JP2018103211A - Method for producing rare earth-transition metal based ferromagnetic alloy, and rare earth-transition metal based ferromagnetic alloy - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a method for producing a rare earth-transition metal based ferromagnetic alloy having a high 1-12 phase ratio, in which compositional variation in the crystal grains of the 1-12 phase or among the 1-12 phase crystal grains is reduced, and suitable for the raw material alloy of a sintered magnet, and to provide a rare earth-transition metal based ferromagnetic alloy.SOLUTION: Provided is a method for producing a rare earth-transition metal based ferromagnetic alloy comprising the steps of: preparing a molten metal of an alloy having a composition represented by (YSm)(FeCo)Ti(in the formula, 0.2≤x≤0.8, 0≤y≤0.3, 11.5≤w≤12.5, and 0.45≤z≤1.2); and feeding the molten metal of the alloy to rotating cooling rolls to obtain a slab continuously solidified to a thickness of 0.1 mm or more and 1 mm or less.SELECTED DRAWING: Figure 2

Description

  The present invention relates to a method for producing a rare earth-transition metal ferromagnetic alloy and a rare earth-transition metal ferromagnetic alloy.

  Permanent magnets are used in various motors such as automobile parts, industrial machines, and home appliances.

  A typical high performance magnet is an Nd—Fe—B magnet. Nd—Fe—B magnets are mainly used in drive motors for electric vehicles (EV, HV, PHV, etc.) and hybrid vehicles. The need for further high efficiency and miniaturization of motors has increased, and development of permanent magnets with higher magnetic properties is expected.

As one of the permanent magnet main phase alloys that exceed the magnetic properties of Nd-Fe-B magnets, attention has been focused on a ThMn ₁₂ type crystal structure and an RT ₁₂ compound having a similar structure. The RT ₁₂ compound is a compound constituting the main phase of the Nd—Fe—B magnet R ₂ T ₁₄ B (R is at least one rare earth element, and T is one or more iron group transition metal elements containing at least Fe) ) High magnetic properties are expected because it contains a higher concentration of iron group transition metals.

In the magnet disclosed in Patent Document 1, a part of the T element is partially partially formed by the structure stabilizing element M (Si, Al, Ti, V, Cr, Mn, Mo, W, Re, Be, Nb, etc.). By substituting, thermal stability is obtained in exchange for high magnetization. Moreover, in the example of Patent Document 2, Y-Sm-Fe- Co -based ferromagnetic compound _{Y 0.6 Sm 0.4 (Fe 0.83 Co} 0.17) 11.5 volume magnetization (saturation magnetization (J _s )) 1.68T, magnetic anisotropy magnetic field ( _HA ) 7.60T (about 6.0 MA / m) and high magnetic properties. Patent Document 2, Y-Sm-Fe- Co -based magnetic properties _{(J s} magnetic properties is _Nd 2 _{Fe 14} B of the ferromagnetic compound is 1.61T, _{H A} is 6.7T (about 5.3MA / m )) Can be surpassed.

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

The rare earth permanent magnet of Patent Document 1 has a problem that the magnetization becomes small because the substitution amount of the M element with respect to the T element is large. On the other hand, the ferromagnetic compound disclosed in Patent Document 2 is notable in that it exhibits high physical properties. However, as a result of studies by the present inventors, the RT ₁₂ compound phase (hereinafter referred to as 1-12 phase) It has been found that the problem is that the distribution is non-uniform. Specifically, there is a problem that composition variations in individual 1-12 phase crystal grains and composition variations between 1-12 phase crystal grains are large, and magnets having excellent magnetic properties cannot be stably supplied. I found out that

  Various embodiments of the present invention provide a raw material alloy for a sintered magnet having a high 1-12 phase ratio and a small composition variation within 1-12 phase grains or between 1-12 phase grains. A method for producing a rare earth-transition metal ferromagnetic alloy and a rare earth-transition metal ferromagnetic alloy suitable for use in the present invention are provided.

  After the content of the claim is confirmed, the description corresponding to each claim will be added.

