JP6614365B2 - Rare earth-transition metal ferromagnetic alloys - Google Patents

Description

  The present invention relates to a rare earth-transition metal ferromagnetic alloy, and more particularly to a rare earth-transition metal ferromagnetic alloy suitably used for permanent magnets.

In recent years, there has been a demand for the development of a magnet with a reduced content of rare earth elements. In this specification, the rare earth element refers to at least one element selected from the group consisting of scandium (Sc), yttrium (Y), and lanthanoid. Here, the lanthanoid is a general term for 15 elements from lanthanum to lutetium. As a ferromagnetic alloy in which the composition ratio of the rare earth element is relatively small, RT ₁₂ (R is at least one rare earth element, T is Fe, Co, or Ni) having a body-centered tetragonal ThMn ₁₂ type crystal structure. Are known. RT ₁₂ has high magnetization, but has a problem that its crystal structure is thermally unstable.

  Patent Document 1 discloses a rare earth permanent magnet in which part of Fe as a T element is partially replaced by Ti as a structure stabilizing element, and the thermal stability is improved in exchange for high magnetization. Yes.

In Patent Document 2, the R element of the RFe _12- based compound is partially substituted with a substitution element M1 such as Zr or Hf, so that the amount of substitution element M2 such as Ti that substitutes the transition metal element is reduced and saturated. A rare earth permanent magnet is disclosed in which the ThMn ₁₂ structure is stabilized while maintaining the magnetization.

Patent Document 3 discloses an R′—Fe—Co based ferromagnetic alloy in which Y or Gd is selected as part of the R element of RFe ₁₂ , and this R′—Fe—Co based ferromagnetic alloy is disclosed. It is described that the alloy exhibits a high magnetic property because it has a ThMn ₁₂ type crystal structure formed by a rapid quenching method.

JP-A 64-76703 JP-A-4-322406 Japanese Patent Laying-Open No. 2015-156436

  Single-crystal-like main phase particles used in anisotropic sintered magnets, which are often used as high-performance magnets, generally coarsen the crystal grains of the raw material alloy, pulverize the raw material alloy, and then sinter In view of the general processing temperature during this sintering step, the main phase compound is required to be stably present at least 900 ° C. or higher, preferably 1000 ° C. or higher.

  Although the rare earth permanent magnet described in Patent Document 1 has improved thermal stability due to elemental substitution of Fe by Ti, since the amount of Fe substitution by Ti is large, magnetization is reduced by that amount, and sufficient magnetic properties are obtained. I can't get it.

On the other hand, in the rare earth permanent magnet described in Patent Document 2, although the ThMn ₁₂ structure is stabilized by substituting the transition metal element with Ti or the like, the effect is not necessarily sufficient.

  The R′—Fe—Co based ferromagnetic alloy described in Patent Document 3 does not replace the Fe element with the structural stabilizing element M, and thus has high magnetization, large magnetic anisotropy, and high Curie temperature. However, since it is a non-equilibrium phase, the main phase compound may decompose in a densification process at a high temperature such as sintering.

  Therefore, an object of the present invention is to provide a rare earth-transition metal ferromagnetic alloy having improved magnetic properties and thermal stability.

  A preferred embodiment of the rare earth-transition metal ferromagnetic alloy according to the present invention has a composition represented by the following formula (1).

R1 in _{1-x R2 x (Fe 1} -y Co y) w-z Ti z ... (1) ( Formula (1), R1 is, Y or Y and Gd, R2 is, Sm, La, Ce, Nd and One or more selected from the group consisting of Pr, and at least Sm-containing rare earth element, x, y, z, and w are 0 <x <1.0, 0 ≦ y ≦ 0.4, and 11 ≦, respectively. (W ≦ 12.5, 1/3 ≦ z ≦ 0.7 and x ≦ 6z−2)

  According to the present invention, a rare earth-transition metal ferromagnetic alloy having improved magnetic properties and thermal stability can be realized.

It is a figure which shows the relationship between the phase ratio of the bcc- (Fe, Co, Ti) phase contained in a ribbon, and the heat processing temperature. A phase fraction Ω phases ferromagnetic compound having ThMn ₁₂ type crystal structure of the ribbon is a diagram showing the relationship between the Sm substitution amount x. It is a figure which shows the composition range map of the composition which shows a high temperature stable phase. It is a figure which shows the relationship between a magnetic anisotropic magnetic field and heat processing temperature. It is a figure which shows the Sm substitution amount x dependence of the lattice constant a, the lattice constant c, the unit cell volume, and the axial ratio c / a for every Co substitution amount y. It is a figure which shows Sm substitution amount x dependence of axial ratio c / a for every Ti substitution amount z. It is a figure which shows the relationship between Co substitution amount y and Curie temperature. It is a figure which shows the initial magnetization curve in the room temperature about the ferromagnetic alloy which concerns on Example 4. FIG. It is a figure of the polarization microscope image of the cross-sectional structure | tissue of the ferromagnetic alloy which concerns on Example 5. FIG.

[Composition of rare earth transition metal ferromagnetic alloys]
The rare earth-transition metal ferromagnetic alloy according to the embodiment is an R1-R2-Fe-Co-Ti ferromagnetic alloy having a composition represented by the following formula (1).

_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z ... (1)
In the above formula (1), “R1” includes at least Y, and may further include Gd. In the formula (1), “R2” has at least Sm, and may further contain at least one rare earth element selected from the group consisting of La, Ce, Nd, and Pr. In the above formula (1), x (R2 substitution amount x) indicating the ratio of R2 atoms to the total of R1 and R2, Co atomic ratio y (Co substitution amount y) to the total of Fe and Co, Fe Z (Ti content z) indicating the atomic ratio of the Ti content to the total amount of Co, Ti, and w indicating the atomic ratio of the total amount of Fe, Co, and Ti to the total amount of R1 and R2, respectively. <X <1.0, 0 ≦ y ≦ 0.4, 11 ≦ w ≦ 12.5, 1/3 ≦ z ≦ 0.7 and x ≦ 6z−2.

