JP2005183630A - Permanent magnetic powder, method for manufacturing the same and bond magnet - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide permanent magnetic powder for easily generating a ThMn<SB>12</SB>phase even at the time of using Nd as rear earth elements, and to provide a method for manufacturing the permanent magnetic powder. <P>SOLUTION: This permanent magnetic powder is constituted of a general formula R(Fe<SB>100-y-w</SB>Co<SB>w</SB>Ti<SB>y</SB>)<SB>x</SB>Si<SB>z</SB>A<SB>v</SB>(R is one or two types or more of rear earth elements, and its 50 mol% or more is Nd, and A is one or two types of N and C), wherein the mol rate of the general formula is x=10 to 12.5, y=(8.3-1.7×z) to 12, z=0.2 to 2, v=0.1 to 3, and w=0 to 30. In this case, the permanent magnetic powder is characterized by having composition satisfying (Fe+Co+Ti+Si)/R>12, and consisting of the group of particles whose mean crystal particle diameter is 200nm or less. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  The present invention relates to a magnet powder that is a material of a permanent magnet used in an electric device such as a motor, and more particularly to a magnet powder having a high remanent magnetization and a coercive force that is optimal as a material for a bonded magnet and the magnet.

Conventionally, SmCo magnet powder and NdFeB magnet powder are known as permanent magnet powders used for bond magnets and the like, and in particular, magnet powder having Nd ₂ Fe ₁₄ B ₁ phase obtained by rapid solidification method is widely used. .
These magnet powders that are currently in practical use use expensive rare earth metals, and in particular, Sm is expensive, and NdFeB magnet powder composed of cheap Nd contains about 30 wt% Nd. For the purpose of reducing raw material costs, a magnet material having a composition based on an inexpensive raw material is desired.

Among them, rare earth-iron-based magnet materials having a body-centered tetragonal or ThMn ₁₂ type crystal structure are disclosed in, for example, Patent Document 1 (Japanese Patent Laid-Open No. 63-273303) and Patent Document 2 (Japanese Patent Laid-Open No. 5-65603). ) And Patent Document 3 (Japanese Patent Laid-Open No. 2000-114017).
Patent Document 1 has the formula RxTiyAzFeaCob (wherein R is a rare earth element including Y, A is one or more of each of B, C, Al, Si, P, Ga, Ge, Sn, S, and N, in weight percentage) x is 12 to 30%, y is 4 to 10%, z is 0.1 to 8%, a is 55 to 85%, and b is 34% or less). Patent Document 1 states that the A element enters between atoms and changes the distance between Fe in a preferable direction.

  In Patent Document 2, R is Y, Th and a combination of one or more elements selected from the group consisting of all lanthanoid elements, and X is N (nitrogen) or B (boron) or C (carbon). ) Or a combination of these elements, the atomic percentage includes R: 3 to 30%, X: 0.3 to 50%, the balance is substantially composed of Fe, and the main phase is a body-centered tetragonal structure. An iron-rare earth permanent magnet material having the following is disclosed. In Patent Document 2, a part of Fe is selected from the group consisting of M elements (Ti, Cr, V, Zr, Nb, Al, Mo, Mn, Hf, Ta, W, Mg, Si, Sn, Ge, Ga). It is further proposed to include M: 0.5 to 30% by atomic percent by substitution with a selected one or a combination of two or more elements. In Patent Document 2, the element M is positioned as an element having a great effect in generating a body-centered tetragonal structure.

Further, Patent Document 3, the general formula _{(R 1-u M u)} (Fe 1-vw Co v T w) x A y (R in the formula, M, T, A, respectively R: rare earth containing Y At least one element selected from elements, at least one element selected from M: Ti, Nb, T: at least one element selected from Ni, Cu, Sn, V, Ta, Cr, Mo, W, Mn, A : At least one element selected from Si, Ge, Al, and Ga, and u, v, w, x, and y are 0.1 ≦ u ≦ 0.7, 0 ≦ v ≦ 0.8, 0, respectively. ≦ w ≦ 0.1, 5 ≦ x ≦ 12, and 0.1 ≦ y ≦ 1.5), and the main hard magnetic phase has a ThMn ₁₂ type crystal structure. Disclosed material. Patent Document 3 states that the amount of Si, Ge, and the like, which are stabilizing elements of a phase having a ThMn ₁₂ type crystal structure (hereinafter referred to as ThMn ₁₂ phase) can be reduced by substituting R element with M element. .

JP-A-63-273303 JP-A-5-65603 JP 2000-1114017 A

While rare earth permanent magnets are required to have high characteristics, they are also required to be low in cost. Among the rare earth elements constituting the rare earth permanent magnet, Nd is less expensive than Sm. Therefore, it is desirable that Nd, which is cheaper than expensive Sm, is the main element of the rare earth element. However, when Nd is used, it is difficult to produce a ThMn ₁₂ phase, and many nonmagnetic impurities and high-temperature, long-time heat treatment are required for its production. In addition, sufficient characteristics could not be obtained because many nonmagnetic impurities were used. For example, in the above-mentioned Patent Document 2, annealing is performed at 900 ° C. for 7 days, and in Patent Document 3, only Sm is used as a rare earth element with some exceptions.
Therefore, an object of the present invention is to provide a permanent magnet powder that can easily generate a ThMn ₁₂ phase even when Nd is used as a rare earth element, and a method for producing the same. Another object of the present invention is to provide a bonded magnet using such permanent magnet powder.

The present inventor can easily produce a hard magnet for a permanent magnet even when a phase having a ThMn ₁₂ type crystal structure uses Nd as a rare earth element by simultaneously adding a predetermined amount of Ti, Si and N. It has been found that sufficient characteristics can be obtained as a composition. However, this composition had a coercive force as small as several hundred Oe, which was insufficient as a practical permanent magnet powder.
Therefore, the present inventors have found that when the crystal structure of the composition is refined, a sufficient coercive force can be expressed as a permanent magnet powder.

