JP2014157845A - Magnet material - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a magnet material more excellent in magnetic characteristic than a conventional rare earth iron fluorine magnet.SOLUTION: In a magnet material composed of alloys of a binary system of R-Fe (R is a 4f transition element or Y in a formula) where F is positioned at an intrusion position of a crystal lattice and a ternary system of R-Fe-T (R is as above, T is a 3d transition element excluding Fe, is Mo, Nb, or W in a formula [however, when T is W, R is a 4f transition element or Y other than W]), a part of an Fe site is partially replaced by at least one of Al, Si, and Ga.

Description

  The present invention relates to a magnet material.

As background art in this technical field, there is JP-A-8-316018 (Patent Document 1). This publication states that “R (R is one or more of rare earth elements, Sm ratio in R is 50 atomic% or more), T (Fe or Fe and Co), N and M (Zr or Zr A part of which is substituted with at least one of Ti, V, Cr, Nb, Hf, Ta, Mo, W, Al, C and P), and the R amount is 4 to 8 atomic%, N The amount is 10 to 20 atomic%, the amount of M is 2 to 10 atomic%, and the balance is substantially T. In the magnet, R, T and N are mainly composed of TbCu ₇ type, Th ₂ Zn ₁₇ type and There is a hard magnetic phase including one or more types of crystal phases of Th ₂ Ni ₁₇ type and a soft magnetic phase composed of a T phase having a bcc structure, and the soft crystal has an average crystal grain size of 5 to 60 nm. The proportion of the magnetic phase is 10 to 60% by volume.This configuration provides a high coercive force, a high squareness ratio, Energy product is realized. Has been described as ".

On the other hand, Non-Patent Document 1 describes that “the most suitable element to be introduced into Sm ₂ Fe ₁₇ from the first principle calculation is F element”. Further, Non-Patent Document 2 describes that “the Curie temperature rises by introducing F element into the Th ₂ Ni ₁₇ structure”.

JP-A-8-316018

P.Uebele, K. Hummler, and M. Fahnle, “Full-potential linear-muffin-tin-orbital calculations of the magnetic properties of rare-earth-transition-metal intermetallics. III. Gd2Fe17Z3 (Z = C, N, O , F) ”, Phys. Rev. B 53 (6) 3296-3303 (1996). J.D.Ardisson, A.I.C.Persiano, L.O.Ladeira and F.A.Batista, “Magnetic improvement of R2Fe17 compounds due to the addition of fluorine”, J.Mat.Sci.Lett.16 (1997) 1658-1661.

Sm ₂ Fe ₁₇ N ₃ , which exhibits high magnetic properties as a parent phase of a permanent magnet material, generally introduces N element into the Sm ₂ Fm ₁₇ structure by thermal diffusion using NH ₃ or N ₂ which are nitriding gases. It is. Since nitriding follows the diffusion rule, a device for uniformly performing nitrogenation in a short time with a trace amount of additive elements has been studied. For example, in Patent Document 1, by adding one or more of Zr, Ti, V, Cr, Nb, Hf, Ta, Mo, W, Al, C, and P, the time required for nitriding treatment is controlled by controlling the structure. It is shortened. However, these elements are added for fine structure control, and there is no description that limits substitution with the Fe elements constituting the lattice and substitution sites. There is no description about the thermal stability of the matrix lattice.

On the other hand, Non-Patent Document 1 proposes Sm ₂ Fe ₁₇ F ₃ as a system having a possibility of exhibiting higher magnetic properties than Sm ₂ Fe ₁₇ N ₃ . Experimentally, Non-Patent Document 2 observes an increase in Curie temperature of 40 ° C. by introducing F element into Sm ₂ Fe ₁₇ . However, the Curie temperature rise is small and cannot be used as a permanent magnet material.

  Therefore, an object of the present invention is to provide a magnet material that is superior in magnetic properties to conventional rare earth iron fluorine-based magnets.

  In order to solve the above-described problems, the present invention provides a binary system of R—Fe [wherein R is a 4f transition element or Y], or R—Fe—T, in which F is arranged at an intrusion position of a crystal lattice. [Wherein, R is as described above, and T is a 3d transition element excluding Fe, or Mo, Nb, or W (provided that when T is W, R is a 4f transition element other than W or Y In the magnet material made of a ternary alloy, a magnetic material in which part of the Fe site is partially replaced by at least one of Al, Si, and Ga is employed.