  As a result of intensive studies by the present inventors, the following points were found and the above-mentioned rare earth-transition metal ferromagnetic alloy (hereinafter simply referred to as a ferromagnetic alloy) was obtained.

That is, in the above formula (1), at least Y is adopted as R1, and Gd is used as necessary, so that the high-phase stability of the main phase is ensured without using a substitution element such as Zr or Hf. It was found that the amount z of Ti as a stabilizing element can be reduced to about z = 1/3. Furthermore, it has been found that even if a part of the rare earth element R1 constituting the ThMn ₁₂ type crystal structure is replaced with R2 such as Sm, the ThMn ₁₂ type crystal structure becomes a high temperature equilibrium phase even with a small Ti addition amount. Thereby, high thermal stability can be maintained, suppressing the fall of the magnetic characteristic by addition of Ti. For this reason, for example, the thermal stability of the main phase is improved when a bulk magnet is produced by a sintering method.

  In the above formula (1), at least Sm is adopted as R2, and at least one rare earth element selected from the group consisting of La, Ce, Nd, and Pr is used as necessary. The present inventors have found that magnetic anisotropy, which is important for increasing the coercive force of alloys, can be improved.

  As described above, Ti is added as a stabilizing element for replacing a part of Fe or Co to enhance thermal stability. There is a trade-off relationship between the amount z of Ti at which the main phase is stable at a high temperature and the upper limit of the amount x of R2, and the amount x of R2 and the amount z of Ti satisfy a predetermined relationship. The present inventors have found that the main phase can be a stable ferromagnetic alloy at high temperatures. Specifically, when the amount x of R2 and the amount z of Ti satisfy the relationship of x ≦ 6z−2, the main phase can exist stably at a high temperature.

  In order to allow the main phase to exist stably at a high temperature, the lower limit value of the Ti amount z is 1/3, and is further defined by x ≦ 6z−2 in relation to x.

When conventional rare earth magnets such as Nd—Fe—B magnets and Sm—Fe—N magnets are used as bonded magnets, the magnet powder filling rate is approximately 80% or less. That is, the rare earth element content per unit volume is reduced as compared with the sintered magnet. The residual magnetic flux density (B _r ) at this time is about 1.2 T at the maximum. Temporarily, when the rare earth-transition metal system according to the present embodiment is produced in an aspect like a sintered magnet that is sufficiently densified, the value of the residual magnetic flux density (B _r ) can be obtained only below that of a conventional bonded magnet. In addition, the merit in terms of both magnetic characteristics and rare earth reduction is reduced. Therefore, it is desired that a sintered magnet obtained from an alloy using a ferromagnetic compound with a reduced amount of rare earth can have a residual magnetic flux density (B _r ) of 1.2 T or more.

The residual magnetic flux density (B _r ) is mainly determined by the saturation magnetization (J _s ) of the main phase compound, the volume ratio, and the degree of orientation. For this reason, in order to increase the residual magnetic flux density (B _r ), it is effective as one means to orient and sinter powder composed of single-crystal-like main phase particles in a magnetic field. Considering the degree of orientation and the main phase ratio obtained, it is required to increase the saturation magnetization (J _s ) of the main phase compound. Specifically, in order to obtain a residual magnetic flux density (B _r ) of 1.2 T or higher, the saturation magnetization (J _s ) of the main phase compound is desirably 1.4 T or higher.

When the amount z of Ti exceeds 0.7, the saturation magnetization (J _s ) of the main phase is difficult to exceed 1.4T, and may not be suitable as a raw material alloy for obtaining a high-performance magnet. By setting the amount z of Ti to 0.7 or less, the saturation magnetization (J _s ) of the main phase may be obtained, for example, a value of 1.4 T or more, and the residual exceeding the conventional rare earth bonded magnet It can be suitably used as a raw material alloy for obtaining a permanent magnet having a magnetic flux density (Br). From the viewpoint of obtaining high magnetic properties while maintaining the stability of the main phase at a high temperature, a more preferable range of Ti amount z is 1/3 ≦ z ≦ 0.6, and a more preferable range of Ti amount z. Is 1/3 ≦ z <0.5.

  In the above formula (1), a part of Ti may be substituted with an element such as Mo, V, etc. within a range of 50 mol% or less. The structure stabilizing element M for substituting a part of Ti is not limited to Mo and V. For example, Si, Al, Cr, Mn, W, Re, Be, Nb, or the like may be used.

  x is 0 <x <1.0. From the viewpoint of increasing the magnetic anisotropy energy, it is preferable that the amount x of R2 is large. However, if the amount x of R2 is too large, the stability of the main phase at a high temperature decreases. Therefore, x is more preferably 0.5 ≦ x ≦ 0.8.

  Regarding x, the following knowledge is obtained from the relationship of x ≦ 6z−2 described above. That is, in the range of z <0.5, the range of x where the main phase is stabilized is defined as a range satisfying the relationship of x ≦ 6z−2, but in the range of z ≧ 0.5, the main phase is The range of x that stabilizes at a particular high temperature (eg, 1050 ° C.) does not depend on the value of z.

From the viewpoint of increasing the magnetic moment and improving the magnetization at practical temperatures accompanying the improvement of the Curie temperature and improving the magnetic anisotropy, it is preferable to perform partial substitution of Fe by Co. However, if the amount of substitution by Co is too large, the magnetization and magnetic anisotropy are lowered, which is not desirable. Specifically, the Co substitution amount y is desirably 0 ≦ y ≦ 0.4, and more desirably 0.1 ≦ y ≦ 0.3. By adopting at least Y as R1 in the above formula (1), a highly magnetically anisotropic ThMn ₁₂ type compound can be obtained even if the amount of substitution by Co in the raw material alloy is relatively small. At this time, in the obtained ThMn ₁₂ type compound, 80 mol% or more of Co occupies 8f or 8j sites.

w represents the total amount of Fe, Co, and Ti, and 11 ≦ w ≦ 12.5. If w is less than 11, formation of a different phase Th ₂ Zn ₁₇ type crystal structure or Th ₂ Ni ₁₇ type crystal structure (hereinafter referred to as 2-17 phase) is not preferable. On the other hand, if w is larger than 12.5, the formation of α- (Fe, Co, Ti) phase or (Fe, Co) ₂ Ti phase becomes remarkable, which is not preferable.