The present invention is a permanent magnet powder based on the above, the general formula _{R (Fe 100-yw Co w} Ti y) x Si z A v ( provided that 50 together with R is one or more rare earth elements The molar ratio is Nd, A is one or two of N and C), and the molar ratio of the general formula is x = 10 to 12.5, y = (8.3-1.7 × z) To 12, z = 0.2 to 2, v = 0.1 to 3, w = 0 to 30 and a composition satisfying (Fe + Co + Ti + Si) / R> 12, and an average crystal grain size of 200 nm or less It is characterized by comprising a set of particles.

In the permanent magnet powder of the present invention, it is desirable that each particle constituting the powder has a phase having a ThMn ₁₂ type crystal structure as a main phase, and particularly has a single-phase structure substantially having a phase having a ThMn ₁₂ type crystal structure.
Further, in the permanent magnet powder of the present invention, even if it is a form in which Nd occupies 70 mol% or more of R, a single-phase structure of a phase having a ThMn ₁₂ type crystal structure can be obtained, which is advantageous for cost reduction. It is.

As described above, the permanent magnet powder of the present invention is characterized by having a fine crystal structure. Such a fine crystal structure is realized by subjecting the amorphous or fine crystalline powder subjected to rapid solidification to a predetermined heat treatment. Accordingly, the present invention has an average particle size of the general formula _{R (Fe 100-yw Co w} Ti y) x Si z ( wherein, R represents one or more than with a more thereof 50 mol% Nd rare earth element) The molar ratio of the general formula is x = 10 to 12.5, y = (8.3-1.7 × z) to 12, z = 0.2 to 2, w = 0 to 30. In addition, a powder having a composition satisfying (Fe + Co + Ti + Si) / R> 12 and subjected to rapid solidification treatment is prepared, and this powder is 0.5 in a temperature range of 600 to 850 ° C. in an inert atmosphere. Provided is a method for producing a permanent magnet powder, characterized by performing a heat treatment for holding for 120 hours, and subjecting the heat treated powder to a nitriding treatment or a carbonizing treatment.

  In the method for producing a permanent magnet powder of the present invention, the powder subjected to the rapid solidification treatment exhibits a structure of any of an amorphous phase, a mixed phase of an amorphous phase and a crystalline phase, or a crystalline phase. Among these, it is desirable to make the mixed phase of the amorphous phase and the crystalline phase, particularly the rich phase of the crystalline phase, from the viewpoint of easy control of the crystal grain size after the subsequent heat treatment.

In the method for producing the permanent magnet powder of the present invention, the specific method of the rapid solidification treatment is not limited. However, it is possible to apply the single roll method to obtain a desired structure stably after cooling and solidification. Desirable for reasons. When applying the single roll method, the peripheral speed of the roll is preferably 10 to 100 m / s. Although there are some differences depending on other conditions such as the composition of the alloy to be obtained, the nozzle hole diameter for discharging the molten metal, the roll material, etc., the rapidly solidified powder in this range is an amorphous phase, a mixed phase of an amorphous phase and a crystalline phase. Alternatively, it can exhibit any texture of the crystalline phase.
In the method for producing a permanent magnet powder of the present invention, the heat treatment performed on the powder subjected to the rapid solidification treatment crystallizes the amorphous phase or adjusts the particle size of the crystal particles constituting the crystal phase. Become.

By using the permanent magnet powder obtained by the present invention, a bonded magnet can be produced. This bonded magnet includes a permanent magnet powder and a resin phase in which the permanent magnet powder is dispersed. The crystalline hard magnetic particles constituting the permanent magnet powder are represented by the general formula R (Fe _100-yw Co _w Ti _y ) _x Si _z A _v (wherein R is one or more of rare earth elements and 50 mol% or more thereof is Nd, and A is one or two of N and C). The molar ratio is x = 10 to 12.5, y = (8.3-1.7 × z) to 12, z = 0.2 to 2, v = 0.1 to 3, w = 0 to 30. In addition, it is characterized by satisfying the composition of (Fe + Co + Ti + Si) / R> 12.

  It is desirable from the viewpoint of magnetic properties that the hard magnetic particles in the bonded magnet of the present invention have an average crystal grain size of 200 nm or less.

According to the present invention, it is possible to provide a permanent magnet powder and a manufacturing method thereof capable of easily generating a ThMn ₁₂ phase even when using Nd as a rare earth element. Moreover, this invention can obtain the bonded magnet using such a permanent magnet powder.

Hereinafter, embodiments including the best mode of the permanent magnet powder, the manufacturing method thereof, and the bonded magnet of the present invention will be described.
First, the reasons for limiting each element in the present invention will be described.
<R (rare earth element)>
R is an element essential for obtaining high magnetic anisotropy. In order to produce a ThMn ₁₂ phase as a hard magnetic phase, it is advantageous to use Sm. However, in the present invention, Nd accounts for 50 mol% or more of R in order to obtain cost merit. The present invention makes it possible to easily produce a ThMn ₁₂ phase while occupying 50 mol% or more of R with Nd. However, the present invention allows rare earth elements to be included in addition to Nd. In that case, one or more of Y, La, Ce, Pr, and Sm can be included together with Nd. Thus, the rare earth element in this invention has the concept containing Y. Among these, Pr exhibits a property substantially equivalent to Nd, and therefore, a value equivalent to Nd can be obtained also in characteristics, which is particularly preferable. According to the present invention, even when the ratio of Nd in R is as high as 70 mol% or more, or 90 mol% or more, a single-phase structure composed of a ThMn ₁₂ phase that is a hard magnetic phase can be obtained. As shown in the experimental examples to be described later, according to the present invention, even when R is only Nd, that is, when Nd occupies 100% of R, a single-phase structure composed of a ThMn ₁₂ phase that is a hard magnetic phase is obtained. be able to.