  According to the present invention, a rare earth-transition metal-based strength suitable for a raw material alloy of a sintered magnet having a high 1-12 phase ratio and a small variation in composition within and between 1-12 phase grains. A magnetic alloy and a method for producing the same can be provided.

It is explanatory drawing of the strip casting method. Sample No. It is a figure which shows the photograph of the alloy cross section in the copper roll surface side of 4. FIG. Sample No. FIG. 4 is a view showing a photograph of an alloy cross section at the center of 4. Sample No. It is a figure which shows the photograph of the alloy cross section in the cool-release surface side of 4. FIG. Sample No. It is a figure which shows the photograph of the alloy cross section in the 10 copper hearth surface side. Sample No. It is a figure which shows the photograph of the alloy cross section in 10 center parts. Sample No. It is a figure which shows the photograph of the alloy cross section in the 10 cool-release surface side. Sample No. 4 and sample no. It is the figure which showed the value of z in 10 1-12 phases.

According to the study by the present inventors, in the Y—Sm—Fe—Co based ferromagnetic compound described in Patent Document 2, the distribution of the RT ₁₂ compound phase (1-12 phase) is non-uniform. Various causes and examinations were conducted on this cause. First, before describing an embodiment of the present disclosure, these examination results will be described below.

  The Y-Sm-Fe-Co ferromagnetic compound disclosed in Patent Document 2 is a single-roll ultra-rapid cooling method, that is, a ribbon-like shape in which a molten metal is injected onto a copper roll rotating at high speed and rapidly solidified. Alloys and those that have been heat treated. However, the ferromagnetic compound present in this alloy is a non-equilibrium phase, and it is difficult to produce the target ferromagnetic compound when the roll peripheral speed is low, that is, when the cooling rate is reduced.

As a method for enhancing the structural stability of the RT ₁₂ compound, as disclosed in Patent Document 1, it is conceivable to substitute a part of T with an element such as Ti or V. Therefore, the present inventors have prepared a method in which a compound obtained by substituting part of Fe and Co with Ti in order to improve the crystal structure stability of the Y-Sm-Fe-Co ferromagnetic compound of Patent Document 2 described above is a casting method. After that, heat treatment was performed for homogenizing the structure. However, in a sample having a small Ti content, in addition to the 1-12 phase, a phase having a Th ₂ Ni ₁₇ type structure (hereinafter referred to as 2-17 phase) or a bcc structure phase containing Fe, Co, and Ti (hereinafter referred to as “phase”). As a result, a lot of α- (described as Fe, Co, Ti) phase was present. The 2-17 phase and the α- (Fe, Co, Ti) phase are soft magnetic phases, and if they are present in the magnet, they cause a significant decrease in residual magnetic flux density (B _r ) and coercive force (H _cJ ). It is necessary to suppress generation as much as possible.

Furthermore, as a result of the experiments by the present inventors, a sample near z = 1.0 having a large Ti content was composed of only about 1-12 phase, but within one 1-12 phase crystal grain, Or it became clear that the dispersion | variation in content of Y and Sm became large between the crystal grains of 1-12 phase, and it was difficult to eliminate this composition dispersion | distribution completely by subsequent heat processing. Even in a sample having a small Ti content, the same variation was observed although the degree of variation was small. It is considered that the contents of Y and Sm greatly affect the magnetic properties of compounds such as _Js and _HA . In order to stably supply a magnet having excellent magnetic properties, it is necessary to suppress the variation in composition of the raw material alloy as much as possible.

  In order to obtain the Y—Sm—Fe—Co-based ferromagnetic compound described in Patent Document 2, a single-roll ultra-quenching method is employed to rapidly cool the molten alloy. It is possible to suppress the composition variation of Y and Sm. However, the crystal grain size of the 1-12 phase obtained by the ultra-quenching method is as fine as about several tens of nanometers, and cannot be employed as a raw material for anisotropic sintered magnets, for example. Conversely, for example, a cooling method with a low cooling rate such as a melt casting method makes it difficult to form a ferromagnetic compound that is a non-equilibrium phase, or the main phase compound decomposes during a high-temperature densification process such as sintering. There is a problem that it ends up.

