IRON-BASED RARE EARTH BORON-BASED ISOTROPIC MAGNET ALLOY

Abstract

An iron-based rare earth boron-based isotropic magnet alloy, which has an alloy composition represented by T.sub.100-x-y-z(B.sub.1-nC.sub.n).sub.xRE.sub.yM.sub.z (where T is a transition metal element containing at least Fe, RE contains at least Nd, and M is one or more metal elements selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb), 4.2 atom %≤x≤5.6 atom %, 11.5 atom %≤y≤13.0 atom %, 0.0 atom %≤z≤5.0 atom %, and 0.0≤n≤0.5, and the iron-based rare earth boron-based isotropic magnet alloy has an average crystal grain size of 10 nm to less than 70 nm as a main phase.

Claims

1. An iron-based rare earth boron-based isotropic magnet alloy having an alloy composition represented by: T.sub.100-x-y-z(B.sub.1-nC.sub.n).sub.xRE.sub.yM.sub.z, wherein T is a transition metal element containing at least Fe; RE comprises at least Nd; M is one or more metal elements selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb; 4.2 atom %≤x≤5.5 atom %; 11.5 atom %≤y≤13.0 atom %; 0.0 atom %≤z≤5.0 atom %; 0.0≤n≤0.5, and wherein the iron-based rare earth boron-based isotropic magnet alloy has an average crystal grain size of 10 nm to less than 70 nm as a main phase.

2. The iron-based rare earth boron-based isotropic magnet alloy according to claim 1, wherein the iron-based rare earth boron-based isotropic magnet alloy has a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1400 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more.

3. The iron-based rare earth boron-based isotropic magnet alloy according to claim 1, wherein the RE further comprises Pr.

4. The iron-based rare earth boron-based isotropic magnet alloy according to claim 1, wherein the T further comprises at least one of Co and Ni.

5. The iron-based rare earth boron-based isotropic magnet alloy according to claim 1, further comprising a grain boundary phase surrounding the main phase.

6. The iron-based rare earth boron-based isotropic magnet alloy according to claim 5, wherein the grain boundary phase comprises RE and Fe as main components thereof.

7. The iron-based rare earth boron-based isotropic magnet alloy according to claim 5, wherein the grain boundary phase is a ferromagnetic phase.

8. The iron-based rare earth boron-based isotropic magnet alloy according to claim 5, wherein a width of the grain boundary phase is 1 nm to less than 10 nm.

9. An iron-based rare earth boron-based isotropic magnet alloy having an alloy composition represented by: T.sub.100-x-y-z(B.sub.1-nC.sub.n).sub.xRE.sub.yM.sub.z, wherein T is a transition metal element containing at least Fe; RE comprises at least Nd; M is one or more metal elements selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb; 4.2 atom %≤x≤5.6 atom %; 11.5 atom %≤y≤13.0 atom %; 0.0 atom %≤z≤5.0 atom %; 0.0≤n≤0.5, wherein the iron-based rare earth boron-based isotropic magnet alloy has a metal structure comprising an RE.sub.2Fe.sub.14B-type tetragonal compound with an average crystal grain size of 10 nm to less than 70 nm as a main phase, and has a B-containing concentration lower than a stoichiometric composition of the RE.sub.2Fe.sub.14B-type tetragonal compound; and a grain boundary phase surrounding the main phase.

10. The iron-based rare earth boron-based isotropic magnet alloy according to claim 9, wherein the grain boundary phase surrounding the main phase comprises RE and Fe as main components thereof.

11. The iron-based rare earth boron-based isotropic magnet alloy according to claim 10, wherein the grain boundary phase is a ferromagnetic phase.

12. The iron-based rare earth boron-based isotropic magnet alloy according to claim 10, wherein a width of the grain boundary phase is 1 nm to less than 10 nm.

13. The iron-based rare earth boron-based isotropic magnet alloy according to claim 10, wherein a ratio of the main phase is 70 vol % to less than 99 vol %, and a ratio of the grain boundary phase is 1 vol % to less than 30 vol %.

14. The iron-based rare earth boron-based isotropic magnet alloy according to claim 9, wherein the iron-based rare earth boron-based isotropic magnet alloy has a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1400 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more.

15. The iron-based rare earth boron-based isotropic magnet alloy according to claim 9, wherein the RE further comprises Pr.

16. A method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy, the method comprising: preparing a molten alloy having a composition represented by T.sub.100-x-y-z(B.sub.1-nC.sub.n).sub.xRE.sub.yM.sub.z, wherein T is a transition metal element containing at least Fe, RE is at least one rare earth element substantially not containing La and Ce, M is one or more metal elements selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb, 4.2 atom %≤x≤5.6 atom %, 11.5 atom %≤y≤13.0 atom %, 0.0 atom %≤z≤5.0 atom %, and 0.0≤n≤0.5; and injecting the molten alloy onto a surface of a rotating roll comprising Cu, Mo, W or an alloy containing at least one of these metals as a main component, at an average metal tapping rate of 200 g/min to less than 2000 g/min per hole of an orifice arranged at a tip of a nozzle to prepare a rapidly solidified alloy having 1 vol % or more of either a crystal phase or an amorphous phase containing an RE.sub.2Fe.sub.14B phase.

17. The method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy according to claim 16, further comprising: performing flash annealing on the rapidly solidified alloy by making a temperature reach a constant temperature range of a crystallization temperature or higher and 850° C. or less at a temperature rising rate of 10° C./sec to less than 200° C./sec; and then quenching the rapidly solidified alloy after a lapse of 0.1 sec to less than 7 min so as to form a metal structure finer than a single magnetic domain critical diameter of an RE.sub.2Fe.sub.14B-type tetragonal compound, having an average crystal grain size of 10 nm to less than 70 nm as a main phase, and having a B-containing concentration lower than stoichiometric composition of the RE.sub.2Fe.sub.14B-type tetragonal compound, and a grain boundary phase surrounding the main phase with a width of 1 nm to less than 10 nm comprising RE and Fe as main components thereof.

18. The method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy according to claim 16, further comprising preparing an iron-based rare earth boron-based isotropic magnet alloy powder by pulverizing the rapidly solidified alloy.

19. A method for manufacturing a resin-bonded permanent magnet, comprising: preparing the iron-based rare earth boron-based isotropic magnet alloy powder according to claim 18; adding a thermosetting resin to the iron-based rare earth boron-based isotropic magnet alloy powder to form a mixture; filling a molding die with the mixture; forming a compression molded body by compression molding; and performing a heat treatment at a temperature equal to or higher than a polymerization temperature of the thermosetting resin.

20. A method for manufacturing a resin-bonded permanent magnet, comprising: preparing the iron-based rare earth boron-based isotropic magnet alloy powder according to claim 18; adding a thermoplastic resin to the iron-based rare earth boron-based isotropic magnet alloy powder to prepare an injection molding compound; and performing injection molding using the injection molding compound.

Description

BRIEF DESCRIPTION OF THE DRAWINGS

[0034] FIG. 1 is a cross-sectional view schematically showing an example of an iron-based rare earth boron-based isotropic magnet alloy of the present invention.

[0035] FIG. 2A is an apparatus configuration diagram of a heat treatment furnace for realizing flash annealing, and FIG. 2B is a diagram showing a state of a rapidly solidified alloy moving in a furnace core tube.

[0036] FIG. 3 is a conceptual diagram of a thermal history by flash annealing performed in the present invention.

[0037] FIG. 4 is a bright field image and elemental mapping obtained by observing an iron-based rare earth boron-based isotropic magnet alloy obtained in Example 13 with a transmission electron microscope.

[0038] FIG. 5 is a bright field image and elemental mapping obtained by observing an iron-based rare earth boron-based isotropic magnet alloy obtained in Comparative Example 38 with a transmission electron microscope.

[0039] FIG. 6 is a powder X-ray diffraction profile of a rapidly solidified alloy obtained in Example 13.

[0040] FIG. 7 is a powder X-ray diffraction profile of a rapidly solidified alloy after flash annealing (crystallization heat treatment) obtained in Example 13.