<Si>
Si is added to R (Nd) and Fe simultaneously with Ti, thereby contributing to stabilization of the ThMn ₁₂ phase as a hard magnetic phase. At this time, Si has an effect of entering the ThMn ₁₂ phase lattice and reducing the crystal lattice. When the amount of Si is less than 0.2 (molar ratio, the same applies hereinafter), a phase having a Mn ₂ Th ₁₇ type crystal structure (hereinafter referred to as Mn ₂ Th ₁₇ phase) precipitates, and when it exceeds 2.3, α-Fe is reduced. It tends to precipitate. Therefore, in the present invention, it is recommended that z, which is the amount of Si, be in the range of 0.2 to 2.3. A desirable Si amount (z) is 0.2 to 2, and a more desirable Si amount (z) is 0.2 to 1. Si is included so as to satisfy the relationship of Fe, Co, Ti, and R (Fe molar ratio + Co molar ratio + Ti molar ratio + Si molar ratio) / (R molar ratio)> 12. However, this point will be described later.

Ti contributes to the formation of a ThMn ₁₂ phase by substituting Fe. In order to obtain this effect sufficiently, it is necessary to set the lower limit of the Ti amount (y) in relation to the Si amount. That is, as shown in an experimental example described later, when the Ti amount (y) is less than (8.3-1.7 × z (Si amount)), α-Fe and Mn ₂ Th ₁₇ phases are precipitated. Further, when the Ti amount (y) exceeds 12, the saturation magnetization is remarkably reduced. Therefore, in the present invention, the Ti amount (y) is set to (8.3-1.7 × z (Si amount)) to 12. Desirable Ti amount (y) is (8.3-1.7 × z (Si amount)) to 12, and more desirable Ti amount (y) is (8.3-1.7 × z (Si amount)) to 9. It is.
Further, when the sum (x) of the Fe amount and the Ti amount is less than 10, both the saturation magnetization and the anisotropic magnetic field are low, and α-Fe exceeding 12.5 is precipitated. Therefore, in the present invention, the sum (x) of the Fe amount and the Ti amount is set to 10 to 12.5. Desirable x is 11 to 12.5.

<A (one or two of N (nitrogen) and C (carbon))>
A is to expand the lattice of ThMn ₁₂ phase by entering the interstitial of ThMn ₁₂ phase is an element effective for improving the magnetic properties. However, if the amount of A (v) exceeds 3, precipitation of α-Fe is observed. On the other hand, if it is 0.1 or less, the effect of improving the magnetic properties cannot be obtained sufficiently. Therefore, in the present invention, the A amount (v) is set to 0.1 to 3. A desirable A amount (v) is 0.3 to 2.5, and a more desirable A amount (v) is 1 to 2.5.

<Fe, Fe-Co>
The hard magnetic composition according to the present invention substantially contains Fe other than the above elements, but it is effective to substitute a part of Fe with Co. As will be described in an experimental example to be described later, by adding Co, the saturation magnetization (σ _s ) and the anisotropic magnetic field ( _HA ) increase. The amount of Co is desirably added in a molar ratio of 30 or less, and more desirably in the range of 5 to 20. Note that the addition of Co is not essential.

<(Fe molar ratio + Co molar ratio + Ti molar ratio + Si molar ratio) / (R molar ratio)>12>
The individual contents of Fe, Co, Ti, and Si are as described above, but satisfying the condition of (Fe + Co + Ti + Si) / R> 12 when the hard magnetic composition of the present invention has a single phase structure of ThMn ₁₂ phase. It is important to. As shown in an experimental example to be described later, the saturation magnetization is low when the above conditions are not satisfied.

The experimental results that serve as the basis for the reasons for limiting the composition range described above are shown below.
<Experimental example 1>
First, the experimental result (Experiment 1) of the Z value (Si amount) dependence on the phase state and magnetic characteristics will be described.
A high purity Nd, Fe, Ti, Si metal is used as a raw material, and a sample is obtained by an arc melting method in an Ar atmosphere so that the alloy composition is Nd— (Ti _8.3 Fe _91.7 ) ₁₂ —Si _z. Was made. Subsequently, this alloy was pulverized by a stamp mill and passed through a sieve having an opening of 38 μm, and then heat treatment (nitriding) was performed in a nitrogen atmosphere at a temperature of 430 to 520 ° C. for 100 hours. About each sample after heat processing, while performing the chemical composition analysis and the identification of the phase comprised, the saturation magnetization ((sigma) s) and the anisotropic magnetic field ( _HA ) were measured. The results are shown in Table 1.

In addition, identification of the phase comprised was performed based on the measurement of the X ray diffraction method and the thermomagnetic curve. X-ray diffraction was measured using a Cu tube at an output of 15 kW, and the presence or absence of peaks other than the ThMn ₁₂ phase was confirmed. However, since the peak of the Mn ₂ Th ₁₇ phase substantially coincides with the peak of the ThMn ₁₂ phase, it may be difficult to confirm only by the X-ray diffraction method. Also used. Further, the thermomagnetic curve was measured by applying a magnetic field of 2 kOe, and the presence or absence of expression of Tc (Curie temperature) corresponding to a phase other than the ThMn ₁₂ phase was confirmed. In the present invention, “the single phase structure of the ThMn ₁₂ phase” means that no peak of the phase other than the ThMn ₁₂ phase is observed by the above X-ray diffraction method, and the above thermomagnetic curve is measured. Tc corresponding to a phase other than the ThMn ₁₂ phase is not confirmed by the above, and the magnetization remaining on the higher temperature side than the Tc is 0.05 or less, including inevitable impurities and unreacted substances that are not detected. It does not matter. For example, in arc melting, the thermal uniformity during melting is insufficient, a slight unreacted phase (for example, Nd, α-Fe, etc.) may remain, and Cu from the sample holder is unavoidable. Although it may be contained as an impurity, it is not considered unless it is detected by measurement of X-ray diffraction and thermomagnetic curve. Specific examples relating to the identification of the phases to be configured will be described with reference to FIGS.