The rare earth-transition metal ferromagnetic alloy (R1-R2-Fe-Co-Ti ferromagnetic alloy) according to the embodiment obtained as described above has a ThMn ₁₂ type crystal structure as a main phase, R1 -R2-Fe-Co-Ti ferromagnetic compound is contained. In the ferromagnetic alloy according to the embodiment, such a phase of a compound having a ThMn ₁₂ type crystal structure (hereinafter also referred to as a ThMn ₁₂ type phase) typically exists stably at 1000 ° C. or higher. it can. That is, the ThMn ₁₂ type phase contained in the ferromagnetic alloy according to the embodiment is a high temperature equilibrium phase in a temperature range of 1000 ° C. or higher. Thus, the R1-R2-Fe-Co-Ti-based ferromagnetic compound has a temperature range that becomes an equilibrium phase at 1000 ° C. or higher. More specifically, the R1-R2-Fe-Co-Ti ferromagnetic compound has a temperature range that becomes an equilibrium phase at 1000 ° C. or higher and 1300 ° C. or lower. Therefore, the ferromagnetic alloy according to the embodiment has good thermal stability of the main phase, and can be suitably used for adopting a high-performance magnet manufacturing process such as a sintering method. Therefore, for example, in a sintered body, a ThMn ₁₂ type phase can be contained in a high ratio.

  In addition, the ferromagnetic alloy according to the embodiment can achieve both the above-described characteristics and high magnetic characteristics. For example, “having a temperature range that becomes an equilibrium phase at 1000 ° C. or more and 1300 ° C. or less” does not require the main phase to be an equilibrium phase in all temperature ranges of 1000 ° C. or more and 1300 ° C. or less. That is, the temperature range used as an equilibrium phase includes the aspect which becomes 1050 degreeC or more and 1200 degrees C or less, for example, and the aspect which becomes 950 to 1250 degreeC. These temperature ranges for the equilibrium phase may vary depending on the detailed alloy composition.

In general, the “ThMn _12- type crystal structure” is a tetragonal crystal. However, in the rare earth-transition metal ferromagnetic alloy according to the embodiment, the tetragonal crystal lattice is slightly distorted and the target is orthorhombic. Even when the periodicity of atoms in the crystal is slightly disturbed, it is regarded as “ThMn ₁₂ type crystal structure”.

[Method for producing R1-R2-Fe-Co-Ti-based ferromagnetic alloy]
Examples of a method for producing a rare earth-transition metal ferromagnetic alloy (R1-R2-Fe-Co-Ti ferromagnetic alloy) according to the embodiment include a die casting method, a centrifugal casting method, a strip casting method, and a liquid ultra rapid cooling method. It is possible to adopt known methods such as For example, α- (Fe, Co, Ti) phase (However, Co and Ti are not essential. Hereinafter, they may be described as “bcc-Fe phase” or “bcc- (Fe, Co, Ti) phase”). In particular, when the generation of a phase that is not preferable as a raw material alloy for a magnet is suppressed as much as possible, a strip casting method or a liquid ultra-quenching method having a relatively high cooling rate can be employed.

  In addition, as a result of earnest research by the inventors, as shown in the examples described later, by controlling the heat removal direction during cooling of the molten metal and solidifying the components contained in the molten metal unidirectionally, it is relatively less reliable. It has been found that even when the Ti amount is z, coarse crystal grains can be formed while suppressing the formation of the α-Fe phase.

  In place of the above-described method, a reduction diffusion method in which an oxide or metal of a constituent element is mixed with granular metal calcium and heated and reacted in an inert gas atmosphere may be used. The reduction diffusion method has an advantage that a uniform structure is formed, so that it is easy to generate particularly during the solidification process and it is difficult to generate an α- (Fe, Co, Ti) phase that causes a decrease in magnetic properties.

In the R1-R2-Fe-Co quaternary nonequilibrium phase containing no Ti as the main component, a continuous structural change from the TbCu _7- type crystal structure to the ThMn _12- type crystal structure occurs with heat treatment. In this non-equilibrium phase, thermal decomposition occurs during heat treatment, and a 2-17 phase or a bcc phase (bcc-Fe phase) is precipitated. Such a non-equilibrium phase is stable only mainly at 900 ° C. or less, depending on the composition, and most of the non-equilibrium phase is decomposed in a temperature range of 1000 ° C. or more typically applied in a sintering process or the like. To do.

The rare earth-transition metal-based ferromagnetic alloy according to the embodiment allows the ThMn _12- type phase to exist stably even in a temperature range of 1000 ° C. or higher by adding a small amount of Ti to the composition system of the non-equilibrium phase. Can do.

In the rare earth-transition metal ferromagnetic alloy according to the embodiment, the formation of the bcc-Fe phase can be suppressed relatively easily, and can be 3 wt% or less with respect to the entire alloy. On the other hand, for example, when the R element of the RFe ₁₂ alloy described in Patent Document 2 is replaced with Zr, the formation of the bcc-Fe phase becomes obvious, and the content becomes 4 wt% or more based on the entire alloy. This seems to be related to the formation of Zr (Fe, Co) ₂ phase.