As a result of the study, the present inventors have determined that (Y _1-x Sm _x ) (Fe _1-y Co _y ) _wz T _iz (where 0.2 ≦ x ≦ 0.8, 0 ≦ y ≦ 0.3). 11.5 ≦ w ≦ 12.5, 0.45 ≦ z ≦ 1.2, and x ≦ 6z−2) a step of preparing a molten alloy having a composition, and a rotating cooling roll Supplying the molten metal of the alloy to obtain a slab continuously solidified to a thickness of 0.1 mm or more and 1 mm or less, that is, a slab of 0.1 mm or more and 1 mm or less by a strip casting method. And the step of cooling at a cooling rate to be high, the 1-12 phase ratio is high, the composition variation within and between the 1-12 phase grains is small, and the 1-12 phase It was found that a ferromagnetic alloy with a grain size suitable for anisotropic sintered magnet production can be produced. It was.

  As a result of experiments by the present inventors, when the Ti content is low and the composition is in the vicinity of z = 0.5, the coarse 2-17 phase and α- A (Fe, Co, Ti) phase is generated. However, since the 1-12 phase is an equilibrium phase at a high temperature around 1100 ° C., for example, when heat treatment is performed at 900 ° C. to 1250 ° C. for 1 hour to 200 hours, the 2-17 phase and α- (Fe, The ratio of the Co, Ti) phase decreases and the ratio of the 1-12 phase increases. However, if there are coarse 2-17 phase and α- (Fe, Co, Ti) phase before heat treatment, the diffusion distance required for the reaction of the phase becomes very large. As a result, the reaction to the 1-12 phase cannot proceed sufficiently, and a lot of the 2-17 phase and α- (Fe, Co, Ti) phase remain even after the heat treatment in a realistic heat treatment time. Furthermore, the distribution ratios of Y and Sm are different between the 2-17 phase and the 1-12 phase, and more Y is distributed to the 2-17 phase. Therefore, the concentrations of Y and Sm are greatly different between the 1-12 phase generated during casting and the 1-12 phase newly generated by the subsequent heat treatment.

  An alloy having a high 1-12 phase ratio can be obtained with a composition having a large Ti content near z = 1.0 even when the cooling rate during casting is lower than the cooling rate in the strip casting method. However, the composition of the 1-12 phase varies and is divided into a region with a high Y content and a region with a high Sm content. This is because a 1-12 phase having a high Y content is generated as a primary crystal. Further, this composition variation is not eliminated even by long-time heat treatment.

The embodiment of the present disclosure is configured by setting the composition of the alloy whose composition formula is (Y _1-x Sm _x ) (Fe _1-y Co _y ) _wz T i _z and the cooling rate of the molten metal to an appropriate range. , Generation of coarse 2-17 phase and α- (Fe, Co, Ti) phase can be suppressed. In addition, it is possible to suppress the generation of the 2-17 phase that occurs when the molten alloy is gradually cooled, and to reduce the uneven distribution of Y and Sm between the 1-12 phases and the composition variation within the 1-12 phase. . Furthermore, when the obtained alloy is heat-treated, the 2-17 phase and the α- (Fe, Co, Ti) phase react with each other without diffusion over a long distance to produce a 1-12 phase. The -12 phase ratio is high, and the composition of the 1-12 phase is more uniform than that of an alloy obtained by a general casting method. The “ThMn _12- type crystal structure” that is the basic structure of the 1-12 phase (RT ₁₂ compound phase) is generally a tetragonal crystal, but in the rare earth-transition metal ferromagnetic alloy according to the embodiment, A case where the tetragonal crystal lattice is slightly distorted to have orthorhombic symmetry, or a case where the periodicity of atoms in the crystal is slightly disturbed is also regarded as a “ThMn _12- type crystal structure”.

[Reason for limiting composition, etc.]
(Ratio of Y and Sm)
The range of x (Sm substitution amount x) indicating the atomic ratio of Sm to the sum of Y and Sm is 0.2 ≦ x ≦ 0.8. When x is less than 0.2, not only is the effect of the present invention difficult to obtain, but also _Js and _HA are lowered, which is not preferable. Further, when x is larger than 0.8, not only is the effect of the present invention difficult to obtain, but particularly when the Ti content z is small, the high temperature stability of the 1-12 phase cannot be secured, which is not preferable.