[0041] FIG. 8 is a powder X-ray diffraction profile of a rapidly solidified alloy after flash annealing (crystallization heat treatment) obtained in Comparative Example 38.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0042] Hereinafter, an iron-based rare earth boron-based isotropic magnet alloy of the present invention, a method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy of the present invention, and a method for manufacturing a resin-bonded permanent magnet of the present invention will be described. Note that the present invention is not limited to the following configuration, and may be appropriately modified without departing from the gist of the present invention. The present invention also includes a combination of a plurality of preferred configurations described below.

[0043] An iron-based rare earth boron-based isotropic magnet alloy of the present invention has, in a first aspect, an alloy composition represented by T.sub.100-x-y-z(B.sub.1-nC.sub.n).sub.xRE.sub.yM.sub.z (wherein T is at least one element selected from the group consisting of Fe, Co, and Ni, and is a transition metal element necessarily containing Fe; RE is at least one rare earth element necessarily containing at least Nd among Nd and Pr; and M is one or more metal elements selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb); 4.2 atom %≤x≤5.6 atom %; 11.5 atom %≤y≤13.0 atom %; 0.0 atom %≤z≤5.0 atom %; and 0.0≤n≤0.5. The iron-based rare earth boron-based isotropic magnet alloy has a metal structure finer than the single magnetic domain critical diameter of an RE.sub.2Fe.sub.14B-type tetragonal compound with an average crystal grain size of 10 nm to less than 70 nm as a main phase, while having a B-containing concentration lower than a stoichiometric composition of the RE.sub.2Fe.sub.14B-type tetragonal compound.

[0044] The iron-based rare earth boron-based isotropic magnet alloy of the present invention has, in a second aspect, an alloy composition represented by T.sub.100-x-y-z(B.sub.1-nC.sub.n).sub.xRE.sub.yM.sub.z (wherein T is at least one element selected from the group consisting of Fe, Co, and Ni, and is a transition metal element necessarily containing Fe; RE is at least one rare earth element necessarily containing at least Nd among Nd and Pr; and M is one or more metal elements selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb); 4.2 atom %≤x≤5.6 atom %; 11.5 atom %≤y≤13.0 atom %; 0.0 atom %≤z≤5.0 atom %; and 0.0≤n≤0.5. The iron-based rare earth boron-based isotropic magnet alloy has a metal structure having the RE.sub.2Fe.sub.14B-type tetragonal compound with an average crystal grain size of 10 nm to less than 70 nm as a main phase, in which a grain boundary phase surrounding the main phase is present, while having a B-containing concentration lower than a stoichiometric composition of the RE.sub.2Fe.sub.14B-type tetragonal compound. An example of the iron-based rare earth boron-based isotropic magnet alloy of the present invention as above is shown in FIG. 1, where the main phase 21 is surrounded by the grain boundary phase 22.

[0045] It is preferable that the iron-based rare earth boron-based isotropic magnet alloy of the present invention, in the second aspect, has a metal structure finer than the single magnetic domain critical diameter of the RE.sub.2Fe.sub.14B-type tetragonal compound, in which the grain boundary phase surrounding the main phase composed of the RE.sub.2Fe.sub.14B-type tetragonal compound contains RE and Fe as main components.

[0046] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, in the second aspect, the grain boundary phase containing RE and Fe as main components and surrounding the main phase composed of the RE.sub.2Fe.sub.14B-type tetragonal compound is preferably a ferromagnetic phase.

[0047] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, in the second aspect, the width of the grain boundary phase containing RE and Fe as main components and surrounding the main phase composed of the RE.sub.2Fe.sub.14B-type tetragonal compound is preferably 1 nm to less than 10 nm.

[0048] The iron-based rare earth boron-based isotropic magnet alloy of the present invention has a low boron content concentration, and the boron (B) content concentration is in a range of 4.2 atom % to 5.6 atom %. Furthermore, in the iron-based rare earth boron-based isotropic magnet alloy of the present invention, the rare earth element (RE) and iron (Fe) are brought into a surplus state in the same alloy structure, so that a grain boundary phase containing surplus RE and Fe which are not necessary for generation of the RE.sub.2Fe.sub.14B phase as the main phase is formed. As a result, the iron-based rare earth boron-based isotropic magnet alloy of the present invention has a unique fine metal structure in which a grain boundary phase with a width of 1 nm to less than 10 nm containing RE and Fe as main components and surrounding the RE.sub.2Fe.sub.14B phase with an average crystal grain size of 10 nm to less than 70 nm is present.

[0049] The present inventors have found that by realizing the above unique uniform and fine metal structure, the RE.sub.2Fe.sub.14B phase as the main phase and the grain boundary phase having RE and Fe as main components, which is uniformly present around the main phase, are bound by a strong exchange interaction in addition to a magnetostatic interaction, and behave as if they are an integrated hard magnetic phase with the grain boundary phase (for example, α-Fe phase) having saturation magnetization equal to or higher than that of the main phase, thereby obtaining a high residual magnetic flux density Br and a high maximum energy product (BH) max by improving squareness of demagnetization curve without impairing the intrinsic coercive force HcJ of the RE.sub.2Fe.sub.14B phase. In particular, it is considered that having the grain boundary phase as described above contributes to developing a high intrinsic coercive force HcJ, and it is considered that having the small average crystal grain size as described above contributes to developing a high residual magnetic flux density Br and a high coercive force HcJ.

[0050] When the boron content concentration is less than 4.2 atom %, the generation of the RE.sub.2Fe.sub.14B phase as the main phase is inhibited, so that both the intrinsic coercive force HcJ and the residual magnetic flux density Br significantly decrease. In addition, when the boron content concentration exceeds 5.6 atom %, a metal structure in which an RE.sub.2Fe.sub.14B single phase is present or a nonmagnetic B-rich phase is present around the RE.sub.2Fe.sub.14B phase is obtained, and thus, although a high intrinsic coercive force HcJ can be maintained, the residual magnetic flux density Br and the maximum energy product (BH) max are not increased, and sufficient magnetic properties, for example, magnetic properties of a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1400 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more are not be obtained.

[0051] On the other hand, when the boron content concentration is set to 4.2 atom % to 5.6 atom %, the grain boundary phase containing RE and Fe as main components is uniformly generated without impairing the generation of the RE.sub.2Fe.sub.14B phase as the main phase, and thus the above magnetic properties are considered to be obtained.

[0052] Patent Literature 2, Patent Literature 3, Patent Literature 4, Patent Literature 5, and Patent Literature 6 all disclose a microcrystalline isotropic permanent magnet material in which an RE.sub.2Fe.sub.14B-type tetragonal compound bears intrinsic coercive force HcJ. However, the magnitude of the intrinsic coercive force HcJ mainly depends on the volume ratio of the RE.sub.2Fe.sub.14B-type tetragonal compound, and the intrinsic coercive force HcJ increases when the volume ratio of the RE.sub.2Fe.sub.14B phase is high, and the intrinsic coercive force HcJ decreases when the volume ratio of the RE.sub.2Fe.sub.14B phase is low.

[0053] On the other hand, in the anisotropic RE.sub.2Fe.sub.14B sintered magnet described in Patent Literature 1, heavy rare earth elements such as Dy and Tb are included in the RE.sub.2Fe.sub.14B-type tetragonal compound as the main phase, and the anisotropic magnetic field of the RE.sub.2Fe.sub.14B-type tetragonal compound is increased, thereby realizing improvement of the intrinsic coercive force HcJ. Although both of the fine isotropic permanent magnet material and the anisotropic sintered magnet have the RE.sub.2Fe.sub.14B-type tetragonal compound as the main phase, the main phase size of the anisotropic sintered magnet is about 1 μm to 10 μm, and is equal to or more than the single magnetic domain critical diameter of the RE.sub.2Fe.sub.14B-type tetragonal compound. Therefore, although the anisotropic sintered magnet is in a multi-magnetic domain state before magnetization, magnetic moments are aligned in the magnetization direction (C-axis direction) by magnetization, and the permanent magnet properties are exhibited by bringing the anisotropic sintered magnet into a single magnetic domain state. Thus, the intrinsic coercive force HcJ of the anisotropic sintered magnet represents an ability to maintain a state in which the magnetic moments are aligned in the same direction. Therefore, the intrinsic coercive force HcJ is improved by increasing the anisotropic magnetic field of the RE.sub.2Fe.sub.14B-type tetragonal compound.