  In addition, the rare earth-transition metal ferromagnetic alloy according to the embodiment can further reduce the heterogeneous phase generated in the solidification process by applying heat treatment, or can be a single crystal-like material useful as a raw material for anisotropic sintered magnets. In order to easily obtain a powder composed of particles by a pulverization method, the crystal grains may be coarsened. The heat treatment temperature at this time is preferably 900 ° C. or higher and 1250 ° C. or lower, and more preferably 1050 ° C. or higher and 1250 ° C. or lower. The heat treatment time is usually from 5 minutes to 48 hours.

  In addition, the alloy obtained by using the liquid ultra-quenching method represented by the single roll ultra-quenching method and the atomizing method may have fine crystal grains, in which case the raw material alloy for anisotropic sintered magnets Not suitable for. However, even in this case, it is suitably used as a magnet powder for an isotropic bonded magnet by performing an appropriate heat treatment. The heat treatment temperature at this time is preferably 900 ° C. or higher and 1000 ° C. or lower. The heat treatment time is usually 5 minutes or more and 4 hours or less.

Since the above-described ThMn ₁₂ type phase is often a high-temperature stable phase, it may be decomposed when the heat treatment temperature is less than 900 ° C. Therefore, care must be taken when setting the heat treatment temperature.

  Examples of the present invention will be specifically described below, but the present invention is not limited to these examples.

In Example 1, the influence of the Sm substitution amount x and the Ti substitution amount z on the phase stability was examined. In the examination of Example 1, in order to surely suppress the formation of the α- (Fe, Co, Ti) phase, a single roll ultra-quenching method, which is a kind of liquid ultra-quenching method, was applied as an alloy production method.
(Process A)
First, represented by a composition is the chemical formula _{Y 1-x Sm x (Fe} 0.83 Co 0.17) 12-z Ti z (x = 0.0~1.0, z = 0.0~1.0) Y (purity 99.9%), Sm (purity 99.9%), electrolytic iron (purity 99.9%), electrolytic cobalt (purity 99.9%), and Ti (purity 99.3%) 9%) was weighed respectively. In consideration of evaporation of Y and Sm at high temperature, weighed so that Y was 3 mass% and Sm was 5 mass% higher than the target composition. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting.
(Process B)
After confirming that the Y-Sm-Fe-Co-Ti-based alloy is sufficiently dissolved in the step (A), the molten metal is emitted onto a copper roll (roll diameter 230 mm) that rotates at high speed and rapidly solidified. A ribbon-shaped alloy (hereinafter referred to as a quenched ribbon) was produced. In Example 1, the roll peripheral speed was set to 40 m / s.
(Process C)
The quenched ribbon obtained in step B was wrapped in Nb foil and heat-treated at 800 ° C. to 1100 ° C. for 20 minutes in an Ar flow environment to obtain a rare earth-transition metal ferromagnetic alloy.
(Evaluation results)
The relationship between the phase ratio of the bcc- (Fe, Co, Ti) phase contained in the ribbon thus produced and the heat treatment temperature was evaluated. The evaluation results are shown in FIG. The phase ratio of the bcc- (Fe, Co, Ti) phase was identified using X-ray Rietveld analysis. Heat treatment temperature and phase of bcc- (Fe, Co, Ti) phase for ferromagnetic alloys of various compositions with x value of 0.4 and z value varied in the range of 0.0 to 0.5 The relationship with the ratio is shown in FIG.

Further, among the ferromagnetic alloys studied in FIG. 1, the phase ratio Ω and Sm substitution of the ThMn ₁₂ type phase contained in the ribbon are those for z = 0.3, z = 0.4, and z = 0.5. The relationship with quantity x was evaluated. The results are shown in FIG. The phase ratio Ω of the ThMn ₁₂ type phase was identified using X-ray Rietveld analysis. The graph of FIG. 2 shows ThMn ₁₂ when the Sm substitution amount x is changed in the range of 0.40 to 0.80 for each of z = 0.3, z = 0.4, and z = 0.5. The change in the phase ratio Ω of the mold phase is shown. In addition, each sample of the plot shown in FIG. 2 used what was heat-processed at 1050 degreeC.

  As shown in FIG. 1, in the compositions of z = 0.4 and 0.5, the bcc- (Fe, Co, Ti) phase (bcc-Fe phase) that is soft magnetic and significantly deteriorates the magnetic properties is 1000 ° C. The amount of the bcc- (Fe, Co, Ti) phase was reduced to 3 wt% or less, especially in the heat treatment at 1050 ° C. by the above heat treatment.

This result shows that the ferromagnetic compound having a ThMn ₁₂ type crystal structure present in the alloy having this composition (z = 0.4, 0.5) is stable even in a temperature range of 1000 ° C. or higher. Note that, at a heat treatment temperature of 1100 ° C., rare earth evaporation occurred in each alloy. Therefore, it is not essential that the phase ratio of the bcc- (Fe, Co, Ti) phase at the heat treatment temperature of 1100 ° C. in FIG. 1 is slightly higher than 1050 ° C.

As shown in FIG. 2, in the composition of z = 0.5, the phase ratio Ω of the ThMn ₁₂ type phase showed a high value of almost constant 90 wt% without depending on the Sm substitution amount x. In the composition of z = 0.4, the phase ratio Ω of the ThMn ₁₂ type phase exceeded 90 wt% at x = 0.4. When x ≦ 0.4, although not shown in FIG. 2, the phase ratio Ω of the ThMn ₁₂ type phase did not depend on the Sm substitution amount x, and was almost constant 90 wt%. On the other hand, when x ≧ 0.4, it was confirmed that the phase ratio Ω of the ThMn ₁₂ type phase decreased linearly (phase ratio Ω = 131−109x) as the Sm substitution amount x increased.

Similarly to the case of z = 0.4, even in the composition of z = 0.3, as the Sm substitution amount x increases, the phase ratio Ω of the ThMn ₁₂ type phase is linear (phase ratio Ω = 74-100x ) Was confirmed to decrease. In the composition of z = 0.3, the phase ratio Ω when extrapolated to x = 0 in FIG. 2 is 74 wt%, and any Sm substitution amount x reaches a level sufficient as a raw material alloy for high-performance magnets. It is estimated not to.