(Fe to Co ratio)
The range of y (Co substitution amount y) indicating the ratio of the number of Co atoms to the total of Fe and Co is 0 ≦ y ≦ 0.3. In order to avoid the possibility that the Curie temperature is lowered, y is more preferably 0.1 or more. On the other hand, if y is larger than 0.3, J _s and _HA are lowered, which is not preferable.

(Total amount of Fe, Co, Ti)
The range of w indicating the atomic ratio of the total amount of Fe, Co, and Ti with respect to the total amount of Y and Sm is 11.5 ≦ w ≦ 12.5. If w is less than 11.5, formation of the 2-17 phase becomes significant, which 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.

(Ti content)
The range of z (Ti content z) indicating the atomic number ratio of the Ti content to the total amount of Fe, Co, and Ti is 0.45 ≦ z ≦ 1.2. Further, the lower limit value of z depends on x, and the relationship of x ≦ 6z−2 is established.

When z is less than 0.45, the 1-12 phase cannot exist stably when the strip casting method employed in this embodiment is employed. Further, if z is larger than 1.2, the (Fe, Co) ₂ Ti phase increases, which is not preferable. In order to obtain higher magnetic properties, particularly _Js , it is preferable that the amount of Ti is small. Specifically, 0.45 ≦ z ≦ 0.7 is more preferable. More preferably, 0.45 ≦ z ≦ 0.6. There is a trade-off relationship between the Sm substitution amount x and the Ti content z, and if the relationship x ≦ 6z−2 is satisfied by experiment, the main phase is stable at a high temperature (for example, 1000 ° C.). I found out that it can exist.

[Embodiment of Alloy Manufacturing Method]
(A) Step of melting raw material alloy Mixing a single metal or an alloy composed of two or more metals so that Y, Sm, Fe, Co and Ti have a predetermined composition, vacuum or inert Dissolves in gas. The melting method is not particularly limited as long as the raw material alloy is uniformly melted. For example, high-frequency melting can be used. The temperature of the molten metal only needs to be equal to or higher than the melting point of the alloy, and it is desirable that the molten metal is raised to a temperature having a margin beyond the melting point so that the molten metal does not solidify before reaching the cooling roll in step (B). Since the viscosity of the molten metal changes depending on the temperature of the molten metal, it is preferable to optimize the molten metal temperature according to the composition and the process. Suppression of composition deviation at the time of melting can be suppressed by shortening the heating time for melting or by adding a rare earth element metal lump. In particular, since Sm has a high vapor pressure and easily evaporates, it is effective to introduce the composition after other alloys are dissolved in order to avoid composition deviation.

(B) Step of solidifying molten raw material alloy In this embodiment, the molten Y-Sm-Fe-Co-Ti alloy produced in step (A) is rapidly cooled by a strip casting method. A schematic example of the strip casting method is shown in FIG. An alloy melt 2 melted in a crucible 1 made of a heat-resistant material such as ceramics is poured into a tundish 3. After the molten metal 2 is temporarily stored in the tundish 3, the molten metal 2 comes into contact with the peripheral surface of the cooling roll 5 that rotates at the end 4 of the tundish. Thereby, the molten metal 2 is pulled up to the cooling roll 5 and rapidly cooled by moving to a certain position 6 in contact with the peripheral surface of the cooling roll 5. The molten metal 2 supplied from the crucible 1 through the tundish 3 to the cooling roll 5 is continuously quenched and solidified by the cooling roll 5 and is continuously strip-shaped slab (rapidly solidified alloy) extending in the circumferential direction. become. The width of the slab is, for example, about 5 mm to 20 cm.

  The tundish 3 can be formed from a heat resistant material such as ceramics. The tundish 3 not only temporarily stores the molten alloy 2 but also rectifies the molten metal 2 flowing toward the cooling roll 5 and adjusts the temperature of the molten metal 2 when the molten metal 2 contacts the cooling roll 5. Also plays.

