of holes and grooves can be arbitrarily changed (both not shown).

In addition to the above function, the chute 4 also has a function of adjusting the temperature of the molten alloy 2 immediately before reaching the cooling roll 3. 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 such as TiB ₂ that adversely affect the alloy properties after rapid cooling 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 4 can be controlled by adjusting the molten metal temperature at the time of pouring from the melting furnace 1 into the chute 4 or the heat capacity of the chute 4 itself. 1 (not shown) may be provided.

The outer peripheral surface of the cooling roll 3 is made of a material having good thermal conductivity such as copper, and can be made of, for example, Al alloy, copper alloy, carbon steel, brass, W, Mo, or bronze. However, from the viewpoint of mechanical strength and economy, it is preferable to form Cu, Fe, or an alloy containing Cu or Fe. If the cooling roll is made of a material other than Cu or Fe, the peelability of the rapidly solidified alloy with respect to the cooling roll is deteriorated, which may cause the rapidly solidified alloy to be wound around the roll. The diameter of the cooling roll 3 is, for example, 300 to 500 mm. The water cooling capacity of the water cooling device provided in the cooling roll 3 is calculated and adjusted according to the solidification latent heat per unit time and the amount of hot water. The cooling roll 3 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 3 can be arbitrarily adjusted. The cooling rate by the strip casting apparatus can be controlled in the range of about 10 ² ° C / second to about 10 ⁷ ° C / second by selecting the rotation speed of the cooling roll 3 and the like. In the present embodiment, the cooling rate is preferably about 10 ³ ° C / second to about 10 ⁶ ° C / second. The rotation speed of the cooling roll 3 is preferably 5 m / s to 20 m / s, although it depends on the amount of tapping water.

  The strip casting apparatus of FIG. 1 may include a scraper gas injector (not shown) that facilitates peeling of the ribbon-like rapidly solidified alloy 5 supplied to the cooling roll 3.

  The molten alloy 2 solidified on the outer peripheral surface of the rotating cooling roll 3 becomes a ribbon-like rapidly solidified alloy 5 and peels from the cooling roll 3. In the case of this embodiment, the molten metal flowing out from the discharge part becomes a band having a predetermined width and solidifies. The rapidly solidified alloy 5 that has been peeled off may be recovered in a recovery device (not shown), or may be recovered as a pulverized alloy by providing a pulverization device (not shown) in front of the recovery device.

  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 molten alloy hot water can be increased, so that the rapidly solidified alloy ribbon can be thickened. In the melt spinning method, it is possible to increase the thickness of the rapidly solidified alloy ribbon by slowing the rotational speed of the cooling roll as described above, but the cooling speed of the alloy at the slow rotational speed as described above. In addition, it is difficult to control the thickness of the ribbon as compared to the strip casting method.

  In the present invention, the average thickness of the ribbon-like rapidly solidified alloy is set to be 60 μm or more and 300 μm or less. A more preferable range of the average thickness of the rapidly solidified alloy is 65 μm or more and 200 μm or less. In consideration of the packing density of the bonded magnet, the thickness of the rapidly solidified alloy is preferably more than 70 μm.

  When a rapidly solidified alloy having the above average thickness is pulverized to produce a powder having an average particle size of 10 μm or more and 300 μm or less, the aspect ratio of each powder particle (ratio of the size in the short axis direction to the size in the short axis direction) Of 0.3 or more and 1 or less can be formed.

  The pulverization can be performed separately in a coarse pulverization step and a fine pulverization step. Part or all of these pulverization steps can be performed at any timing before or after the “heat treatment” described below.

[Heat treatment]
In the present embodiment, the heat treatment is performed in an inert gas atmosphere such as an argon atmosphere. The heat treatment may be performed in a vacuum of 0.1 kPa or less. Preferably, the rate of temperature rise is 0.5 ° C./second or more and 10 ° C./second or less, and the temperature is maintained at 500 ° C. or more and 850 ° C. or less for 30 seconds or more and 30 minutes or less, and then cooled to room temperature. By this heat treatment, the α-Fe phase is finely precipitated while the R ₂ Fe ₁₄ B phase is sufficiently crystallized.