  In order to briefly show the above results, a composition range map of the composition showing the high temperature stable phase is shown in FIG. In FIG. 3, evaluation was made with a white circle mark indicating that the phase stability was confirmed, and a cross mark indicating that the phase was unstable and partially decomposed.

The phase stability was judged by whether or not the 2-17 phase, which is a different phase, precipitated in the alloy. However, since the diffraction pattern is similar between the 2-17 phase and the 1-12 phase which is a ThMn ₁₂ type phase, if either one is a trace amount, the presence or absence of the 2-17 phase can be determined only by X-ray diffraction. It is difficult to judge. Therefore, for the phase stability, the presence or absence of the 2-17 phase was evaluated using a thermomagnetic balance capable of detecting a trace phase in combination with the X-ray Rietveld analysis. The thermomagnetic balance used in this study is a magnetic field (10 to 15 mT) applied to a sample installation portion by attaching a permanent magnet to a thermogravimetric measuring device (thermobalance) (TG: TGA / SDTA851 ^e manufactured by METTTLER TOLEDO). The magnetic attractive force acting on the ferromagnetic phase in the sample can be detected as the weight value of TG. In addition to the composition of the plot of each graph examined in FIG. 2, in FIG. 3, 0.0 <x <0.4, 0.8 <x <1.0, 0.5 <z <0.6 For any composition in the range, an alloy was prepared by the same procedure as described above, and the phase stability was evaluated.

As shown in Example 2 and Example 3 to be described later, higher magnetic characteristics are obtained toward the upper left side in the region in the graph shown in FIG. Therefore, the composition of the region in the vicinity of the upper left line among the lines indicating the boundary of the region where the phase stability can be obtained is most desirable as the magnetic material. When this boundary was investigated in detail, it was found that the upper left boundary in FIG. 3 was given by the relational expression x = 6z−2. That is, when x and z satisfy the relationship of x ≦ 6z−2, it was confirmed that a ferromagnetic compound having a ThMn ₁₂ type crystal structure can exist stably.

  Considering FIG. 3, a specific preferable numerical range of x and z is examined. When z is a value satisfying 1/3 ≦ z ≦ 0.6 or 1/3 ≦ z <0.5. In particular, it includes a range in which the magnetic characteristics are suitable, and if x is a value satisfying 0.5 ≦ x ≦ 0.8, it is understood that the range in which the magnetic characteristics are stable is included.

In Example 2, the relationship between the Sm substitution amount x and the magnetic property value was examined.
(Process A)
First, the composition is Y _1-x Sm _x (Fe _0.83 Co _0.17 ) _11.5 Ti _{0.5 in the} chemical formula.
In order to obtain a raw material alloy represented by (x = 0.4 to 0.8), Y (purity 99.9%), Sm (purity 99.9%), electrolytic iron (purity 99.9%), and electrolytic cobalt (Purity 99.9%) and Ti (purity 99.9%) were weighed respectively. In consideration of evaporation of Y and Sm at a high temperature, the sample was weighed so that Y was 3 mass% and Sm was 5 mass% higher than the target composition. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting.
(Process B)
After confirming that the Y-Sm-Fe-Co-Ti-based alloy is sufficiently dissolved in the step (A), the molten metal is emitted onto a copper roll (roll diameter 230 mm) that rotates at high speed and rapidly solidified. A ribbon-shaped alloy having a uniform solidified structure was produced. In Example 2, the roll peripheral speed was set to 40 m / s.
(Process C)
The quenched ribbon obtained in step B was wrapped in Nb foil and heat-treated at 1050 ° C. for 20 minutes in an Ar flow environment.
(Evaluation results)
About each ferromagnetic alloy produced by the above, the magnetic anisotropy magnetic field at room temperature was measured. The magnetic anisotropic magnetic field was identified by a singular point detection method using a vibrating sample magnetometer (VSM) capable of applying a 10T magnetic field. Table 1 shows the measurement results of the magnetic anisotropic magnetic field for each ferromagnetic alloy.

As shown in Table 1, it was confirmed that the magnetic anisotropic magnetic field increased as the Sm substitution amount x increased. Therefore, from the viewpoint of increasing the magnetic anisotropy, it is desirable that the Sm substitution amount x is large. As shown in Table 1, when x = 0.5, the magnetic anisotropy magnetic field (6.7 μ ₀ H _a (T)) equivalent to Nd ₂ Fe ₁₄ B is shown, and therefore x ≧ 0.5 is preferable. .

  Next, for each ferromagnetic alloy shown in Table 1, each magnetic anisotropic magnetic field was measured for each alloy obtained by changing the heat treatment temperature in the range of 20 to 140 ° C. Identification of the magnetic anisotropy field was performed in the same manner as in Table 1 except that it was performed in dilute He to suppress oxidation of the sample.

  FIG. 4 shows the relationship between the magnetic anisotropic magnetic field and the heat treatment temperature. As shown in FIG. 4, in the composition having a large Sm substitution amount x, the temperature coefficient of magnetic anisotropy was poor, and the magnetic anisotropy magnetic field rapidly decreased as the temperature increased. This result is considered to be due to the fact that the temperature coefficient of the exchange interaction between the rare earth element and the Fe element is worse than the temperature coefficient of the exchange interaction between the Fe elements. Therefore, it is possible to design a rare earth alloy having a smaller temperature coefficient of magnetic anisotropy magnetic field as the concentration of non-magnetic Y element is higher.

However, for example, Nd ₂ Fe ₁₄ B has a magnetic anisotropy field of approximately 5 T in the vicinity of 140 ° C., which is a typical driving temperature of a practically important hybrid vehicle or electric vehicle motor. It can be confirmed that the ferromagnetic alloy exhibits a magnetic anisotropy magnetic field superior to this in the composition range of x ≧ 0.5. From the viewpoint of increasing the magnetic anisotropy magnetic field at a high temperature, x preferably satisfies x ≧ 0.5.