[0054] In the iron-based rare earth boron-based isotropic magnet alloy having a low boron content concentration of the present invention, by realizing a unique metal structure having a grain boundary phase containing RE and Fe as main components, when a heavy rare earth element such as Dy is added to the alloy composition, the anisotropic magnetic field of not only the RE.sub.2Fe.sub.14B-type tetragonal compound as the main phase but also the grain boundary phase is improved. Therefore, it has been found that demagnetization of magnetic moment of the main phase with a single magnetic domain crystal grain size or less can be suppressed by the grain boundary phase, and improvement of the intrinsic coercive force HcJ by addition of the heavy rare earth element, which was not effective in the conventional fine crystal type isotropic RE.sub.2Fe.sub.14B permanent magnet material, can be realized. Accordingly, according to the iron-based rare earth boron-based isotropic magnet alloy of the present invention, a high-performance isotropic RE.sub.2Fe.sub.14B permanent magnet that has a high intrinsic coercive force HcJ and has not been conventionally obtained is obtain without causing a significant decrease in residual magnetic flux density Br.

[0055] In addition, it has been found that in the iron-based rare earth boron-based isotropic magnet alloy having a low boron content concentration of the present invention, improvement of the intrinsic coercive force HcJ is realized without causing a decrease in the residual magnetic flux density Br by substituting a part of the boron (B) with carbon (C), and further, the effect of improving the intrinsic coercive force HcJ can be increased by combining the substitution with the carbon (C) and the addition of a heavy rare earth element.

[0056] [Alloy Composition]

[0057] The alloy composition of the iron-based rare earth boron-based isotropic magnet alloy of the present invention is represented by formula T.sub.100-x-y-z(B.sub.1-n C.sub.n).sub.xRE.sub.yM.sub.z (wherein T is at least one element selected from the group consisting of Fe, Co, and Ni, and is a transition metal element necessarily containing Fe; RE is at least one rare earth element necessarily containing at least Nd among Nd and Pr; and M is one or more metal element selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb); 4.2 atom %≤x≤5.6 atom %; 11.5 atom %≤y≤13.0 atom %; 0.0 atom %≤z≤5.0 atom %; and 0.0≤n≤0.5. The composition of the entire magnet alloy according to the present invention is analyzed by ICP mass spectrometry. In addition, a combustion-infrared absorption method may be used in combination as necessary.

[0058] Transition metal element T containing Fe as an essential element occupies the remainder of the content of the above-described elements. Even if a part of Fe is substituted with one or two of Co and Ni which are ferromagnetic elements like Fe, desired hard magnetic properties can be obtained. However, when the amount of substitution for Fe exceeds 30%, the magnetic flux density is significantly reduced, and therefore the amount of substitution is preferably in the range of 0% to 30%. It is to be noted that the addition of Co not only contributes to improvement of magnetization, but also has an effect of lowering the viscosity of the molten metal to stabilize the metal tapping rate from a nozzle at the time of quenching the molten metal. Therefore, the amount of substitution by Co is more preferably 0.5% to 30%, and from the viewpoint of cost effectiveness, the amount of substitution by Co is still more preferably 0.5% to 10%.

[0059] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, when the composition ratio x of B+C is less than 4.2 atom %, the amount of B+C required for producing an RE.sub.2Fe.sub.14B-type tetragonal compound cannot be secured, and the magnetic properties are deteriorated and amorphous forming ability is greatly deteriorated, so that an α-Fe phase is precipitated during molten metal rapid solidification, and as a result, the squareness of the demagnetization curve is impaired. In addition, when the composition ratio x of B+C exceeds 5.6 atom %, a grain boundary phase containing RE and Fe as main components is not generated, and there is a possibility that the above-described magnetic properties cannot be secured. Accordingly, the composition ratio x is limited to a range of 4.2 atom % to 5.6 atom %. The composition ratio x is preferably 4.2 atom % to 5.2 atom %, and more preferably 4.4 atom % to 5.0 atom %.

[0060] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, by substituting a part of B with C, a melting point of the molten alloy is lowered, and the wear amount of a refractory used at the time of rapid solidification is reduced, so that the process cost related to rapid solidification can be reduced, and the effect of improving the intrinsic coercive force HcJ is obtained. However, it is not preferable that the substitution rate of C for B exceeds 50% since the amorphous forming ability is greatly deteriorated. Accordingly, the substitution rate of C for B is limited to a range of 0% to 50%, that is, 0.0≤n≤0.5. From the viewpoint of the effect of improving the intrinsic coercive force HcJ, the substitution rate of C for B is preferably 2% to 30%, and more preferably 3% to 15%.

[0061] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, when the composition ratio y of at least one rare earth element RE necessarily containing at least Nd among Nd and Pr is less than 11.5 atom %, a grain boundary phase containing RE and Fe as main components is not generated, and there is a possibility that the above-described magnetic properties cannot be secured. In addition, when the composition ratio y exceeds 13.0 atom %, the magnetization decreases. Accordingly, the composition ratio y is limited to a range of 11.5 atom % to 13.0 atom %. Moreover, the composition ratio y is preferably 11.76 atom % to 13.0 atom %, which is the stoichiometric composition of the RE.sub.2Fe.sub.14B-type tetragonal compound, from the viewpoint of ensuring stability of the intrinsic coercive force HcJ, and more preferably 11.76 atom % to 12.5 atom % from the viewpoint of ensuring a high residual magnetic flux density Br.

[0062] Further, the rare earth RE may be RE.sub.y=(Nd.sub.1-lPr.sub.l).sub.y in order to obtain a higher intrinsic coercive force HcJ, and at that time, l is limited to 0.05 to 0.7. It is to be noted that if substitution ratio l of Pr for Nd is too low, the effect of improving HcJ is small, and if l is too high, an absolute value of temperature coefficient β related to coercive force of the magnet alloy becomes small, so that there is a concern that the heat resistance may deteriorate. Therefore, l is preferably 0.15 to 0.6, and further preferably 0.2 to 0.5.

[0063] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, one or more metal elements M selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb may be added. By the addition of the metal element M, effects such as improvement of the amorphous forming ability, improvement of the intrinsic coercive force HcJ by uniform refinement of a metal structure after crystallization heat treatment, improvement in the squareness of the demagnetization curve, and the like are obtained, and the magnetic properties are improved. However, when the composition ratio z of these metal elements M exceeds 5.0 atom %, the magnetization decreases, and thus the composition ratio z is limited to a range of 0.0 atom % to 5.0 atom %. In addition, the composition ratio z is preferably 0.0 atom % to 4.0 atom %, and more preferably 0.0 atom % to 3.0 atom %.

[0064] [Metal Structure]

[0065] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, when the average crystal grain size of the RE.sub.2Fe.sub.14B-type tetragonal compound as the main phase is less than 10 nm, the intrinsic coercive force HcJ decreases, and when the average crystal grain size is 70 nm or more, the squareness of the demagnetization curve decreases due to a decrease in exchange interaction acting between crystal grains. Therefore, for example, in order to realize magnetic properties of a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1400 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more, an average crystal grain size of the RE.sub.2Fe.sub.14B-type tetragonal compound is limited to a range of 10 nm to less than nm. The average crystal grain size of the RE.sub.2Fe.sub.14B-type tetragonal compound is preferably 15 nm to 60 nm, and more preferably 15 nm to 50 nm.

[0066] The average crystal grain size of the RE.sub.2Fe.sub.14B-type tetragonal compound means the average value of the equivalent circle diameters of particles present in the field of view when the particle size of each particle is measured at 3 or more points by a line segment method using a transmission electron microscope (TEM).