  At a heat treatment temperature of less than 500 ° C., crystallization from amorphous does not occur and desired magnetic properties cannot be obtained. On the other hand, if the temperature exceeds 850 ° C., the exchange interaction between the phases is weakened due to the growth of each crystal grain, so that a residual magnetic flux density of 0.75 T or more cannot be obtained. Preferably it is 550 degreeC or more and 850 degrees C or less, More preferably, 600 degreeC or more and 800 degrees C or less are good. As for the rate of temperature rise during the crystallization heat treatment, if it is less than 0.5 ° C./second, a uniform fine metal structure cannot be obtained and a residual magnetic flux density of 0.75 T or more cannot be obtained. In addition, there is no particular limitation to obtain a uniform fine metal structure at the upper limit of the temperature increase rate, but if the temperature increase rate becomes too fast, it takes time to reach the temperature reached and stabilize at that temperature. From the viewpoint of heat treatment apparatus design, the temperature rise is preferably 0.5 ° C./second or more and 10 ° C./second or less. More preferably, it is 1 ° C./second or more and 7 ° C./second or less, and further preferably 1 ° C./second or more and 6 ° C./second or less. The holding time is not particularly affected as long as it reaches the ultimate temperature, but a holding time of 1 minute or longer is preferable in order to obtain a stable heat treatment state.

The nanocomposite magnet obtained after the heat treatment has a structure in which α-Fe is mainly present at the grain boundary triple point of the R ₂ Fe ₁₄ B type crystal phase having an average crystal grain size of 5 nm or more and 100 nm or less. The average crystal grain size of the phase is 1 nm or more and 50 nm or less. The existence ratio of α-Fe is considered to be 5% by volume or more of the whole, and the residual magnetic flux density of the whole magnet is improved.

  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. Note that the ribbon of the rapidly solidified alloy may be roughly cut or pulverized before the heat treatment. The aspect ratio of the obtained magnet powder shape is 0.3 or more and 1 or less. When producing a bonded magnet, the iron-based rare earth alloy magnetic powder is mixed with a resin such as epoxy, polyamide, polyphenylene sulfide (PPS), or liquid crystal polymer, 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 of this embodiment 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. More preferably, the particle size distribution has two peaks and the average particle size is 50 μm or more and 200 μm or less.

  In addition, by performing surface treatment such as coupling treatment, chemical conversion treatment, and plating on the surface of the powder, the formability at the time of forming the bonded magnet and the corrosion resistance and heat resistance of the obtained bonded magnet can be improved regardless of the forming method. Further, when the surface of the bonded magnet after molding is subjected to surface treatment such as resin coating, chemical conversion treatment or plating, the corrosion resistance and heat resistance of the bonded magnet can be improved in the same manner as the powder surface treatment.

(Magnetic powder magnetic properties)
A rapidly solidified alloy slab was produced from a raw material alloy having the composition shown in the following tables using a strip casting apparatus. The rapid cooling conditions such as the roll peripheral speed were adjusted so that the average thickness of the rapidly solidified alloy slab was 60 μm or more and 300 μm or less. The obtained slab was roughly pulverized to a size of 850 μm or less, and then heat-treated at a predetermined temperature shown in the following table. The heat treatment temperature was set to a temperature region where Nd ₂ Fe ₁₄ B was not coarsened and a fine metal structure was obtained in each alloy composition. The heat treatment was performed in an argon atmosphere. By this heat treatment, the amorphous portion of the rapidly solidified alloy slab powder (coarse pulverized powder) was crystallized to obtain magnetic powder having a magnet structure. Magnetic properties (residual magnetic flux density B _r , intrinsic coercive force H _cJ , maximum energy product (BH) _max ) were measured for the magnetic powder of each sample after the heat treatment.

Table 1 shows the alloy compositions and heat treatment temperatures of Examples 1 to 3 and Comparative Example 1, and Table 2 shows the magnetic properties after heat treatment of Examples 1 to 3 and Comparative Example 1. Comparing the magnetic characteristics of Examples 1 to 3 and Comparative Example 1, in the comparative example in which the Nd composition ratio is relatively low, the residual magnetic flux density B _r , the intrinsic coercive force H _cJ , and the maximum energy product are compared with those in the example. (BH) Both _max are low.

Table 3 below shows the alloy compositions and heat treatment temperatures of Example 4 and Comparative Examples 2 to 4, and Table 4 shows the magnetic properties after heat treatment of Example 4 and Comparative Examples 2 to 4. In Comparative Examples 2 and 3 to which Ti is not added, the intrinsic coercive force H _cJ and the maximum energy product (BH) _max are lower than those in Example 4. The reason why the magnetic characteristics of the comparative example deteriorated relatively was that Nd ₂ Fe ₁₄ B was not coarsened and heat treatment was performed in a temperature region where a fine metal structure was obtained, but Ti was not added. This is probably because α-Fe was not sufficiently precipitated in the region.