5 shows a sample No. described later. Although the charts show X-ray diffraction measurement results of 4, 7, and 45, only the peak showing the ThMn ₁₂ phase was observed. However, Sample No. 7, the peak of α-Fe can be confirmed. As described above, since the peak of the Mn ₂ Th ₁₇ phase overlaps the peak of the ThMn ₁₂ phase, the two cannot be distinguished on this chart. In addition, FIG. The thermomagnetic curves of 4, 7, 45 and 33 are shown. ThMn ₁₂ phase Tc exists in the vicinity of 400 ° C. Further, the Tc of the Mn ₂ Th ₁₇ phase (2-17 phase) is confirmed on the lower temperature side than the Tc of the ThMn ₁₂ phase as shown in FIG. 6 (Sample No. 33). Here, Tc other than Tc of the ThMn ₁₂ phase was not confirmed, and when the magnetization remaining on the higher temperature side than this Tc was 0.05 or less, it was recognized as a single phase. That is, sample no. 4 and 45 were identified as single-phase structures of the ThMn ₁₂ phase because Tc other than the Tc of the ThMn ₁₂ phase was not confirmed and the magnetization remaining at a higher temperature than the Tc was 0.05 or less. Sample No. 7, Tc other than the Tc of the ThMn ₁₂ phase was not confirmed, but the magnetization remaining on the higher temperature side than this Tc exceeds 0.05, and from FIG. 5, in addition to the ThMn ₁₂ phase, α-Fe Is identified as deposited. Furthermore, sample no. No. 33, since the Tc of the Mn ₂ Th ₁₇ phase is confirmed and the magnetization remaining on the higher temperature side than the Tc of the ThMn ₁₂ phase exceeds 0.05, the Mn ₂ Th ₁₇ phase and the ThMn ₁₂ phase It is identified that α-Fe is precipitated.
As described above, in both FIG. 5 (X-ray diffraction) and FIG. 6 (thermomagnetic curve), when a phase configuration other than the ThMn ₁₂ phase is not confirmed, in the present invention, a single phase structure of the ThMn ₁₂ phase is used. Define that there is.

Further, the saturation magnetization (σs) and the anisotropic magnetic field (H _A ) are the magnetization curve and the magnetization difficulty in the easy axis direction measured with a maximum applied magnetic field of 20 kOe using a VSM (Vibrating Sample Magnetometer). It is obtained based on the magnetization curve in the axial direction. However, for convenience of measurement, the saturation magnetization (σs) is the maximum magnetization value on the magnetization curve in the easy axis direction. The anisotropic magnetic field (H _A ) was defined as the value of the magnetic field at which the tangent at 10 kOe on the magnetization curve in the hard axis direction intersects the saturation magnetization (σs) value.

As shown in Table 1, sample no. In addition to the ThMn ₁₂ phase (hereinafter referred to as 1-12 phase), Mn ₂ Th ₁₇ phase (hereinafter referred to as 2-17 phase) and α-Fe are present, and in particular, the anisotropic magnetic field (H _A ) is present. Low. On the other hand, sample no. It can be seen that 1 to 5 become a single phase of the 1-12 phase and the 1-12 phase is stabilized. The compositions of these 1-12-phase single-phase can be obtained 130 emu / g or more saturated magnetization ([sigma] s), or more anisotropy field 50kOe (H _A). However, sample no. In No. 7, α-Fe precipitates and the characteristics deteriorate. Further, Sample No. with Fe + Ti amount of less than 10 and Si amount of 2.5. No. 8 significantly decreases both the saturation magnetization (σs) and the anisotropic magnetic field ( _HA ). If α-Fe, which is soft magnetic, is present, that portion generates a reverse magnetic domain with a low magnetic field (demagnetizing field). Therefore, since the coercive force is lowered as a result of easily reversing the magnetic domain of the hard magnetic phase component, the presence of α-Fe is not desirable for a permanent magnet that requires a coercive force.
Sample No. In the range from 1 to 5, the anisotropic magnetic field (H _A ) increases as the Si amount increases, and conversely, the saturation magnetization (σs) tends to increase as the Si amount decreases.

<Experimental example 2>
In the same manner as in Experimental Example 1, a sample was prepared so as to have a composition of Nd— (Ti _8.3 Fe _91.7 ) _x —Si _z —N _1.5 , analysis of the chemical composition, identification of the constituted phase, saturation magnetization (σs ) And anisotropic magnetic field (H _A ) were measured. The results are shown in Table 2.
Experimental Example 2 is for confirming the influence of x (Fe amount + Ti amount) and x + z (Fe amount + Ti amount + Si amount) on the phase structure, saturation magnetization (σs) and anisotropic magnetic field (H _A ). This is the experiment that was conducted.

As shown in Table 2, when x is less than 10 (sample Nos. 17 and 21), the saturation magnetization (σs) is less than 120 emu / g, and z (Si amount) is as low as 1.1. 17, the anisotropic magnetic field (H _A ) is as low as about 30. Conversely, if x exceeds 12.5 (Sample Nos. 20 and 22), α-Fe is precipitated. Even if x is in the range of 10 to 12.5, when x + z is 12 or less (Sample Nos. 18 and 19), the saturation magnetization (σs) is less than 120 emu / g, and the anisotropic magnetic field (H _A ) Is as low as about 30 kOe.
On the other hand, when x is in the range of 10 to 12.5 and x + z exceeds 12 (sample Nos. 9 to 16), saturation magnetization (σs) of 120 emu / g or more, difference of 50 kOe or more. It has a characteristic of isotropic magnetic field ( _HA ), and a 1-12 phase single phase structure can be obtained.