[0067] When the width of the grain boundary phase containing RE and Fe as main components and surrounding the main phase composed of the RE.sub.2Fe.sub.14B-type tetragonal compound is less than 1 nm, the bonding force acting between the main phase grains increases, leading to a decrease in the intrinsic coercive force HcJ. In addition, when the width of the grain boundary phase is 10 nm or more, conversely, interparticle bonding is weakened, and the square shape of the demagnetization curve decreases. Therefore, the width of the grain boundary phase is preferably 1 nm to less than 10 nm, more preferably 2 nm to 8 nm, and still more preferably 2 nm to 5 nm. The width of the grain boundary phase was determined by performing image analysis on a bright field image taken using a scanning transmission electron microscope under the conditions of an acceleration voltage of 200 kV and an observation magnification of 900,000 times.

[0068] In the iron-based rare earth boron-based isotropic magnet alloy of the present invention, in the second aspect, in the constituent ratio of the grain boundary phase containing RE and Fe as main components and surrounding the main phase composed of the RE.sub.2Fe.sub.14B-type tetragonal compound, it is preferable that the ratio of the main phase is 70 vol % to less than 99 vol %, and the ratio of the grain boundary phase is 1 vol % to less than 30 vol %. This makes it easy to realize magnetic properties of, for example, a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1400 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more. The ratio of the main phase is preferably 80 vol % to less than 99 vol %, and more preferably 90 vol % to less than 98 vol %. The constituent ratio of the main phase and the grain boundary phase was determined by performing image analysis on a bright field image taken using a scanning transmission electron microscope under the conditions of an acceleration voltage of 200 kV and an observation magnification of 900,000 times.

[0069] [Magnetic Properties]

[0070] As will be described later, the iron-based rare earth boron-based isotropic magnet alloy of the present invention can exhibit magnetic properties of, for example, a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1200 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more. However, in a case of a magnetic circuit configuration in which a reverse magnetic field is likely to be applied to a permanent magnet such as a surface magnet rotor (SPM rotor), when using the iron-based rare earth boron-based isotropic magnet alloy in various rotating machines optimum for electrical equipment and white goods of about 1 horsepower (750 W) or less, the intrinsic coercive force HcJ is preferably 800 kA/m or more and more preferably 950 kA/m or more. When the intrinsic coercive force HcJ is 1400 kA/m or more, magnetizability is significantly reduced, and thus the intrinsic coercive force HcJ is preferably 1300 kA/m or less, and more preferably 1250 kA/m or less. In addition, regarding the residual magnetic flux density Br, in a case where an interior permanent magnet rotor (IPM rotor) or the like is adopted, it is possible to drive at a higher operating point (permeance) than the SPM type. Therefore, although the residual magnetic flux density Br is preferably as high as possible, in consideration of the balance with the intrinsic coercive force HcJ, the residual magnetic flux density Br is preferably 0.87 T or more, and more preferably 0.9 T or more.

[0071] The reason why the residual magnetic flux density Br was set to 0.85 T or more as an example is that, in the case of applying to a DC brushless motor as an isotropic bonded magnet, an operating point (permeance Pc) of the magnet is about 3 to 10, and thus, in the residual magnetic flux density Br 0.85 T, an execution magnetic flux Bm which is equivalent to that of an anisotropic Nd—Fe—B sintered magnet with a maximum energy product (BH) max of 300 kJ/m.sup.3 or more can be obtained within the range of this Pc. The residual magnetic flux density Br is still more preferably 0.86 T or more.

[0072] In addition, the reason why the intrinsic coercive force HcJ was set to 700 kA/m or more as an example is that when the intrinsic coercive force HcJ is less than 700 kA/m, in the case of applying to a DC brushless motor as an isotropic bonded magnet, a heat resistance temperature of the motor of 100° C. cannot be secured, and there is a possibility that desired motor characteristics cannot be obtained due to thermal demagnetization. In addition, the reason why the intrinsic coercive force HcJ was set to less than 1400 kA/m is that magnetization is difficult when the intrinsic coercive force HcJ is 1400 kA/m or more, and multipolar magnetization for securing Pc of 3 to 10 is difficult.

[0073] Furthermore, the reason why the maximum energy product (BH) max was set to 120 kJ/m.sup.3 or more as an example is that when the maximum energy product (BH) max is less than 120 kJ/m.sup.3, the squareness ratio of the demagnetization curve (residual magnetization Jr/saturation magnetization Js) is 0.8 or less, and thus in the case of applying to a DC brushless motor as an isotropic bonded magnet, magnetic properties may be deteriorated due to a reverse magnetic field generated during motor operation, and there is a possibility that desired motor properties cannot be obtained.

[0074] A method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy of the present invention includes preparing a molten alloy having a composition represented by formula T.sub.100-x-y-z(B.sub.1-nC.sub.n).sub.xRE.sub.yM.sub.z (wherein T is at least one element selected from the group consisting of Fe, Co, and Ni, and is a transition metal element necessarily containing Fe; RE is at least one rare earth element substantially not containing La and Ce; and M is one or more metal elements selected from the group consisting of Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb); 4.2 atom %≤x≤5.6 atom %; 11.5 atom %≤y≤13.0 atom %; 0.0 atom %≤z≤5.0 atom %; and 0.0≤n≤0.5; and injecting the molten alloy onto a surface of a rotating roll containing Cu, Mo, W or an alloy containing at least one of these metals as a main component, at an average metal tapping rate of 200 g/min to less than 2000 g/min per hole of an orifice arranged at the tip of the nozzle to prepare a rapidly solidified alloy having 1 vol % or more of either a crystal phase or an amorphous phase containing an RE.sub.2Fe.sub.14B phase. Note that RE is at least one rare earth element substantially not containing La and Ce, but as an example, as described above, RE can be at least one rare earth element necessarily containing at least Nd among Nd and Pr. Details are as described above.

[0075] [Molten Metal Quenching]

[0076] In the method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy of the present invention, a raw material prepared so as to have a predetermined alloy composition is dissolved to form a molten alloy, and then the molten alloy is injected onto the surface of a rotating roll containing Cu, Mo, W or an alloy containing at least one of these metals as a main component, at an average metal tapping rate of 200 g/min to less than 2000 g/min per hole of an orifice arranged at the tip of the nozzle to prepare a rapidly solidified alloy having 1 vol % or more of either a crystal phase or an amorphous phase containing an RE.sub.2Fe.sub.14B phase, but when the average metal tapping rate is less than 200 g/min, productivity is poor, and when the average metal tapping rate is 2000 g/min or more, since a molten metal quenched alloy structure containing a coarse α-Fe phase is obtained, there is a possibility that the above-described magnetic properties cannot be obtained even if the crystallization heat treatment is performed. Accordingly, the average metal tapping rate per hole of the orifice arranged at the tip of the nozzle is limited to a range of 200 g/min to less than 2000 g/min. The average metal tapping rate is preferably 300 g/min to 1500 g/min, and more preferably 400 g/min to 1300 g/min.

[0077] The hole arranged at the tip of the nozzle and through which molten metal is tapped is not limited to a circular orifice, but may have any shape such as a square, a triangle, or an ellipse, and have a slit shape as long as the hole has a hole shape that can secure a predetermined molten metal tapping rate. In addition, the nozzle material is allowed as long as it is a refractory material that does not react with or hardly reacts with the molten alloy, but is preferably a ceramic material, SiC, C, or BN with less wear of a nozzle orifice due to the molten metal during tapping, more preferably BN, and still more preferably hard BN containing an additive.

[0078] When preparing the rapidly solidified alloy, the rapidly solidified atmosphere is preferably an oxygen-free or low-oxygen atmosphere since an increase in molten metal viscosity can be suppressed by preventing oxidation of the molten alloy, and a stable metal tapping rate can be maintained. In order to realize this atmosphere, it is necessary to perform rapid solidification after evacuating inside of a rapid solidification device to 20 Pa or less, preferably 10 Pa or less, and more preferably 1 Pa or less, then introducing an inert gas into the rapid solidification device, and setting the oxygen concentration in the rapid solidification device to 500 ppm or less, preferably 200 ppm or less, and more preferably 100 ppm or less. As the inert gas, a rare gas such as helium or argon or nitrogen can be used, but since nitrogen is relatively easily reacted with a rare earth element and iron, a rare gas such as helium or argon is preferable, and an argon gas is more preferable from the viewpoint of cost.