In Example 3, the relationship between the Ti substitution amount z and the magnetic property value was examined.
(Process A)
First, since the composition to obtain a _{Y 0.6 Sm 0.4 (Fe 0.83 Co} 0.17) at _{12-z Ti z (z =} 0.0~0.5) material alloy represented by the chemical formula, Y (Purity 99.9%), Sm (purity 99.9%), electrolytic iron (purity 99.9%), electrolytic cobalt (purity 99.9%), and Ti (purity 99.9%) were weighed, respectively. In consideration of evaporation of Y and Sm at a high temperature, the sample was weighed so that Y was 3 mass% and Sm was 5 mass% higher than the target composition. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting.
(Process B)
After confirming that the Y-Sm-Fe-Co-Ti-based alloy is sufficiently dissolved in the step (A), the molten metal is emitted onto a copper roll (roll diameter: 230 mm) rotating at high speed and rapidly solidified. A ribbon-shaped alloy having a uniform solidified structure was produced. In Example 3, the roll peripheral speed was set to 40 m / s.
(Process C)
The quenched ribbon obtained in step B was wrapped in Nb foil, and was subjected to heat treatment at 900 ° C. for 20 minutes, under which the decomposition of the ThMn ₁₂ type phase can be suppressed over the entire composition range in an Ar flow environment. This is because when heat treatment is performed at 1000 ° C., the ThMn ₁₂ type phase is decomposed in the sample having a composition range of z <1/3, and consistent evaluation based on the difference in Ti substitution amount z cannot be performed.
(Evaluation results)
About each obtained ferromagnetic alloy, the volume magnetization of the ThMn ₁₂ type phase at room temperature was measured. The measurement results of the volume magnetization of each ferromagnetic alloy are shown in Table 2 together with the Sm substitution amount x. However, the volume magnetization was identified from the magnitude of the internal magnetic field measured by ⁵⁷ Fe Mossbauer. At that time, 0.77 times the 15.7T / μ proportional, the magnetic moment of Fe 8f site magnetic moment of Co is selecting coordination _B between the magnetic moment inside the magnetic field of Fe, Ti is 8i A probable assumption was made that the site has selective coordination and the magnetic moment of Sm is 1.49 μ _B. As shown in Table 2, it was confirmed that the volume magnetization was remarkably lowered as the Ti addition amount increased. Therefore, it was shown that it is better to reduce the Ti substitution amount z as much as possible from the viewpoint of improving the volume magnetization. In particular, when z increased to 0.8, Js <1.4T.

In Example 4, the relationship between the Sm substitution amount x and the Co substitution amount y, the lattice change, and the magnetic property value were examined.
(Process A)
First, the composition is the chemical formula _{Y 1-x Sm x (Fe} 1-y Co y) 11.5 Ti 0.5 (x = 0.4~1.0, y = 0.17 (= 2/12), In order to obtain a raw material alloy represented by 0.25 (= 3/12), 0.33 (= 4/12)), Y (purity 99.9%), Sm (purity 99.9%) and electrolytic iron ( Purity 99.9%), electrolytic cobalt (purity 99.9%) and Ti (purity 99.9%) were weighed, respectively. In consideration of evaporation of Y and Sm at a high temperature, the sample was weighed so that Y was 3 mass% and Sm was 5 mass% higher than the target composition. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting.
(Process B)
After confirming that the Y-Sm-Fe-Co-Ti-based alloy is sufficiently dissolved in the step (A), the molten metal is emitted onto a copper roll (roll diameter: 230 mm) rotating at high speed and rapidly solidified. A ribbon-shaped alloy having a uniform solidified structure was produced. In Example 4, the roll peripheral speed was set to 40 m / s.
(Process C)
The quenched ribbon obtained in step B was wrapped in Nb foil and heat-treated at 1050 ° C. for 20 minutes in an Ar flow environment.

(Evaluation results)
FIG. 5 shows the dependency of the lattice constant a, the lattice constant c, and the axial ratio c / a on the Sm substitution amount for each Co substitution amount. From FIG. 5, it was confirmed that when the Sm substitution amount x increases, the lattice constants a and c increase, whereas when the Co substitution amount y increases, the lattice constants a and c decrease. The size of the crystal structure is related to the phase stability of the ThMn ₁₂ type crystal structure. The axial ratio c / a in the crystal structure is generally in the range of 0.56 to 0.57. In general, the axial ratio c / a tends to decrease as the M element substitution amount increases and the phase stability improves. is there.

FIG. 6 shows the Sm substitution amount dependence of the axial ratio c / a at y = 2/12 for each Ti substitution amount z. As shown in FIG. 6, z = 0.3, z = 0.4, z = 0.5, and the axial ratio c / a decreases as the value of z increases. Accordingly, it can be understood from FIG. 6 that the axial ratio c / a tends to decrease as the Ti substitution amount z increases and the phase stability improves. That is, it is determined that the smaller the axial ratio c / a, the higher the phase stability of the ThMn ₁₂ type phase.

  As discussed in FIGS. 5A, 5B, and 5D, the lattice constants a and c and the axial ratio c / a vary depending on the Co substitution amount y. Therefore, it is considered that Co substitution as well as Sm substitution contributes to the stability of the crystal structure. It is important from the viewpoint of phase stability to adjust both the Sm substitution amount x and the Co substitution amount y to be within an appropriate range. On the other hand, the axial ratio c / a indicates the degree of tetragonal crystal and is also related to the magnetic anisotropy. Specifically, although depending on the crystal field of the rare earth site, qualitatively, the magnetic anisotropy tends to be larger when the axial ratio c / a is smaller.