Table 5 below shows the alloy compositions and heat treatment temperatures of Examples 4-7 and Comparative Examples 5-7, and Table 6 shows the magnetic properties after heat treatment of Examples 4-7 and Comparative Examples 5-7. Yes. In Comparative Examples 5 and 6 to which Zr is not added, the intrinsic coercive force H _cJ and the maximum energy product (BH) _max are significantly lower than those in Examples 4 to 6. In Comparative Example 7 in which Ti is larger than Zr, the decrease in the intrinsic coercive force H _cJ is suppressed as compared with Comparative Examples 5 and _6. However, compared with Examples 4 to 6, the intrinsic coercive force H _cJ and the maximum energy product are _reduced. (BH) _max is an insufficient level.

  Table 7 below shows the alloy compositions and heat treatment temperatures of Example 8 and Comparative Example 8, and Table 8 shows the magnetic properties after heat treatment of Example 8 and Comparative Example 8. In Comparative Example 8 with a small amount of B, the magnet characteristics are deteriorated as compared with Example 8.

  Table 9 below shows the alloy compositions and heat treatment temperatures of Examples 9-12, and Table 10 shows the magnetic properties after heat treatment of Examples 9-12.

As is apparent from the measurement results of the magnetic characteristics regarding the above-described embodiments, according to the present invention, the residual magnetic flux density B _r ≧ 0.75 T, the intrinsic coercive force H _cJ ≧ 640 kA / m, and the maximum energy product (BH) _max An iron-based rare earth nanocomposite magnet having permanent magnet characteristics of ≧ 100 kJ / m ³ is obtained.

In addition, about the coarsely pulverized powders of Examples 1 to 12, the crystal phase was examined by powder XRD. As a result, it was a metal structure composed of an Nd ₂ Fe ₁₄ B phase and an α-Fe phase. Further, when the metal structure was examined using a transmission electron microscope, the average crystal grain size of the α-Fe phase was 1 nm to 50 nm, and the average crystal grain size of the Nd ₂ Fe ₁₄ B phase was 5 nm to 100 nm. . Furthermore, when each powder particle of the finely pulverized powder having an average particle diameter of 10 μm or more and 300 μm or less was observed with a microscope, the aspect ratio of the powder particles was in the range of 0.3 or more and 1 or less.

(Magnetization characteristics)
A compression-molded bonded magnet was produced and the magnetization characteristics were evaluated. First, magnetic powders were produced in the same manner as in Example 1 for the alloy compositions shown in Table 11. Table 12 shows the magnetic powder magnetic properties after the heat treatment.

  After finely pulverizing the magnetic powder to a size of 250 μm or less, the finely pulverized magnetic powder was kneaded with 2 mass% of an epoxy resin and compounded. Further, 0.05 to 0.1 mass% of calcium stearate as a lubricant was added and mixed uniformly. This compound was put into a φ10 mm cavity and compression molded at a molding pressure of about 980 MPa. The compound amount was adjusted so that the molded body height after molding was about 7 mm. In order to solidify the resin component of the molded body, a curing treatment was performed in a reduced pressure atmosphere at 150 ° C. for 1 hour to obtain a compression molded bond magnet.

  A predetermined magnetizing magnetic field was applied to the obtained compression-molded bonded magnet with a pulse magnetizer, and magnetized. Thereafter, the magnetization rate was determined by measuring the flux. The magnetization rate was obtained by the following formula.

  FIG. 2 is a graph showing the magnetization characteristics of Example 5 and Comparative Example 9. The horizontal axis of the graph is the magnetization magnetic field, and the vertical axis is the magnetization rate. FIG. 3 is an enlarged graph of a part of FIG.

  According to the bonded magnet of this example, the magnetization magnetic field required for 90% magnetization when the magnetization field of 4.8 MA / m is 100% is 1.25 MA / m, which is 1.5 MA / m or less. It was confirmed that it was lowered. The magnetization field of the comparative example was 1.57 MA / m.

  According to the present invention, there is provided an iron-based rare earth nanocomposite magnet with improved magnetizing characteristics and wear properties such as a mold while maintaining high coercive force and excellent compound fluidity, which are the characteristics of a Ti-containing nanocomposite magnet. Provided and suitably used for electronic devices that require high magnetization characteristics such as small motors for small-diameter, multi-pole information appliances.