[0079] In the preparing a rapidly solidified alloy, the rotating roll that quenches the molten alloy contains Cu, Mo, W or an alloy containing at least one of these metals, as a main component, and preferably has a base material containing such a main component. This is because these base materials are excellent in thermal conductivity and durability. In addition, by plating Cr, Ni or a combination thereof on a surface of the base material of the rotating roll, heat resistance and hardness of the surface of the base material of the rotating roll can be enhanced, and melting and deterioration of the surface of the base material of the rotating roll during rapid solidification can be suppressed. The diameter of the rotating roll is, for example, Φ200 mm to Φ20,000 mm. When the rapid solidification time is a short time of 10 sec or less, it is not necessary to cool the rotating roll with water, but when the rapid solidification time exceeds 10 sec, it is preferable to flow cooling water into the rotating roll to suppress the temperature rise of the rotating roll base material. It is preferred that the water cooling capacity of the rotating roll is calculated according to the latent heat of solidification per unit time and the metal tapping rate, and optimally adjusted as appropriate.

[0080] [Flash Annealing]

[0081] The method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy of the present invention preferably further includes performing flash annealing on the rapidly solidified alloy by making the temperature reach a constant temperature range of a crystallization temperature or higher and 850° C. or less at a temperature rising rate of 10° C./sec to less than 200° C./sec, and then quenching after a lapse of 0.1 sec to less than 7 min, in which, by the performing flash annealing, the method forms a metal structure finer than the single magnetic domain critical diameter of an RE.sub.2Fe.sub.14B-type tetragonal compound, and has an average crystal grain size of 10 nm to less than 70 nm as a main phase, in which a grain boundary phase with a width of 1 nm to less than 10 nm containing RE and Fe as main components and surrounding the main phase is present, while having a B-containing concentration lower than a stoichiometric composition of the RE.sub.2Fe.sub.14B-type tetragonal compound.

[0082] When the temperature rising rate during flash annealing (crystallization heat treatment) is less than 10° C./sec, a fine metal structure cannot be obtained due to excessive grain growth, leading to decreases in the intrinsic coercive force HcJ and the residual magnetic flux density Br. When the temperature rising rate is 200° C./sec or more, the crystal grain growth cannot be made in time, and a metal structure finer than the single magnetic domain critical diameter of the RE.sub.2Fe.sub.14B-type tetragonal compound having the RE.sub.2Fe.sub.14B-type tetragonal compound with an average crystal grain size of 10 nm to less than 70 nm necessary for expression of permanent magnet as a main phase, in which a grain boundary phase with a width of 1 nm to less than 10 nm containing RE and Fe as main components and surrounding the main phase is present, is not obtained, leading to deterioration of the magnetic properties as in the case of less than 10° C./sec. Accordingly, the temperature rising rate is preferably 10° C./sec to less than 200° C./sec, more preferably 30° C./sec to 200° C./sec, and still more preferably 40° C./sec to 180° C./sec.

[0083] In the flash annealing (crystallization heat treatment) in the method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy of the present invention, in order to obtain good magnetic properties, it is preferable to immediately quench the alloy after reaching a crystallization heat treatment temperature (holding temperature) in a constant temperature range of a crystallization temperature or higher and 850° C. or less. More specifically, it is sufficient that the holding time after reaching the crystallization heat treatment temperature until quenching is substantially 0.1 sec or more, and it is not preferable that the holding time is 7 min or more since uniform and fine metal structures are impaired, leading to deterioration of various magnetic properties. Accordingly, the holding time is preferably 0.1 sec to less than 7 min, more preferably 0.1 sec to 2 min, and still more preferably 0.1 sec to 30 sec.

[0084] In the flash annealing (crystallization heat treatment) in the method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy of the present invention, it is preferable to cool the rapidly solidified alloy to 400° C. or less at a temperature drop rate of 2° C./sec to 200° C./sec. When the temperature drop rate is less than 2° C./sec, coarsening of the crystal structure proceeds, and when the temperature drop rate exceeds 200° C./sec, the alloy may be oxidized. Accordingly, the temperature drop rate is preferably 2° C./sec to 200° C./sec, more preferably 5° C./sec to 200° C./sec, and still more preferably 5° C./sec to 150° C./sec.

[0085] The atmosphere of the flash annealing (crystallization heat treatment) is preferably an inert gas atmosphere in order to prevent oxidation of the rapidly solidified alloy. As the inert gas, a rare gas such as helium or argon or nitrogen can be used, but since nitrogen is relatively easily reacted with a rare earth element and iron, a rare gas such as helium or argon is preferable, and an argon gas is more preferable from the viewpoint of cost.

[0086] [Pulverization and Molding]

[0087] The method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy of the present invention may further include preparing an iron-based rare earth boron-based isotropic magnet alloy powder by pulverizing the rapidly solidified alloy or the rapidly solidified alloy subjected to the flash annealing.

[0088] In the rapidly solidified alloy obtained through the above step, a thin band-shaped rapidly solidified alloy may be roughly cut or pulverized to, for example, 50 mm or less before flash annealing (crystallization heat treatment). Furthermore, by forming the magnet alloy of the present invention after flash annealing (crystallization heat treatment) into magnet alloy powder pulverized to an appropriate average powder particle diameter in a range of an average powder particle diameter of 20 μm to 200 μm, various resin-bonded permanent magnets (commonly called “plastic magnet” or “bonded magnet”) can be manufactured by known processes using the magnet alloy powder.

[0089] A method for manufacturing a resin-bonded permanent magnet of the present invention, in the first aspect, includes preparing an iron-based rare earth boron-based isotropic magnet alloy powder manufactured by the method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy; and adding a thermosetting resin to the iron-based rare earth boron-based isotropic magnet alloy powder, then filling a molding die with the mixture, forming a compression molded body by compression molding, and then performing heat treatment at a temperature equal to or higher than a polymerization temperature of the thermosetting resin.

[0090] A method for manufacturing a resin-bonded permanent magnet of the present invention, in the second aspect, includes preparing an iron-based rare earth boron-based isotropic magnet alloy powder manufactured by the method for manufacturing an iron-based rare earth boron-based isotropic magnet alloy; and adding a thermoplastic resin to the iron-based rare earth boron-based isotropic magnet alloy powder to prepare an injection molding compound, and then performing injection molding.

[0091] In the case of preparing the resin-bonded permanent magnet, iron-based rare earth-based nanocomposite magnet powder is mixed with epoxy, polyamide, polyphenylene sulfide (PPS), a liquid crystal polymer, acrylic, polyether, or the like, and molded into a desired shape. At this time, for example, hybrid magnet powder obtained by mixing permanent magnet powder such as SmFeN-based magnet powder or hard ferrite magnet powder may be used.

[0092] It is possible to manufacture various rotating machines and various magnetic sensors applicable to automobiles (also including electric vehicles and hybrid vehicles) and white goods as a brushless DC motor of about 1 horsepower (750 W) or less by using the above-described resin-bonded permanent magnet.

[0093] When the magnet alloy powder of the present invention is used for an injection-molded bonded magnet, the pulverization is preferably performed so that an average grain size is 100 μm or less, and the more preferable average crystal grain size of the powder is 20 μm to 100 μm. Also, when the magnet alloy powder is used for a compression-molded bonded magnet, the pulverization is preferably performed so that an average grain size is 200 μm or less, and the more preferable average crystal grain size of the powder is 50 μm to 150 μm. Still more preferably, the magnet alloy powder has two peaks in the particle size distribution, and the average crystal grain size is 80 μm to 130 μm.

[0094] By subjecting the surface of the magnet alloy powder of the present invention to surface treatment such as coupling treatment or chemical conversion treatment (including phosphoric acid treatment and glass coating treatment), it is possible to improve moldability at the time of molding the resin-bonded permanent magnet and corrosion resistance and heat resistance of the resin-bonded permanent magnet to be obtained regardless of the molding method. In addition, even when the surface of the resin-bonded permanent magnet after molding is subjected to surface treatment such as resin coating, chemical conversion treatment, or plating, it is possible to improve the corrosion resistance and heat resistance of the resin-bonded permanent magnet similarly to the surface treatment of the magnet alloy powder.