  Specifically, in the case of z = 0.5, the composition range of 0.6 ≦ x ≦ 0.8 (see FIG. 6) at y = 2/12 (see FIG. 6), and 0.6 ≦ x ≦ 0 at y = 3/12. .8 composition range (see FIG. 5 (d)), and y = 4/12, the composition range of 0.7 ≦ x ≦ 0.8 (see FIG. 5 (d)) is large and high in magnetic anisotropy, respectively. It can be said that the aspect ratio (axial ratio c / a) can provide phase stability. Further, when z = 0.4, 0.6 ≦ x ≦ 0.8 at y = 2/12 can be said to be an appropriate aspect ratio for obtaining large magnetic anisotropy and high phase stability ( (See FIG. 6).

  Further, in the three compositions examined in FIG. 5, the axial ratio c / a tends to be smaller as the Co substitution amount y is smaller (see FIG. 5D). For this reason, from the viewpoint of improving phase stability and magnetic anisotropy, y and z are preferably y = 2/12 and 0.6 ≦ x ≦ 0.8. The range of the desirable Sm substitution amount x described above does not change significantly in the range of 1/3 ≦ z ≦ 0.6, which is the range indicated as the appropriate Ti substitution amount z in Example 1.

  Next, for the three compositions examined in FIG. 5, the relationship between the Co substitution amount y and the Curie temperature was examined. The Curie temperature was measured using a thermomagnetic balance. Specifically, the temperature was lowered from 750 ° C., and the inflection point was defined as the Curie temperature. FIG. 7 shows the relationship between the Co substitution amount y and the Curie temperature when x = 0.7. From FIG. 7, it was confirmed that the Curie temperature increased with the increase of the Co substitution amount y. Therefore, from the viewpoint of increasing the Curie temperature, it is desirable that the Co substitution amount y is large.

  FIG. 8 shows the technical magnetization process (initial magnetization curve) at room temperature for samples of y = 2/12 and y = 3/12 with x = 0.7 composition. The technical magnetization process was evaluated using a vibrating sample magnetometer. As shown in FIG. 8, it was confirmed that the sample with y = 3/12 with a large amount of Co substitution was more easily magnetically saturated than the sample with y = 2/12 with a small Co substitution amount. Reflecting this point, if the magnetization in a magnetic field larger than 10T is estimated from the saturation asymptotic rule, it is presumed that the saturation magnetization of the sample of y = 2/12 is larger than that of the sample of y = 3/12. . It was confirmed that this was not dependent on the value of the Sm substitution amount x, and the same was true in the temperature range up to 140 ° C., which is important for practical use. This change in magnetization due to Co substitution is in accordance with the Slater-Polling curve, and reflects that when the Co substitution amount y is too large, the magnetization characteristics deteriorate.

  From the above points, in the case of the rare earth-transition metal-based ferromagnetic alloy according to the embodiment, if the Co substitution amount y is too large, not only magnetization but also magnetic anisotropy is reduced, which is not desirable. From the viewpoints of magnetization and magnetic anisotropy, it can be determined that the Co substitution amount y is appropriate in the vicinity of 2/12. Considering that it conforms to the Slater-Polling curve, specifically, it can be determined that the Co substitution amount y is preferably 0 ≦ y ≦ 0.4, and more preferably 0.1 ≦ y ≦ 0.3. The Sm substitution amount x at that time is preferably 0.6 ≦ x ≦ 0.8.

In Example 5, the coarsening of crystal grains was examined.
(Process A)
First, the composition is the chemical formula Sm _0.4 Y _0.6 (Fe _0.83 Co _0.17 ) _12-z T i _z
In order to obtain a raw material alloy represented by (z = 0.4), Y (purity 99.9%), Sm (purity 99.9%), electrolytic iron (purity 99.9%), and electrolytic cobalt (purity 99.99). 9%) and Ti (purity 99.9%) were weighed respectively. In consideration of evaporation of Y and Sm at a high temperature, the sample was weighed so that Y was 3 mass% and Sm was 5 mass% higher than the target composition. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting.
(Process B)
After confirming that the Y-Sm-Fe-Co-Ti-based alloy was sufficiently dissolved in the step (A), fine droplets having a maximum particle size of 100 μm or less were generated by a gas atomization method to form a Cu-based substrate. Solidified on top. Fine droplets were deposited on the Cu-based substrate by repeated rapid solidification.
(Process C)
The deposit thus obtained was removed from the substrate, and heat treatment was performed at 1050 ° C. for 4 hours in an Ar flow environment. FIG. 9 shows a cross-sectional view of the structure of the obtained sample using a polarizing microscope. However, in FIG. 9, the lower side is the base material side. It can be seen that the ThM _n12 type phase is unidirectionally solidified from the substrate side along the heat removal direction (vertical direction). Since the domain wall is substantially along the vertical direction, it can be seen that the c-axis direction, which is the easy axis of magnetization, is oriented in the vertical direction. The crystal grains are about 10 to 30 μm in the direction parallel to the base material, and it can be seen that 30 to 100 μm or more grows in the heat removal direction. From the above, it was shown that coarse crystal grains long in the c-axis direction can be formed by controlling the heat removal direction and solidifying in one direction. Such a shape long in the c-axis direction is advantageous from the viewpoint of shape magnetic anisotropy.