1 Melting furnace 2 Molten alloy 3 Cooling roll 4 Chute (guide means)

Claims (8)

Transition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T , including a Fe or one or more elements selected from the group consisting of Co and Ni and Fe Metal element, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb One or more metal elements selected from the group consisting of:
The composition ratios x, y, z, t, m, and p are respectively
7 ≦ x ≦ 9 atomic%,
6.5 ≦ y ≦ 9 atomic%,
2 ≦ z ≦ 5 atomic%,
0.5 ≦ t ≦ 3 atomic%,
4.5 ≦ z + t ≦ 7 atomic%,
0.5 ≦ z / (z + t) ≦ 0.95,
0 ≦ m ≦ 5 atomic%,
0 ≦ p ≦ 0.5,
Satisfied ,
The average particle size has a particle size of 10 μm to 300 μm,
An iron-based rare earth nanocomposite magnet powder having a ratio of a minor axis size to a major axis size of a powder particle of 0.3 or more and 1 or less . An α-Fe phase having an average crystal grain size of 1 nm to 50 nm as a soft magnetic phase;
The iron-based rare earth-based nanocomposite magnet powder according to claim 1, which has a metal structure including an R ₂ Fe ₁₄ B phase having an average crystal grain size of 5 nm to 100 nm as a hard magnetic phase. Residual magnetic flux density B _r ≧ 0.75 T,
Maximum energy product (BH) _max ≧ 100 kJ / m ³ ,
The iron-based rare earth-based nanocomposite magnet powder according to claim 1 or 2, which has a permanent magnet characteristic with an intrinsic coercive force H _cJ ≧ 640 kA / m. The iron-based rare earth nanocomposite magnet powder according to any one of claims 1 to 3 ,
A bonded magnet including a binder that binds the iron-based rare earth nanocomposite magnet powder. The bonded magnet according to claim 4 , wherein the magnetized magnetic field required for 90% magnetization when compression-molded and the 4.8 MA / m magnetized magnetic field is 100% is 1.5 MA / m or less. Transition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T , including a Fe or one or more elements selected from the group consisting of Co and Ni and Fe Metal element, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb One or more metal elements selected from the group consisting of:
The composition ratios x, y, z, t, m, and p are respectively
7 ≦ x ≦ 9 atomic%,
6.5 ≦ y ≦ 9 atomic%,
2 ≦ z ≦ 5 atomic%,
0.5 ≦ t ≦ 3 atomic%,
4.5 ≦ z + t ≦ 7 atomic%,
0.5 ≦ z / (z + t) ≦ 0.95,
0 ≦ m ≦ 5 atomic%,
0 ≦ p ≦ 0.5,
Satisfied,
A rapidly solidified alloy for iron-based rare earth nanocomposite magnets having an average thickness of 60 μm to 300 μm. Transition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T , including a Fe or one or more elements selected from the group consisting of Co and Ni and Fe Metal element, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb One or more metal elements selected from the group consisting of:
The composition ratios x, y, z, t, m, and p are respectively
7 ≦ x ≦ 9 atomic%,
6.5 ≦ y ≦ 9 atomic%,
2 ≦ z ≦ 5 atomic%,
0.5 ≦ t ≦ 3 atomic%,
4.5 ≦ z + t ≦ 7 atomic%,
0.5 ≦ z / (z + t) ≦ 0.95,
0 ≦ m ≦ 5 atomic%,
0 ≦ p ≦ 0.5,
Forming a molten alloy that satisfies the following conditions:
Cooling the molten metal by a strip casting method , thereby forming a rapidly solidified alloy having a thickness of 60 μm or more and 300 μm or less;
Crystallization of the rapidly solidified alloy by heat treatment to produce an alloy having permanent magnet properties;
Of producing an iron-based rare earth nanocomposite magnet comprising Transition formula _{T 100-xyztm (B 1-} p + C p) x R y Zr z Ti t M m (T , including a Fe or one or more elements selected from the group consisting of Co and Ni and Fe Metal element, R is one or more rare earth elements, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb One or more metal elements selected from the group consisting of:
The composition ratios x, y, z, t, m, and p are respectively
7 ≦ x ≦ 9 atomic%,
6.5 ≦ y ≦ 9 atomic%,
2 ≦ z ≦ 5 atomic%,
0.5 ≦ t ≦ 3 atomic%,
4.5 ≦ z + t ≦ 7 atomic%,
0.5 ≦ z / (z + t) ≦ 0.95,
0 ≦ m ≦ 5 atomic%,
0 ≦ p ≦ 0.5,
Forming a molten alloy that satisfies the following conditions:
Cooling the molten metal by a strip casting method , thereby forming a rapidly solidified alloy having a thickness of 60 μm or more and 300 μm or less;
Crystallization of the rapidly solidified alloy by heat treatment to produce an alloy having permanent magnet properties;
Crushing the alloy to form a powder having an average particle size of 10 μm or more and 300 μm or less and a ratio of the minor axis direction size to the major axis size of the powder particles of 0.3 or more and 1 or less;
A method for producing an iron-based rare earth-based nanocomposite magnet powder.

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