[0095] Incidentally, the method for manufacturing the iron-based rare earth boron-based isotropic magnet alloy of the present invention is not limited to the above-described methods, and other manufacturing methods can be adopted as long as the iron-based rare earth boron-based isotropic magnet alloy having the above-described composition, average crystal grain size, and the like can be manufactured. For example, when flash annealing is used, it is possible to form a fine metal structure having an RE.sub.2Fe.sub.14B-type tetragonal compound with an average crystal grain size of 10 nm to less than 70 nm as a main phase, but in order to form such a fine metal structure, the method is not limited to the flash annealing, and other methods can be adopted. For example, even in the case of adopting a normal annealing step instead of flash annealing, when the surface velocity of the rotating roll for quenching the molten alloy is adjusted to form a rapidly solidified alloy structure as a homogeneous fine metal structure composed of crystal grains about 5% to 20% smaller than those of the alloy structure from which optimal magnetic properties can be obtained, good magnetic properties can be obtained.

Examples

[0096] Hereinafter, examples of the present invention will be described. Note that the present invention is not limited only to these examples.

Examples

[0097] In order to obtain the alloy composition shown in Table 1, 100 g of a raw material in which additive elements such as Co, Al, Si, V, Cr, Ti, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb were blended in addition to main elements of Nd, Pr, Dy, B, C, and Fe with a purity of 99.5% or more was put into an alumina melting crucible, and then set in a work coil in a vacuum melting furnace. Then, the inside of the vacuum melting furnace was evacuated to 0.02 Pa or less, argon gas was then introduced to normal pressure, and a molten alloy was formed by high frequency induction heating. Thereafter, a molten alloy was cast into a water-cooled copper mold to prepare a mother alloy.

[0098] Subsequently, the obtained mother alloy was divided into an appropriate size, and then 40 g of the mother alloy was inserted into a transparent quartz nozzle having, at the bottom, an orifice with an appropriately different diameter (0.7 mm to 1.2 mm) so as to have an average metal tapping rate (in Table 1, simply shown as “metal tapping rate”) described in Table 1, and then the mother alloy was set in a work coil in a single roll quenching device. Then, the inside of the vacuum melting furnace was evacuated to 0.02 Pa or less, argon gas was then introduced until reaching the quenching atmospheric pressure shown in Table 1, the mother alloy was redissolved by high-frequency induction heating, and the molten alloy was tapped from a nozzle orifice at an injection pressure of 30 kPa onto the surface of the rotating roll rotating at the roll surface velocity (Vs) shown in Table 1 to prepare a rapidly solidified alloy. At this time, the distance between the tip of the nozzle and the surface of the rotating roll was set to 0.8 mm. Also, the main component of the rotating roll was copper. In addition, the obtained rapidly solidified alloy had 1 vol % or more of either a crystal phase or an amorphous phase containing an Nd.sub.2Fe.sub.14B phase.

[0099] As a representative example, FIG. 6 shows a powder X-ray diffraction profile of the rapidly solidified alloy obtained in Example 13. From FIG. 6, the presence of the Nd.sub.2Fe.sub.14B phase was already confirmed in a rapidly solidified state.

[0100] The rapidly solidified alloy obtained in the above step was coarsely pulverized to several mm or less to form a rapidly solidified alloy powder, and then, using a flash annealing furnace (crystallization heat treatment furnace, furnace core tube made of transparent quartz and having an outer diameter of 15 mm×an inner diameter of 12.5 mm×a length of 1000 mm, a heating zone of 300 mm, a cooling zone of 500 mm by a cooling fan), the coarse powder of the rapidly solidified alloy was put into a raw material hopper and heat treatment was performed at a workpiece cutting speed of 20 g/min. Note that furnace core tube inclination angle, furnace core tube rotation speed, and furnace core tube vibration frequency were appropriately adjusted together with the heat treatment temperature and the heat treatment time shown in Table 2 so as to achieve the temperature rising rate shown in Table 2. As a result, the rapidly solidified alloy powder passes through the furnace core tube while performing a movement in which stirring by the furnace core tube rotational movement and a hopping phenomenon by the furnace core tube vibration are combined, so that the rapidly solidified alloy powder was placed under a specific heat treatment condition in which the rapidly solidified alloy powder receives thermal history not integrally but individually. Examples of the heat treatment furnace and the thermal history in the performing flash annealing are shown in FIG. 2A and FIG. 2B, and FIG. 3, respectively.

[0101] The constituent phase of the rapidly solidified alloy powder after the flash annealing (crystallization heat treatment) was confirmed by powder X-ray diffraction, and the presence of the Nd.sub.2Fe.sub.14B phase was confirmed. As a representative example, FIG. 7 shows a powder X-ray diffraction profile of the rapidly solidified alloy after flash annealing (crystallization heat treatment) obtained in Example 13. In addition, a peak of α-Fe that was not observed in FIG. 6 was observed in FIG. 7 after flash annealing (crystallization heat treatment), and it was confirmed to be a metal structure in which the Nd.sub.2Fe.sub.14B phase and the α-Fe phase were mixed.

[0102] As a representative example, FIG. 4 shows a bright field image and elemental mapping obtained by observing the iron-based rare earth boron-based isotropic magnet alloy obtained in Example 13 with a transmission electron microscope. From the bright field image, the presence of a Nd.sub.2Fe.sub.14B phase with an average crystal grain size of 50 nm or less and a clear grain boundary phase surrounding the Nd.sub.2Fe.sub.14B phase was confirmed. In addition, in the elemental mapping, it could be confirmed that a grain boundary phase in which Nd and Fe were concentrated was present at the crystal grain boundary of a main phase composed of the main constituent elements of Nd, Fe, and B, and, it was presumed that Fe present at the grain boundary is present as an α-Fe phase, based on the results of the powder X-ray diffraction. It has been confirmed by the present inventor that the grain boundary phase as shown in FIG. 4 is formed in all Examples.

[0103] The iron-based rare earth boron-based isotropic magnet alloys obtained by performing the flash annealing (crystallization heat treatment) described in Table 2 were made into samples for evaluation of magnetic properties with a length of about 7 mm×a width of about 0.9 mm to 2.3 mm or less×a thickness of 18 μm to 25 μm, and then magnetized in the longitudinal direction by a pulse application magnetic field of 3.2 MA/m. Thereafter, the sample for evaluation of magnetic properties was set in the longitudinal direction in order to suppress the influence of demagnetizing field, and the results of measuring room temperature magnetic properties with a vibrating sample magnetometer (VSM) are shown in Table 3. From Table 3, it was found that magnetic properties of a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1400 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more described above were obtained by the alloy composition and manufacturing method described in Examples 1 to 39. In particular, it was found that the intrinsic coercive force HcJ of Examples 32 to 39 containing Pr was higher than that of Examples 1 to 31.

[0104] Subsequently, the magnetic powder subjected to flash annealing (crystallization heat treatment) obtained in Example 13 was pulverized with a pin disc mill so as to have an average grain size of 125 μm. Then, 2 mass % of an epoxy resin diluted with methyl ethyl ketone (MEK) was added to the pulverized magnetic powder, and the mixture was mixed and kneaded. Thereafter, 0.1 mass % of calcium stearate was added thereto as a lubricant to prepare a compound for a compression-molded bonded magnet.

[0105] The compound for a compression-molded bonded magnet was compression molded at a pressure of 1568 MPa (16 ton/cm.sup.2) to obtain a compression molded body having a shape of a diameter of 10 mm×a height of 7 mm, and then this compression molded body was subjected to a curing heat treatment (curing) at 180° C. for 1 hour in an argon gas atmosphere to obtain an isotropic compression-molded bonded magnet. Since the obtained isotropic compression-molded bonded magnet had a molded body density of 6.3 g/cm.sup.3 (true specific gravity of magnetic powder: 7.5 g/cm.sup.3), the magnetic powder filling rate was 84 vol %.

[0106] The magnetic properties of the isotropic compression-molded bonded magnet obtained using the magnetic powder of Example 13 were measured by a BH tracer after being magnetized in the longitudinal direction with a pulse applied magnetic field of 3.2 MA/m, and it was found that the isotropic compression-molded bonded magnet exhibits magnetic properties of a residual magnetic flux density Br of 0.74 T, an intrinsic coercive force HcJ of 1028 kA/m, and a maximum energy product (BH) max of 89.4 kJ/m.sup.3.