  The cooling roll 5 can be formed from a material having high thermal conductivity such as copper or chromium. The peripheral speed of the cooling roll 5 is typically 0.2 to 3.0 m / s, and the cooling speed is 10 ° C./second to 1000 ° C./second, preferably 10 ° C./second to 500 ° C./second.

  By this strip casting method, an alloy having a thickness of several hundred μm can be produced. The thickness of the alloy is closely related to the cooling rate of the molten alloy. In general, the thinner the alloy, the higher the cooling rate because heat removal in the thickness direction is performed in a shorter time. In the single roll ultra-quenching method, the thickness of the alloy ribbon is about several tens of μm, and in the melt casting method, the thickness of the alloy is several mm or more. In the strip casting method, an alloy can be produced at a cooling rate between the single roll ultra-quenching method and the melt casting method, and an alloy thickness intermediate between the two can be obtained. In the present embodiment, the conditions of the strip casting method are adjusted so that the alloy thickness is 0.1 mm or more and 1 mm or less. In order to obtain an alloy having a thickness of less than 0.1 mm, it is necessary to increase the peripheral speed of the cooling roll or extremely slow the supply rate of the molten metal, so that stable operation becomes difficult. Furthermore, if the cooling rate is too high, the crystal grains become too fine. If the crystal grains become fine, single crystal grains cannot be obtained by subsequent pulverization, and an anisotropic sintered magnet cannot be produced. Moreover, when the alloy thickness is thicker than 1 mm, the cooling rate on the surface (free surface or cooling surface) side that is not in contact with the peripheral surface of the cooling roll is particularly low. In a region where the cooling rate of the alloy is relatively low, generation of coarse 2-17 phase and α- (Fe, Co, Ti) phase and compositional variation of 1-12 phase are not preferable. The alloy thus produced is pulverized into flakes having a size of 1 mm to 10 mm as necessary.

(C) Step of heat-treating rapidly solidified alloy The rapidly solidified alloy produced in step (B) may be subjected to heat treatment at a temperature of 900 ° C. to 1250 ° C. for 10 minutes to 200 hours as necessary. good. By such heat treatment, the crystal grains of the 1-12 phase which is the main phase can be coarsened and the main phase ratio can be increased. The increase in the main phase ratio due to the heat treatment is particularly effective for an alloy having a composition of z ≦ 0.9. The heat treatment temperature is preferably controlled by the temperature of the rapidly solidified alloy itself or the ambient temperature in the vicinity of the rapidly solidified alloy.

  The 1-12 phase is often a high-temperature stable phase and decomposes into a 2-17 phase and an α- (Fe, Co, Ti) phase at a low temperature (less than 900 ° C.). For this reason, the heat processing below 900 degreeC is unpreferable. Further, if the heat treatment temperature exceeds 1250 ° C., the alloy is dissolved depending on the composition, which is not preferable.

  When the heat treatment time is less than 10 minutes, the 1-12 phase crystal grains do not grow sufficiently. In addition, the reaction in which the 2-17 phase and the α- (Fe, Co, Ti) phase react to produce a 1-12 phase does not proceed sufficiently, and a high 1-12 phase ratio cannot be obtained. When the heat treatment time exceeds 200 hours, not only a problem from the viewpoint of productivity but also a problem that Y and Sm evaporate from the alloy surface may occur. If Y or Sm evaporates, it is not preferable because the amount of a different phase such as an α- (Fe, Co, Ti) phase cannot be ignored in a region close to the surface of the alloy. The heat treatment time is more preferably 1 hour or more and 24 hours or less.

  The heat treatment is desirably performed in a vacuum or in an inert gas such as Ar. However, since the saturated vapor pressure of Sm is particularly high, the heat treatment is preferably performed in an inert gas such as Ar. Moreover, it is more preferable to enclose Sm powder so as to be heat treatment in Sm vapor to form a closed system, or to prepare a raw material alloy in which Sm is increased in advance in the step (A). An alloy with z> 0.9 does not necessarily require heat treatment since the 1-12 phase ratio is high when it is solidified by the strip cast method and the variation of Sm in the 1-12 phase is small. As described above, the same heat treatment as described above may be performed to coarsen the crystal grains of the 1-12 phase.