<Experimental example 3>
In the same manner as in Experimental Example _{1 Nd- (Ti y Fe 100-} y) -Si 1.0 -N 1.5, Nd- (Ti y Fe 100-y) -Si 1.5 -N 1.5, Nd- (Ti y Fe 100-y ) A sample was prepared so as to have a composition of -Si _2.0 -N _1.5 , and the chemical composition was analyzed, the phases to be constituted were identified, and the saturation magnetization (σs) and the anisotropic magnetic field ( _HA ) were measured. . The results are shown in Table 3.
Experimental Example 3 is an experiment conducted to confirm the influence of y (Ti amount) on the phase configuration, saturation magnetization (σs), and anisotropic magnetic field (H _A ).

When z (Si content) is 1.0, 1.5 or 2.0, and y (Ti content) is less than (8.3-1.7 × z), α-Fe, 2-17 phase precipitates (Sample Nos. 33, 34, 36 to 38). On the other hand, when y (Ti amount) exceeds 12.5 and 12, the saturation magnetization (σs) decreases to less than 120 emu / g (sample No. 35).
On the other hand, when y (Ti amount) is in the range of (8.3-1.7 × z) to 12, a 1-12 phase single phase, in other words, a hard magnetic phase single phase structure, and A saturation magnetization of 130 or 140 emu / g or more and an anisotropic magnetic field of 50 or 55 kOe or more can be obtained (Sample Nos. 23 to 32).

<Experimental example 4>
In the same manner as in Experimental Example 1, a sample was prepared so as to have a composition of Nd— (Ti _8.3 Fe _91.7 ) ₁₂ —Si _2.0 —N _v , analysis of the chemical composition, identification of the constituted phase, saturation magnetization (σs ) And anisotropic magnetic field (H _A ) were measured. The results are shown in Table 4.
Experimental Example 4 is an experiment conducted to confirm the influence of v (N amount) on the phase configuration, saturation magnetization (σs), and anisotropic magnetic field ( _HA ).

When v (N content) is 0, both the saturation magnetization (σs) and the anisotropic magnetic field ( _HA ) are low (sample No. 43). On the other hand, when v (N amount) exceeds 3.5 and 3, α-Fe precipitates (Sample No. 44).
On the other hand, when v (N content) is in the range of 0.1 to 3, a 1-12 phase single phase, in other words, a hard magnetic phase single phase structure, and a saturation magnetization of 120 emu / g or more, An anisotropic magnetic field of 30 kOe or more can be obtained (Sample Nos. 39 to 42). From the viewpoint of saturation magnetization (σs) and anisotropic magnetic field (H _A ), v (N content) is preferably in the range of 0.5 to 2.7, and more preferably 1.0 to 2.5. .

<Experimental example 5>
Each sample shown in Table 5 was produced in the same manner as in Experimental Example 1, and the identification of the constituted phase, the saturation magnetization (σs), and the measurement of the anisotropic magnetic field ( _HA ) were performed. The results are shown in Table 5.
Experimental Example 5 is an experiment for confirming the dependence of w (Co amount) on Nd— (Ti _8.3 Fe _91.7-w Co _w ) 12-Si _z —N _1.5 .

In both cases where z (Si content) is 0.25 and 1.0, increasing w (Co content) improves saturation magnetization (σs) and anisotropic magnetic field (H _A ), and w (Co It can be seen that when the amount is about 20, the effect becomes a peak. Therefore, considering that Co is expensive, w (Co amount) is preferably 30 or less, and more preferably in the range of 10 to 25. Further, in this range of w (Co amount), the structure is a single phase of 1-12 phase.

<Experimental example 6>
A high purity Nd, Fe, Ti, Si metal is used as a raw material, and the alloy composition is Nd- (Ti _8.3 Fe _91.7-w Co _w ) ₁₂ -Si _z. A sample was prepared by the melt dissolution method. Subsequently, this alloy was pulverized by a stamp mill, passed through a sieve having an opening of 38 μm, mixed with C powder having an average particle size of 1 μm or less, and maintained in an Ar atmosphere at a temperature of 400 to 600 ° C. for 24 hours. A heat treatment was performed. About each sample after heat processing, while analyzing a chemical composition and identifying the phase comprised, the saturation magnetization ((sigma) s) and the anisotropic magnetic field ( _HA ) were measured. The results are shown in Table 6.

By adding C instead of N, it is possible to obtain a 1-12 phase single phase structure, a saturation magnetization of 120 emu / g or more, and an anisotropic magnetic field of 30 kOe or more. At this time, C plays the same role as N.
Moreover, even when 1 to 25% of Nd is substituted with Pr, the same result as in the example can be obtained.

Next, the structure of the permanent magnet powder of the present invention will be described.
The permanent magnet powder of the present invention has an average crystal grain size of 200 nm or less, desirably 100 nm or less, and more desirably 80 nm or less. By having such a fine structure, the present invention can realize a coercive force necessary as a permanent magnet powder. Means for obtaining such a fine structure in the present invention will be described later. The crystal grain size is a value calculated as the diameter of a circle having the same area as that obtained by observing the heat-treated quenched alloy with a TEM and recognizing the individual particles, then obtaining the area of each individual particle by image processing. . The average crystal grain size was measured for about 100 crystal grains per sample, and was taken as the average value of the crystal grain sizes of all the measured particles.
The permanent magnet powder of the present invention having a fine crystal structure has a ThMn ₁₂ phase as a main phase, more preferably a single phase structure of a ThMn ₁₂ phase. Whether the structure is a single-phase structure of ThMn ₁₂ phase is determined according to the above-mentioned criteria.