[0107] Next, the magnetic powder subjected to flash annealing (crystallization heat treatment) obtained in Example 13 was pulverized with a pin disc mill so as to have an average grain size of 75 μm. Then, the pulverized magnetic powder was subjected to a coupling treatment by spraying a titanate-based coupling agent so as to be 0.75 mass % while heating and stirring the pulverized magnetic powder, 0.5 mass % of stearic acid amide as a lubricant and 4.75 mass % of nylon 12 resin powder were added and mixed, and then a compound for an injection-molded bonded magnet was prepared at an extrusion temperature of 170° C. using a continuous extrusion kneader.

[0108] Using the compound for an injection-molded bonded magnet, injection molding was performed at an injection temperature of 250° C. to prepare an isotropic injection-molded bonded magnet having a shape of a diameter of 10 mm×a height of 7 mm. Since the obtained isotropic injection-molded bonded magnet had a molded body density of 4.6 g/cm.sup.3 (true specific gravity of magnetic powder: 7.5 g/cm.sup.3), the magnetic powder filling factor was 61 vol %.

[0109] The magnetic properties of the isotropic injection-molded bonded magnet obtained using the magnetic powder of Example 13 were measured by a BH tracer after being magnetized in the longitudinal direction with a pulse application magnetic field of 3.2 MA/m, and it was found that the isotropic injection-molded bonded magnet exhibits magnetic properties of a residual magnetic flux density Br of 0.54 T, an intrinsic coercive force HcJ of 1014 kA/m, and a maximum energy product (BH) max of 63.4 kJ/m.sup.3, and magnetic properties equivalent to those of a general isotropic Nd—Fe—B compression-molded bonded magnet were obtained even by injection molding.

Comparative Example

[0110] In order to obtain the alloy composition shown in Table 1, 100 g of a raw material in which additive elements such as Co, Si, Ti, and Zr were blended in addition to main elements of Nd, Dy, B, and Fe with a purity of 99.5% or more was put into an alumina melting crucible, and then set in a work coil in a vacuum melting furnace. Then, the inside of the vacuum melting furnace was evacuated to 0.02 Pa or less, argon gas was then introduced to normal pressure, and a molten alloy was formed by high frequency induction heating. Thereafter, a molten alloy was cast into a water-cooled copper mold to prepare a mother alloy.

[0111] Subsequently, the obtained mother alloy was divided into an appropriate size, and then 40 g of the mother alloy was inserted into a transparent quartz nozzle having, at the bottom, an orifice with an appropriately different diameter (0.7 mm to 1.2 mm) so as to have an average metal tapping rate (in Table 1, simply shown as “metal tapping rate”) described in Table 1, and then the mother alloy was set in a work coil in a single roll quenching device. Then, the inside of the vacuum melting furnace was evacuated to 0.02 Pa or less, argon gas was then introduced until reaching the quenching atmospheric pressure shown in Table 1, the mother alloy was redissolved by high-frequency induction heating, and the molten alloy was tapped from a nozzle orifice at an injection pressure of 30 kPa onto the surface of the rotating roll rotating at the roll surface velocity (Vs) shown in Table 1 to prepare a rapidly solidified alloy. At this time, the distance between the tip of the nozzle and the surface of the rotating roll was set to 0.8 mm.

[0112] The rapidly solidified alloy obtained in the above step was coarsely pulverized to several mm or less to form a rapidly solidified alloy powder, and then, using a flash annealing furnace (crystallization heat treatment furnace, furnace core tube made of transparent quartz and having an outer diameter of 15 mm×an inner diameter of 12.5 mm×a length of 1000 mm, a heating zone of 300 mm, a cooling zone of 500 mm by a cooling fan), the coarse powder of the rapidly solidified alloy was put into a raw material hopper and heat treatment was performed at a workpiece cutting speed of 20 g/min. Note that furnace core tube inclination angle, furnace core tube rotation speed, and furnace core tube vibration frequency were appropriately adjusted together with the heat treatment temperature and the heat treatment time shown in Table 2 so as to achieve the temperature rising rate shown in Table 2.

[0113] The constituent phase of the rapidly solidified alloy powder after the flash annealing (crystallization heat treatment) was confirmed by powder X-ray diffraction, and the presence of the Nd.sub.2Fe.sub.14B phase was confirmed. As a representative example, FIG. 8 shows a powder X-ray diffraction profile of the rapidly solidified alloy after flash annealing (crystallization heat treatment) obtained in Comparative Example 7. From FIG. 8, it was confirmed that Comparative Example 7 is a single-phase metal structure having the Nd.sub.2Fe.sub.14B phase as a main phase.

[0114] As a representative example, FIG. 5 shows a bright field image and elemental mapping obtained by observing the iron-based rare earth boron-based isotropic magnet alloy obtained in Comparative Example 7 with a transmission electron microscope. In the bright field image, the Nd.sub.2Fe.sub.14B phase with an average crystal grain size of 50 nm or less could be confirmed, but a clear grain boundary phase could not be confirmed. In addition, also from the elemental mapping, it was found that there was no grain boundary phase in which Nd and Fe were concentrated as seen in Example 13 at the crystal grain boundary of the main phase composed of the main constituent elements of Nd, Fe, and B. The same applies to the other comparative examples.

[0115] The iron-based rare earth boron-based isotropic magnet alloys obtained by performing the flash annealing (crystallization heat treatment) described in Table 2 were made into samples for evaluation of magnetic properties with a length of about 7 mm×a width of about 0.9 mm to 2.3 mm×a thickness of 18 μm to 25 μm, and then magnetized in the longitudinal direction by a pulse application magnetic field of 3.2 MA/m. Thereafter, the sample for evaluation of magnetic properties was set in the longitudinal direction in order to suppress the influence of demagnetizing field, and the results of measuring room temperature magnetic properties with a vibrating sample magnetometer (VSM) are shown in Table 3. From Table 3, it was found that magnetic properties of a residual magnetic flux density Br of 0.85 T or more, an intrinsic coercive force HcJ of 700 kA/m to less than 1400 kA/m, and a maximum energy product (BH) max of 120 kJ/m.sup.3 or more described above were not obtained by the alloy composition and manufacturing method described in Comparative Examples 1 to 12.