(D) Characteristics of the obtained alloy The rare earth-transition metal-based ferromagnetic alloy obtained by the manufacturing method in the above embodiment is (Y _1-x Sm _x ) (Fe _1-y Co _y ) _wz Ti _z (where 0.2 ≦ x ≦ 0.8, 0 ≦ y ≦ 0.3, 11.5 ≦ z ≦ 12.5, 0.45 ≦ z ≦ 1.2, and x ≦ 6z−2). The phase constituted by the compound having the composition represented and having the ThMn ₁₂ type structure is 85% by mass or more. Further, when the maximum value of the atomic fraction of Sm with respect to the total rare earth of the compound having the ThMn ₁₂ type structure is z _max and the minimum value is z _min , z _max / z _min ≦ 1.3, typically z _max / Z _min ≦ 1.2 is achieved, the main phase ratio is high, and the composition variation of Y and Sm in the 1-12 phase is also reduced. Typically, the crystal grain size of 1-12 phase (which can be obtained from the equivalent circle diameter of 1-12 phase crystal grains when a cross section parallel to the thickness direction of the alloy is observed) is 1 μm or more, preferably 5 μm or more. Such alloys are a common process of anisotropic sintered magnet, milling, sintered magnet expressing molded in a magnetic field, a high performance magnet obtained by performing sintering, a particularly high B _r material Suitable as an alloy.

  The embodiment of the present disclosure will be described in more detail by way of examples. However, the embodiment of the present disclosure is not limited to the following examples.

<Experimental example 1>
First, the total weight of 4 kg represented by the chemical formula (Y _0.6 Sm _0.4 ) (Fe _0.83 Co _0.17 ) ₁₁ Ti _z (z = 0.4, 0.5, 1.0). Y-metal (purity 99.9%), Sm metal (purity 99.9%), electrolytic iron (purity 99.9%), electrolytic Co (purity 99.9) %) And Ti metal (purity 99.9%) were weighed respectively. Each of the weighed metals was mixed and put into a silica crucible, and heated to 1500 ° C. by high frequency induction heating to dissolve the raw material. Thereafter, the temperature of the molten metal was lowered to 1450 ° C., temporarily stored in a tundish, and then poured onto a copper cooling roll rotating at a peripheral speed of 1.5 m / s to be cooled. The cooled alloy was crushed by a crusher installed at the lower part of the cooling roll. This step was performed using a strip casting apparatus having a configuration as shown in FIG.

  Using an outer micrometer, the thicknesses of 20 locations were randomly measured for the alloys having the respective compositions obtained in the above steps.

  About each alloy of each composition obtained at the said process, 150g was weighed and put into the container made from molybdenum. Thereafter, the alloy was inserted together with the above vessel into a tubular heat treatment furnace in an atmosphere of Ar at a flow rate of 2 L / min, and heat treatment was performed at 1100 ° C. for 20 hours. After the heat treatment, the heat treatment furnace was opened to cool the alloy. At this time, the average cooling rate from 1100 ° C. to 100 ° C. was 10 ° C./min or more.

A part of the alloy before and after heat treatment was embedded in a carbon-containing thermosetting resin so that the thickness direction of the alloy was parallel to the observation surface. After roughly polishing the alloy using abrasive paper, finish polishing was performed using diamond slurry in the order of particle sizes of 3 μm, 1 μm, and 0.25 μm. Regarding the polished sample, a reflected electron image at an acceleration voltage of 15 kV was observed with a desktop scanning electron microscope (JCM-6000 NeoScope ^™ ) manufactured by JEOL Ltd. In addition, point analysis using an energy dispersive X-ray analyzer (JED-2300) was performed on a plurality of points.