Next, the manufacturing method of the permanent magnet powder of this invention is demonstrated.
The permanent magnet powder of the present invention is characterized by having a fine crystal structure as described above, and there are several methods for obtaining this fine crystal structure. For example, a method using a molten metal quenching method, a method using mechanical grinding or mechanical alloying, and a method using an HDDR (Hydrogenation-Decomposition-Desorption-Recombination) method. Below, the manufacturing method using the molten metal quenching method is demonstrated.
The manufacturing method using the molten metal quenching method has three main processes of a molten metal quenching process, a heat treatment process, and a nitriding process. Hereinafter, each process will be described sequentially.

<Melting quenching process>
In the molten metal quenching step, the raw material metal blended so as to have the composition described above is melted to obtain a molten metal, and then the molten metal is rapidly solidified. Specific solidification methods include a single roll method, a twin roll method, a centrifugal quenching method, a gas atomization method, and the like, but it is desirable to use a single roll method. In the single roll method, the molten alloy is discharged from a nozzle and collided with the peripheral surface of the cooling roll, whereby the molten alloy is rapidly cooled to obtain a ribbon-like or flaky quenched alloy. The single roll method has higher mass productivity and better reproducibility of the quenching conditions than other molten metal quenching methods.

  The rapidly solidified alloy varies depending on its composition and the peripheral speed of the cooling roll, but exhibits an amorphous single phase, a mixed phase of an amorphous phase and a crystalline phase, or a single crystalline phase. The amorphous phase is microcrystallized by a heat treatment performed later. As one measure, the higher the peripheral speed of the cooling roll, the higher the proportion occupied by amorphous.

When the peripheral speed of the cooling roll is increased, the obtained quenched alloy becomes thinner, so that a more homogeneous quenched alloy can be obtained. It is most desirable for the present invention to have a microcrystalline structure that is finally obtained in a state of being cooled and solidified, but this is not easy to achieve. On the other hand, after obtaining a single-phase structure of the amorphous phase, it is of course possible to microcrystallize by heat treatment, but the crystal grains based on the nuclei formed in advance grow abnormally to produce coarse crystal grains. There is a fear. Therefore, a desirable form for the present invention is to obtain a solidified structure in which the microcrystalline phase is rich and the balance is the amorphous phase.
For this purpose, the peripheral speed of the cooling roll is usually 10 to 100 m / s, preferably 15 to 75 m / s, and more preferably 25 to 75 m / s. When the peripheral speed of the cooling roll is less than 10 m / s, the crystal grains are coarsened and it is difficult to obtain a desired fine structure. When the peripheral speed of the cooling roll exceeds 100 m / s, the adhesion between the molten alloy and the peripheral surface of the cooling roll. Becomes worse and heat transfer is not effectively performed. In addition, the equipment cost increases. The molten metal quenching step is preferably performed in a non-oxidizing atmosphere such as Ar gas or N 2 gas.

<Heat treatment process>
The quenched alloy obtained by the molten metal quenching process is then subjected to a heat treatment. This heat treatment produces microcrystals having a particle size required by the present invention when the quenched alloy is an amorphous phase single phase. Further, when the quenched alloy is a mixed phase of an amorphous phase and a crystalline phase, the amorphous phase is microcrystallized, and in addition, the crystal grains are controlled to the grain size required in the present invention. Further, when the quenched alloy has a single phase structure of a crystal phase, the crystal grain is controlled to a grain size required in the present invention. Therefore, it is necessary to perform this heat treatment unless the fine structure required by the permanent magnet powder of the present invention is obtained in the quenched alloy state.
The treatment temperature in this heat treatment is 600 to 850 ° C, desirably 650 to 800 ° C, and more desirably 670 to 750 ° C. The treatment time depends on the treatment temperature, but is usually about 0.5 to 120 hours. This heat treatment is preferably performed in a non-oxidizing atmosphere such as Ar, He, or vacuum.

<Nitriding process>
After the heat treatment, the quenched alloy is subjected to nitriding treatment. For N which is an interstitial element, a raw material originally containing N can be used. However, after manufacturing a composition containing an element other than N, it is treated by nitriding in a gas or liquid containing N. It is desirable to intrude. N ₂ gas, N ₂ + H ₂ mixed gas, NH ₃ gas, or a mixed gas thereof can be used as the gas into which N can enter. Moreover, it is desirable to treat these gases as high-pressure gases in order to speed up the nitriding treatment.

The temperature of the nitriding treatment is 200 to 450 ° C., desirably 350 to 420 ° C., and the nitriding treatment time may be appropriately selected within the range of 0.2 to 200 hr. The same applies to the treatment for invading C (carbonization treatment), and a raw material originally containing C can be used. After a composition containing an element other than C is produced, heating is performed in a gas or liquid containing C. It can also be processed. Alternatively, C can be infiltrated by heat treatment with a solid containing C. Examples of gases that can infiltrate C include CH ₄ and C ₂ H ₆ . Carbon black can be used as the solid containing C. Also in carbonization by these, conditions can be appropriately set within the same temperature and processing time range as nitriding.

The above is the basic process for obtaining the permanent magnet powder of the present invention. The alloy obtained by the molten metal quenching method is pulverized at any stage before the heat treatment process, before the nitriding process or after the nitriding process. can do. This is because the alloy obtained by the molten metal quenching method is usually different from the size required for the permanent magnet powder for bonded magnets. The pulverization is performed in an inert gas such as Ar or N ₂ .
The average particle size of the permanent magnet powder is not particularly limited, but it is desirable that the particle size is such that there is as much as possible no significant difference in crystallinity in the same particle, and the particle size can be used as a permanent magnet powder. It is desirable. Specifically, when applied to a bonded magnet, the average particle size is usually desirably 10 μm or more, but in order to obtain sufficient oxidation resistance, the average particle size is desirably 30 μm or more, and more desirably. Is 50 μm or more, more preferably 70 μm or more. Moreover, it can be set as a high-density bond magnet by setting it as this average particle diameter. On the other hand, the upper limit of the average particle diameter is desirably 500 μm, and more desirably 250 μm. In addition, the average particle diameter here can be specified by the median diameter D50. D50 is the particle size when the mass is added from particles having a small diameter and the total mass becomes 50% of the total mass of all particles, that is, the cumulative frequency in the particle size distribution graph.