TABLE-US-00001 TABLE 1 Quenching Metal Roll surface Alloy composition atmospheric tapping velocity (atom %) pressure (kPa) rate (g/min) (m/sec) Ex-  1 Nd.sub.11.5Fe.sub.84.3B.sub.4.2 61.3 340 33 ample  2 Nd.sub.12Fe.sub.83.2B.sub.4.8 41.3 430 30  3 Nd.sub.11.5Fe.sub.84.3B.sub.5 41.3 430 30  4 Nd.sub.12Fe.sub.82.4B.sub.5.6 31.3 600 22  5 Nd.sub.13Fe.sub.82.2B.sub.4.8 41.3 430 30  6 Nd.sub.11.5Pr.sub.0.5Fe.sub.83.2B.sub.4.8 81.3 510 30  7 Nd.sub.11.6Dy.sub.0.4Fe.sub.83B.sub.5 31.3 510 25  8 Nd.sub.12.1Fe.sub.83B.sub.4.6C.sub.0.3 41.3 430 22  9 Nd.sub.12Fe.sub.83.1B.sub.4.7C.sub.0.2 41.3 430 22 10 Nd.sub.12Fe.sub.83.1B.sub.2.5C.sub.2.4 41.3 510 25 11 Nd.sub.12.1Fe.sub.82.6B.sub.4.6C.sub.0.3Al.sub.0.4 21.3 510 22 12 Nd.sub.12.1Fe.sub.82.4B.sub.5Al.sub.0.5 41.3 510 30 13 Nd.sub.12.7Fe.sub.81.9B.sub.4.9Ti.sub.0.5 21.3 600 35 14 Nd.sub.12Fe.sub.82.5B.sub.5Si.sub.0.5 41.3 510 27 15 Nd.sub.11.8Fe.sub.80.9Co.sub.2B.sub.4.8Ti.sub.0.5 41.3 600 35 16 Nd.sub.12Fe.sub.82B.sub.5Ga.sub.1 41.3 430 30 17 Nd.sub.12Fe.sub.82B.sub.5V.sub.1 41.3 860 30 18 Nd.sub.12Fe.sub.82B.sub.5Cr.sub.1 41.3 600 30 19 Nd.sub.12Fe.sub.80B.sub.5Ag.sub.3 41.3 430 25 20 Nd.sub.12Fe.sub.79B.sub.5Mn.sub.4 41.3 1290 30 21 Nd.sub.12Fe.sub.78B.sub.5Ta.sub.5.0 41.3 430 30 22 Nd.sub.12.1Fe.sub.84.2B.sub.5Cu.sub.0.5 41.3 430 25 23 Nd.sub.11.8Fe.sub.81.1Co.sub.2B.sub.4.8Pt.sub.0.3 41.3 600 30 24 Nd.sub.11.8Fe.sub.81.1Co.sub.2B.sub.4.8Au.sub.0.3 41.3 600 30 25 Nd.sub.12Fe.sub.82.5B.sub.5Zn.sub.0.5 41.3 1290 35 26 Nd.sub.12Fe.sub.82.5B.sub.5Zr.sub.0.5 41.3 860 24 27 Nd.sub.12Fe.sub.82.5B.sub.5Nb.sub.0.5 41.3 860 22 28 Nd.sub.12Fe.sub.82B.sub.5Hf.sub.1 41.3 600 30 29 Nd.sub.12Fe.sub.82B.sub.5Mo.sub.1 41.3 600 30 30 Nd.sub.12Fe.sub.82.5B.sub.5Pb.sub.0.5 41.3 430 25 31 Nd.sub.12Fe.sub.82B.sub.5W.sub.2 41.3 1800 35 32 Nd.sub.6.1Pr.sub.6Fe.sub.81B.sub.4.9Nb.sub.2 41.3 770 20 33 Nd.sub.10.4Pr.sub.2Fe.sub.82.7B.sub.4.9 41.3 770 23 34 Nd.sub.8.4Pr.sub.4Fe.sub.82.7B.sub.4.9 41.3 770 25 35 Nd.sub.6.4Pr.sub.6Fe.sub.82.7B.sub.4.9 41.3 770 23 36 Nd.sub.8.4Pr.sub.4Fe.sub.82.2B.sub.4.9Ga.sub.0.5 41.3 770 25 37 Nd.sub.8.8Pr.sub.3.8Fe.sub.82.5B.sub.4.9 41.3 770 23 38 Nd.sub.8.1Pr.sub.4.3Fe.sub.82.7B.sub.4.9 41.3 770 22 39 Nd.sub.7.2Pr.sub.4.8Fe.sub.82.1B.sub.4.9Nb.sub.1 41.3 770 20 Com-  1 Nd.sub.12Fe.sub.82B.sub.6 101.3 600 30 parative  2 Nd.sub.12Fe.sub.82B.sub.6 41.3 600 30 Ex-  3 Nd.sub.12Fe.sub.81B.sub.6Si.sub.1 41.3 600 25 ample  4 Nd.sub.11.6Dy.sub.0.4Fe.sub.82B.sub.6 31.3 600 25  5 Nd.sub.14Fe.sub.80B.sub.6 41.3 600 30  6 Nd.sub.12Fe.sub.80Co.sub.2B.sub.6 41.3 600 30  7 Nd.sub.11.7Fe.sub.80.5B.sub.6.5Nb.sub.1.3 21.3 860 35  8 Nd.sub.9Fe.sub.84B.sub.6Ti.sub.1 21.3 600 30  9 Nd.sub.10.5Fe.sub.83B.sub.6Ti.sub.0.5 41.3 600 35 10 Nd.sub.10Fe.sub.81B.sub.9 41.3 860 25 11 Nd.sub.9Fe.sub.80B.sub.7Ti.sub.1Zr.sub.3 61.3 430 12 12 Nd.sub.4Fe.sub.77.5B.sub.18.5 41.3 2150 10

TABLE-US-00002 TABLE 2 Temperature Heat treatment Heat rising rate temperature treatment (° C./sec) (° C.) time (sec) Example  1 120 620 5  2 70 640 10  3 70 640 10  4 70 640 10  5 130 650 5  6 130 640 5  7 120 620 5  8 130 650 5  9 130 650 5 10 130 640 5 11 125 640 5 12 125 620 5 13 125 620 5 14 125 640 5 15 130 630 5 16 130 650 5 17 130 660 5 18 140 670 5 19 130 630 5 20 130 650 5 21 125 620 5 22 130 630 5 23 160 630 4 24 180 630 3.5 25 180 650 3.5 26 140 670 5 27 130 660 5 28 30 660 20 29 25 650 20 30 70 660 10 31 125 620 5 32 90 680 15 33 90 670 15 34 90 670 15 35 90 670 15 36 90 680 15 37 90 650 15 38 90 680 15 39 90 660 15 Com-  1 4 660 180 parative  2 60 670 10 Example  3 60 650 10  4 130 660 5  5 70 680 10  6 130 660 5  7 140 680 5  8 80 735 10  9 70 690 10 10 140 679 5 11 70 680 10 12 60 620 10

TABLE-US-00003 TABLE 3 Magnetic properties Br HcJ (BH)max (T) (kA/m) (kJ/m.sup.3) Example  1 0.93 720.8 132.5  2 0.92 801.8 134.2  3 0.89 739.4 122.1  4 0.88 868.1 122.9  5 0.87 976.2 124.5  6 0.90 832.3 131.4  7 0.89 1030.8 122.0  8 0.91 1044.2 134.2  9 0.88 814.0 126.6 10 0.87 1088.1 120.3 11 0.88 934.7 123.4 12 0.91 980.0 128.9 13 0.88 1040.4 128.3 14 0.89 985.6 124.5 15 0.92 995.6 137.2 16 0.92 942.6 130.1 17 0.87 1021.6 121.1 18 0.88 1038.2 123.9 19 0.89 966.1 127.2 20 0.87 1094.7 123.6 21 0.87 1133.6 120.2 22 0.91 990.7 128.8 23 0.93 975.9 132.5 24 0.92 943.2 129.7 25 0.89 998.8 126.9 26 0.88 1012.4 125.4 27 0.88 1003.2 124.7 28 0.87 1106.1 122.5 29 0.88 1053.5 121.1 30 0.88 1029.4 128.3 31 0.90 910.6 127.6 32 0.86 1202.0 120.3 33 0.88 1116.3 122.4 34 0.87 1211.0 124.8 35 0.87 1241.8 121.6 36 0.88 1181.8 124.2 37 0.89 1010.2 123.4 38 0.86 1057.8 120.0 39 0.86 1014.3 123.6 Comparative  1 0.82 735.6 110.8 Example  2 0.84 751.5 118.3  3 0.84 700.7 122.1  4 0.82 873.5 119.3  5 0.80 1250.8 111.4  6 0.85 723.1 123.6  7 0.84 978.3 120.2  8 0.92 569.4 121.1  9 0.92 628.8 124.5 10 0.93 558.2 123.7 11 0.89 664.5 124.8 12 0.88 410.6 116.3

REFERENCE SIGNS LIST

[0116] 1 raw material hopper [0117] 2 raw material supply feeder [0118] 3 furnace core tube [0119] 3a furnace core tube enlarged view [0120] 3b furnace core tube cross-sectional enlarged view [0121] 4 tubular furnace [0122] 5 cooling tower [0123] 6 collection hopper [0124] 7 vibrator [0125] 8 furnace core tube rotating motor [0126] 9 furnace core tube rotation axis [0127] 10 device frame [0128] 11 furnace core tube inclination angle [0129] 12 cooling fan air [0130] 13 rapidly solidified alloy powder (workpiece) [0131] 14 moving direction of workpiece [0132] 15 hopping phenomenon of workpiece [0133] 16 temperature rising rate [0134] 17 holding temperature [0135] 18 temperature drop rate [0136] 21 main phase [0137] 22 grain boundary phase