  Part of the alloy before and after heat treatment was pulverized in a mortar, and powder X-ray diffraction was performed using the pulverized powder that passed through a 75 μm mesh. The apparatus used was a Bragg-Brentano focused beam wide-angle X-ray diffractometer (X-ray diffractometer, XRD, D8 ADVANCED / TXS manufactured by Bruker Ax Co., Ltd.). The voltage applied to the Cu rotating cathode was 45 kV, the current was 360 mA, and the Kβ filter was Ni. The scanning axis was 2θ / θ linked operation, the interval was 0.04 °, the speed was 0.6 s / step, and the range of 20 ° ≦ 2θ ≦ 80 ° was investigated at room temperature. For X-ray Rietveld analysis, DIFFRACplus Professional TOPAS 4 (manufactured by Bruker AEX Co., Ltd.) was used.

Table 1 shows the results of Experimental Example 1. The minimum thickness and the maximum thickness are a minimum value and a maximum value when 20 thicknesses are randomly measured. The constituent phase ratio is based on the X-ray Rietveld analysis. Further, the point analysis of the 1-12 phase was performed by randomly selecting 10 points, z was calculated at each point, and the ratio between the maximum value and the minimum value of z was set to z _max / z _min .

  No. For alloys with low Ti content such as z <0.45, such as 1 and 2, the 1-12 phase ratio was as low as 25.0% by mass without heat treatment. In such an alloy, the 1-12 phase ratio was improved by heat treatment at 1100 ° C. for 20 hours, but the 1-12 phase ratio was a low level of 71.4% by mass, and the 2-17 phase was 24.2%. Mass% remained.

  No. For alloys with z = 0.5 such as 3 and 4, the 1-12 phase ratio is as low as 55.4% by mass without heat treatment, but the 1-12 phase ratio is reduced by heat treatment at 1100 ° C. for 20 hours. The level was as high as 95.6% by mass, and the 2-17 phase was not confirmed. Moreover, the variation of Sm content in the 1-12 phase was also small.

  2, FIG. 3, and FIG. 4 shows backscattered electron images of the copper roll surface side, the center portion, and the cooling surface side. In any field of view, an α- (Fe, Co, Ti) phase having a fine and low brightness was observed, but it was found that most of the others were composed of a 1-12 phase.

  No. Alloys with a large Ti content such as z> 0.9 such as 5 and 6 have a high 1-12 phase ratio regardless of the presence or absence of heat treatment, and there is also a variation in the Sm content in the 1-12 phase. The result was small.

<Experimental example 2>
First, a total weight of 600 g whose composition is represented by a chemical formula (Y _0.6 Sm _0.4 ) (Fe _0.83 Co _0.17 ) ₁₁ T i _z (z = 0.4, 0.5, 1.0) Y (purity 99.9%), Sm (purity 99.9%), electrolytic iron (purity 99.9%), electrolytic Co (purity 99.9%), and Ti (purity 99.9%). 9%) was weighed respectively. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting. Thereafter, the molten metal at 1400 ° C. was spread on a water-cooled copper hearth and solidified to obtain an alloy ingot.

  Using an outer micrometer, the thicknesses of 20 locations were randomly measured for the alloys having the respective compositions obtained in the above steps.

  About each alloy of each composition obtained at the said process, 150g was weighed and put into the container made from molybdenum. Thereafter, the alloy was inserted together with the above vessel into a tubular heat treatment furnace in an atmosphere of Ar at a flow rate of 2 L / min, and heat treatment was performed at 1100 ° C. for 20 hours. After the heat treatment, the heat treatment furnace was opened to cool the alloy. At this time, the average cooling rate from 1100 ° C. to 100 ° C. was 10 ° C./min or more.

  A portion of the alloy before and after heat treatment was cut with an outer peripheral blade cutter, and then embedded in a carbon-containing thermosetting resin so that the thickness direction of the alloy was parallel to the observation surface. And it grind | polished by the method similar to Experimental example 1, and the point analysis by an energy dispersive X-ray analyzer was performed by observation with a desktop scanning electron microscope.

  For some of the alloys before and after heat treatment, powder X-ray diffraction was performed in the same manner as in Experimental Example 1, and X-ray Rietveld analysis was performed on the obtained results.

  Table 2 shows the results of Experimental Example 2.