  The permanent magnet powder obtained as described above can be used for a bonded magnet. The bonded magnet is produced by binding particles constituting the permanent magnet powder with a binder. There are several types of bonded magnets depending on the manufacturing method. For example, there is a compression bond magnet using press molding and an injection bond magnet using injection molding. As the binder, it is desirable to use various resins, but it is also possible to use a metal binder as a metal bond magnet. The type of the resin binder is not particularly limited, and may be appropriately selected from various thermosetting resins such as epoxy resin and nylon and various thermoplastic resins according to the purpose. The type of metal binder is not particularly limited. There are no particular restrictions on the various conditions such as the binder content to the permanent magnet powder and the pressure during molding, and the conditions may be appropriately selected from the normal range. However, in order to prevent coarsening of crystal grains, it is preferable to avoid a method that requires high-temperature heat treatment.

Although the example which obtains a fine crystal structure using the molten metal quenching method was demonstrated above, this invention is not limited to this method. As another method, there is a method using mechanical grinding. This method has three main processes: a mechanical grinding process, a heat treatment process, and a nitriding process. Since the heat treatment step and the nitriding step are the same as the method using the molten metal quenching method described above, description thereof is omitted.
Mechanical grinding can change what was a crystal structure into an amorphous phase by continuously applying mechanical impact to alloy particles having a predetermined particle size. The mechanical impact can be imparted by using a ball mill, a shaker mill, or a vibration mill known as a pulverizer. By processing the alloy particles with these pulverizers, the particle structure can be made amorphous.

  The alloy particles can be produced according to a conventional method. For example, it can be obtained by preparing an ingot having a predetermined composition and then pulverizing the ingot. Or the ribbon or slice obtained by the molten metal quenching method can also be made into the object of mechanical grinding. In this case, it goes without saying that it is not necessary to apply to a ribbon or flake that has been in an amorphous state from the beginning.

  The alloy powder made amorphous by mechanical grinding can obtain the permanent magnet powder of the present invention by sequentially undergoing a heat treatment step and a nitriding treatment step. Moreover, the bonded magnet of this invention can be obtained using this permanent magnet powder.

  As a technique for obtaining a fine crystal structure, there is a heat treatment (HDR: Hydrogenation-Decomposition-Desorption-Recombination) that removes hydrogen after being kept at a high temperature in a hydrogen atmosphere. The present invention can also obtain a fine crystal structure using this HDDR. The permanent magnet powder of the present invention can be obtained by sequentially performing a heat treatment step and a nitriding step on the HDDR-treated powder. Moreover, the bonded magnet of this invention can be obtained using this permanent magnet powder.

Hereinafter, the present invention will be described based on specific examples.
The raw materials weighed to the following composition were dissolved in an Ar gas atmosphere and rapidly solidified. The rapid cooling conditions are as follows.
The obtained alloy was flaky with a thickness of 20 μm. These were subjected to heat treatment for 2 hours at 800 ° C. in an Ar gas atmosphere.
Furthermore, it grind | pulverized to the magnitude | size which passes a 75 micrometer sieve with a stamp mill, and nitrided the ground powder. The nitriding conditions are 400 ° C. × 64 hr, N ₂ flow (atmospheric pressure).

Composition: Nd ₁ Fe _9.15 Co _2.0 Ti _0.85 Si _0.2
・ Single roll method (roll material: Cu) ・ Nozzle hole diameter: φ1 mm ・ Gas pressure: 0.5 kg / cm ²
-Molten metal temperature: 1400 ° C-Roll peripheral speed (Vs): 15, 25, 50, 75 m / s

About the flakes (sample) after rapid solidification and the sample after heat processing, the phase structure was observed by XRD (X-Ray Diffractometer, X-ray diffractometer). The results are shown in FIGS. FIG. 1 shows the observation results for the sample after rapid solidification, and FIG. 2 shows the observation results for the sample after heat treatment.
As shown in FIG. 1, in the sample obtained at roll peripheral speed (Vs) of 15 m / s and 25 m / s, the peak of ThMn ₁₂ phase is observed, whereas the roll peripheral speed (Vs) is 50 m. In the sample obtained at / s and 75 m / s, the peak of the ThMn ₁₂ phase is not observed, and the diffraction line is unique to amorphous.
As shown in FIG. 2, it was confirmed that the ThMn ₁₂ phase occupies the main phase at any roll peripheral speed after the heat treatment.

  3 shows the structure of the sample after heat treatment obtained at a roll peripheral speed (Vs) of 25 m / s, and FIG. 4 shows the structure of the sample after heat treatment obtained at a roll peripheral speed (Vs) of 75 m / s. It is a figure which shows the result observed with (Transmission Electron Microscope, a transmission electron microscope). As shown in FIGS. 3 and 4, it was confirmed that an extremely fine crystal structure was exhibited after the heat treatment. However, the structure after heat treatment has the following differences depending on the roll peripheral speed (Vs). In the sample obtained at a roll peripheral speed (Vs) of 25 m / s, many crystals having a particle size of about 25 nm are observed, and the maximum particle size is about 50 nm. On the other hand, in the sample obtained at 75 m / s, many crystals having a particle size of about 10 nm are observed, and the maximum particle size is about 100 nm.