In Example 6, the rare earth-transition metal ferromagnetic alloy was analyzed by neutron diffraction.
(Process A)
First, in order to obtain a raw material alloy whose composition is represented by the chemical formula Y _0.2 Sm _0.8 (Fe _0.8 Co _0.2 ) _11.5 Ti _0.5 , Y (purity 99.9%) and Sm (Purity 99.9%), electrolytic iron (purity 99.9%), electrolytic cobalt (purity 99.9%), and Ti (purity 99.9%) were weighed, respectively. In consideration of evaporation of Y and Sm at high temperature, weighed so that Y was 3 mass% and Sm was 5 mass% higher than the target composition. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting.
(Process B)
After confirming that the Y-Sm-Fe-Co-Ti-based alloy is sufficiently dissolved in the step (A), the molten metal is emitted onto a copper roll (roll diameter 230 mm) that rotates at high speed and rapidly solidified. A ribbon-shaped alloy (hereinafter referred to as a quenched ribbon) was produced. In Example 6, the roll peripheral speed was set to 40 m / s.
(Process C)
The quenched ribbon obtained in step B was wrapped in Nb foil and heat-treated at 1050 ° C. for 20 minutes in an Ar flow environment to obtain a rare earth-transition metal ferromagnetic alloy.

  Next, the obtained rare earth-transition metal ferromagnetic alloy was pulverized to recover an alloy powder of 75 μm or less. A hollow vanadium cylinder (outer diameter: 5.0 mm) is placed in a vanadium container (inner diameter: 5.8 mm) so that Sm, which has a large neutron absorption correction, can be corrected. Transition metal based ferromagnetic alloy). After that, the vanadium container containing the powder was subjected to Ibaraki Prefecture material structure analysis installed at BL20 in the Materials and Life Science Experiment Facility (MLF) of the Japan Proton Accelerator Research Facility (J-PARC). The device was set in an apparatus (iMATERIA), irradiated with pulsed neutrons in a single frame mode (25 Hz) with an output of 300 kW at room temperature, and held until the accumulated count reached 50 MCounts or more, and neutron diffraction data was acquired.

The analysis of the acquired data was performed using the data of the back bank obtained by subtracting the background of the vanadium container, and the interplanar spacing d was in the range of 0.49 mm <d <2.57 mm. The analysis code used was Z-Riedverd 1.0.2. Since the magnetic structure is collinear with ferromagnetic coupling and the diffraction patterns of nuclear scattering and magnetic scattering match, the nuclear scattering intensity does not depend on d except for the Debye-Waller factor, and the magnetic scattering intensity is d ^−2. Is proportional to Using these features, after fitting the structural parameters roughly using the diffraction peak at 0.49 <d <1.4 on the low d side, 1.4 ≦ Å <d <on the high d side The diffraction pattern at 2.57 was taken in and fitting was performed using the magnetic moment as a variable. In the fitting, it was assumed that the magnetic moments of Y, Zr and Ti were zero, and Fe and Co were the same magnetic moment at the same site. In addition, it can be understood that the obtained diffraction pattern includes diffraction patterns from Y or Zr and a 3d transition metal grating and contains almost no information on magnetic scattering from Sm having a large absorption. Therefore, the magnetic moment of Sm was set to a literature value of 1.50 μ _B (g factor was 1.07), and refinement was not performed.

As a result of a series of studies, the bcc-Fe phase in the produced rare earth-transition metal ferromagnetic alloy was 1.9 wt%, and the rest were all ThMn ₁₂ type compounds. On the other hand, as a reference example, a Zr _0.2 Sm _0.8 (Fe _0.8 Co _0.2 ) _11.5 Ti _0.5 composition alloy produced by the same method was analyzed by the same method. The bcc-Fe phase showed a high value of 4.6 wt%.

Further, as a result of estimating the occupied sites of Co and Ti in the ThMn ₁₂ type compound in the rare earth-transition metal ferromagnetic alloy of this example, most of Co, for example, 80 mol% or more is at least one of the 8f site and the 8j site. It was confirmed that the occupation rates in the sites were 26% and 27% for 8f and 8j, respectively, and that most of Ti occupied the 8i site and the occupation rate in the sites was 20%.

  The present invention can be used particularly for permanent magnets and the like.

Claims (10)

A rare earth-transition metal ferromagnetic alloy having a composition represented by the following formula (1):
_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z ... (1)
(In the formula (1), R1 is Y or Y and Gd, and R2 is one or more selected from the group consisting of Sm, La, Ce, Nd and Pr, and includes at least Sm, x, y, z
, W are values satisfying 0 <x <1.0, 0 ≦ y ≦ 0.4, 11 ≦ w ≦ 12.5, 1/3 ≦ z ≦ 0.7 and x ≦ 6z−2, respectively. is there. )   2. The rare earth-transition metal-based ferromagnetic alloy according to claim 1, wherein in the formula (1), z is a value satisfying 1/3 ≦ z ≦ 0.6.   2. The rare earth-transition metal ferromagnetic alloy according to claim 1, wherein z is a value satisfying 1/3 ≦ z <0.5 in the formula (1).   2. The rare earth-transition metal ferromagnetic alloy according to claim 1, wherein in the formula (1), y is a value satisfying 0.1 ≦ y ≦ 0.3.   2. The rare earth-transition metal ferromagnetic alloy according to claim 1, wherein in the formula (1), x is a value satisfying 0.5 ≦ x ≦ 0.8. 2. The rare earth element according to claim 1, wherein the rare earth-transition metal ferromagnetic alloy contains an R1-R2-Fe—Co—Ti ferromagnetic compound having a ThMn ₁₂ type crystal structure as a main phase. Transition metal ferromagnetic alloy.   The rare earth-transition metal ferromagnetic alloy according to claim 6, wherein the R1-R2-Fe-Co-Ti ferromagnetic compound has a temperature range that becomes an equilibrium phase at 1000 ° C or higher.   The rare earth-transition metal ferromagnetic alloy according to claim 6, wherein the R1-R2-Fe-Co-Ti ferromagnetic compound has a temperature range that becomes an equilibrium phase at 1000C or higher and 1300C or lower.   The rare earth-transition metal ferromagnetic alloy according to any one of claims 1 to 8, wherein the amount of the bcc-Fe phase is 3 wt% or less. 10. The rare earth-transition metal ferromagnetic alloy according to claim 1, wherein 80 mol% or more of Co occupies 8j sites or 8f sites in the ThMn ₁₂ type ferromagnetic compound.

Images