  No. 7 and 8 are alloys having a low Ti content such as z <0.45. In the case of no heat treatment, the 1-12 phase ratio was as low as 14.1% by mass, and even after the heat treatment, the 1-12 phase ratio was as low as 43.2% by mass. Moreover, the dispersion | variation in Sm content in 1-12 phase became large by heat processing.

  No. 9 and 10 are alloys of z = 0.5. In an experimental example in which an alloy having the same composition was manufactured by the strip cast method, a high level of 1-12 phase ratio was obtained after the heat treatment, but in an example manufactured by a method of developing on a copper hearth, 1 was also obtained after the heat treatment. The -12 phase ratio was a low level of 61.3% by mass. Moreover, the dispersion | variation in Sm content in 1-12 phase became large by heat processing.

  5, FIG. 6, and FIG. 10 shows backscattered electron images on the copper roll surface side, the center portion, and the cooling surface side. In particular, there are coarse α- (Fe, Co, Ti) phases (a phase having a low brightness) and 2-17 phases (a phase having a bright brightness) at the center and the cooling surface side, and the 1-12 phase ratio is low. I understood it.

  In FIG. 4 and Experimental Example 2 The value of z in 10 1-12 phases is shown. The points 01 to 10 in FIGS. 4 and the points 11 to 20 in FIGS. It corresponds to 10 analysis points. Sample No. produced by strip casting method Compared with sample No. 4, sample No. 4 produced by a method of developing on copper hearth. No. 10 resulted in large variations in z.

  No. 11 and 12 are alloys having a composition with a large amount of Ti such as z> 0.9. Regardless of heat treatment, the 1-12 phase ratio was as high as 98% by mass. However, the Sm variation was already large before the heat treatment, and the variation was not eliminated even after the heat treatment.

  The rare earth-transition metal ferromagnetic alloy and the manufacturing method thereof according to the present disclosure provide a rare earth-transition metal ferromagnetic alloy suitable for a raw material alloy of a sintered magnet.

1 crucible 2 molten metal 3 tundish 4 end of tundish 5 cooling roll 6 position

Claims (7)

_{(Y 1-x Sm x)} (Fe 1-y Co y) w-z Ti z ( wherein 0.2 ≦ x ≦ 0.8,0 ≦ y ≦ 0.3,11.5 ≦ w ≦ 12. 5, preparing a molten alloy having a composition represented by 0.45 ≦ z ≦ 1.2 and x ≦ 6z-2);
Supplying a molten metal of the alloy onto a rotating cooling roll, and continuously obtaining a slab solidified to a thickness of 0.1 mm or more and 1 mm or less;
A method for producing a rare earth-transition metal ferromagnetic alloy containing   The method for producing a rare earth-transition metal ferromagnetic alloy according to claim 1, wherein the composition further satisfies 0.45 ≦ z ≦ 0.7.   3. The rare earth-transition according to claim 1, wherein after the step of obtaining the slab, a step of heat-treating the slab at 900 ° C. to 1250 ° C. for 10 minutes to 200 hours is performed. A method for producing a metallic ferromagnetic alloy. _{(Y 1-x Sm x)} (Fe 1-y Co y) w-z Ti z ( wherein 0.2 ≦ x ≦ 0.8,0 ≦ y ≦ 0.3,11.5 ≦ w ≦ 12. 5, 0.45 ≦ z ≦ 1.2, and x ≦ 6z-2),
The phase composed of a compound having a ThMn ₁₂ type structure is 85% by mass or more of the total,
Rare earth-transition where z _max / z _min ≦ 1.3 when the maximum value of atomic fraction of Sm with respect to the total rare earth contained in the compound having a ThMn ₁₂ type structure is z _max and the minimum value is z _min Metallic ferromagnetic alloy.   The rare earth-transition metal ferromagnetic alloy according to claim 4, wherein the thickness is 0.1 mm or more and 1 mm or less. 6. The rare earth-transition metal ferromagnetic alloy according to claim 4, wherein a crystal grain size of a phase composed of the compound having the ThMn ₁₂ type structure is 1 μm or more.   The rare earth-transition metal ferromagnetic alloy according to claim 4, wherein the composition further satisfies 0.45 ≦ z ≦ 0.7.

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