Next, the magnetic properties of the samples after rapid solidification, heat treatment and nitriding treatment were measured by VSM (applied magnetic field: 20 kOe). The results are shown in Table 7. The N content of the sample after nitriding is as follows.
Roll peripheral speed (Vs) = 25 m / s: 2.93 wt%
Roll peripheral speed (Vs) = 75 m / s: 2.79 wt%
As shown in Table 7, it was confirmed that by performing nitriding after the heat treatment, both the coercive force (Hcj) and the remanent magnetization (σr) are improved, and sufficient characteristics as a permanent magnet powder can be obtained. Table 7 also shows the measurement results of the magnetic properties of the powders according to the following comparative examples, but both the coercive force (Hcj) and the remanent magnetization (σr) remain low compared to the examples. ing.

  Comparative Example: The raw materials were weighed so as to have the same composition as in the Examples, dissolved by high-frequency melting, and then cast into a water-cooled Cu mold (alloy thickness 10 mm). This alloy was pulverized with a stamp mill in the same manner as in the example, and then heat-treated and nitrided in the same manner as in the example to obtain a powder.

Next, 3 wt% epoxy resin was mixed and stirred with the nitrided powder (with a roll peripheral speed (Vs) of 50 m / s), and a molding pressure of 6 ton / cm ² was obtained using a mold having a cylindrical cavity of φ10 mm. The molded body was cured at 150 ° C. for 4 hours to obtain a bonded magnet. The magnetic properties of the bonded magnet were measured with a BH tracer (applied magnetic field: 25 kOe). The results are as follows.
Br = 6700G, Hcj = 7980 Oe, (BH) max = 8.5 MGOe

  After producing a rapidly solidified alloy having the composition shown in Table 8, heat treatment and nitriding treatment were performed. The conditions for rapid solidification, heat treatment, and nitriding are as follows. Table 8 shows the results of measuring the magnetic properties after the nitriding treatment.

<Rapid solidification>
・ Single roll method (roll material: Cu) ・ Nozzle hole diameter: φ1 mm ・ Gas pressure: 0.5 kg / cm ²
Melting temperature: 1400 ° C. Roll peripheral speed (Vs): 50 m / s
<Heat treatment>
Hold at 800 ° C. for 2 hr in Ar gas atmosphere <nitriding>
Hold at 400 ° C for 64 hours in N ₂ gas flow (atmospheric pressure)

It is a figure which shows the result of having observed the phase structure by XRD about the flakes after rapid solidification. It is a figure which shows the result of having observed the phase structure by XRD about the sample after heat processing. It is a figure which shows the result of having observed the structure | tissue after heat-processing the flake obtained by roll peripheral speed (Vs) at 25 m / s by TEM. It is a figure which shows the result of having observed the structure | tissue after heat-treating the flakes obtained with the roll peripheral speed (Vs) of 75 m / s by TEM. It is a chart which shows the result of the X-ray diffraction of the sample obtained by the experiment example. It is a thermomagnetic curve of the sample obtained by the experiment example.

Claims (10)

Formula _{R (Fe 100-yw Co w} Ti y) x Si z A v ( provided that 50 mol% or more with R is one or more rare earth elements Nd, A is N and C 1 The molar ratio of the general formula is x = 10 to 12.5, y = (8.3-1.7 × z) to 12, z = 0.2 to 2, v = 0.1-3, w = 0-30, and a composition satisfying (Fe + Co + Ti + Si) / R> 12,
A permanent magnet powder comprising a set of particles having an average crystal grain size of 200 nm or less. The permanent magnet powder according to claim 1, wherein the particles have a phase having a ThMn ₁₂ type crystal structure as a main phase. The permanent magnet powder according to claim 1, wherein the particles are substantially composed of a single-phase structure of a phase having a ThMn ₁₂ type crystal structure.   The permanent magnet powder according to any one of claims 1 to 3, wherein Nd occupies 70 mol% or more of R. Formula _{R (Fe 100-yw Co w} Ti y) x Si z ( wherein, R is more than the 50 mol% with is one or more rare earth elements are Nd) consists, the molar ratio of the formula X = 10 to 12.5, y = (8.3-1.7 × z) to 12, z = 0.2 to 2, w = 0 to 30 and (Fe + Co + Ti + Si) / R> 12 A powder having a composition satisfying
Heat-treating the powder in an inert atmosphere at a temperature range of 650 to 850 ° C. for 0.5 to 120 hours,
A method for producing a permanent magnet powder, wherein the powder subjected to the heat treatment is subjected to nitriding treatment or carbonization treatment.   6. The method for producing a permanent magnet powder according to claim 5, wherein the powder subjected to the rapid solidification treatment has a structure of any of an amorphous phase, a mixed phase of an amorphous phase and a crystalline phase, or a crystalline phase. .   The method for producing permanent magnet powder according to claim 5 or 6, wherein the rapid solidification treatment is performed by a single roll method, and the peripheral speed of the roll used is 10 to 100 m / s.   The method for producing permanent magnet powder according to any one of claims 5 to 7, wherein the heat treatment crystallizes the amorphous phase or adjusts the particle size of the crystal particles constituting the crystal phase. Permanent magnet powder,
A bonded magnet comprising a resin phase in which the permanent magnet powder is dispersed,
Hard magnetic particles of crystalline constituting the permanent magnet powder of the general formula _{R (Fe 100-yw Co w} Ti y) x Si z A v ( where together R is one or more rare earth elements 50 mol% or more thereof is composed of Nd, A is one or two of N and C), and the molar ratio of the above general formula is x = 10 to 12.5, y = (8.3-1.7 × z) -12, z = 0.2-2, v = 0.1-3, w = 0-30, and satisfies the composition of (Fe + Co + Ti + Si) / R> 12.   The bonded magnet according to claim 9, wherein the hard magnetic particles have an average crystal grain size of 200 nm or less.

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