  According to the present invention, it is possible to provide a magnet material that has better magnetic properties than conventional rare earth iron fluorine-based magnets.

Temperature dependence of desorption gas (mass-to-charge ratio m / z = 19) in fluorinated heat-treated Sm ₂ Fe ₁₇ magnetic powder and Sm ₂ Fe ₁₆ Al magnetic powder.

As a method for fluorinating the parent phase, ammonium fluoride (NH ₄ F), acidic ammonium fluoride (NH ₄ F · HF), ammonium silicofluoride ((NH ₄ F) ₂ SiF ₆ ), ammonium borofluoride ( NH ₄ BF ₄ ) such as hydrofluoric acid generated by thermal decomposition and sublimation of hydrofluoric acid, a method using ammonia, and hydrogen fluoride (HF), carbon tetrafluoride (CF ₄ ), trifluoride There is a method using a gas of nitrogen fluoride (NF ₃ ), boron trifluoride (BF ₃ ), or sulfur hexafluoride (SF ₆ ). Since the F element has a small diffusion coefficient in the iron-based alloy and is difficult to diffuse into the magnetic powder, the reactivity can be improved by performing heat treatment simultaneously with an element having a larger diffusion coefficient in the iron-based alloy than the F element. In particular, the H element is desirable because it exhibits a significant improvement in reactivity over other elements due to reduction diffusion. In addition, element H is characterized by being released out of the system by heating at a relatively low temperature of 300 ° C. or higher.

However, when the elements used at the same time are H, B, C, and N, they have a diffusion coefficient larger than that of the F element and are arranged at the same intrusion position as the F element, so these elements are arranged at the intrusion position first. This makes it difficult to increase the amount of F element introduced and hinders improvement in magnetic properties. In particular, the H element must be removed to reduce the magnetic anisotropy. An ideal fluorination reaction is a reaction in which an element used for improving reactivity during fluorination is released out of the system and an F element is introduced instead. As a result of intensive studies, it was found that the Sm ₂ Fe ₁₇ F _x system decomposes at 320 ° C. Therefore, an ideal fluorination reaction is possible only in simultaneous use with hydrogen element, and the ideal heat treatment temperature is 300 ° C. or higher and 320 ° C. or lower. In order to introduce the F element in a shorter time, it is necessary to increase the diffusion rate of the F element by heat treatment at a high temperature. That is, it is necessary to improve the decomposition temperature of Sm ₂ Fe ₁₇ F _x . In addition, an increase in decomposition temperature means an increase in work / treatment temperature, and it is also advantageous when producing bulk bodies from magnetic powder, such as increasing the density with increasing hot forming temperature and increasing the choice of low melting point alloys. To work.

  Since the factors that determine the decomposition temperature of an interstitial rare earth iron alloy are generally not clear, an exhaustive element search is required for each intrusion system. However, we focused on the large change in crystal lattice constant around the magnetic phase transition temperature of interstitial rare earth iron alloys. It is a giant magnetostrictive system with a lattice that greatly expands with the occurrence of magnetism and that has a distortion of two orders of magnitude or more than the normal magnetovolume effect. In all magnetic systems, the crystal lattice dominates magnetism, but the intrusive rare earth iron alloy has led to the idea that magnetism also dominates the crystal lattice. We thought that the crystal lattice could be controlled by controlling the magnetism.

  In the interstitial rare earth iron alloy, it is the Fe site called dumbbell site that most affects the magnetism. In the Bethe-Slater curve, it is known that the exchange interaction between Fe elements of 0.245 nm or less is negative (antiferromagnetic), and the Fe elements arranged in dumbbell sites correspond to this. In this case, since it is a metal, the exchange interaction is described by the motion exchange interaction and the direct exchange interaction. The former is ferromagnetic between degenerate orthogonal orbits, and the latter is antiferromagnetic. Dumbbell sites are antiferromagnetic because the direct exchange interaction is larger than the motion exchange interaction. It was thought that the stability of the crystal lattice could be improved by increasing the shared state of conduction electrons by increasing the ferromagnetic motion exchange interaction. In other words, it was considered that the thermal stability of the crystal lattice can be improved by preferentially arranging a nonmagnetic element or an element that exhibits a ferromagnetic exchange interaction between dumbbellsites. Therefore, we selected Al, Si, and Ga as nonmagnetic elements preferentially placed on dumbbell sites, and studied fluorination treatment. By substituting the Fe element with these nonmagnetic elements, it should be possible to change the thermal stability of the crystal lattice through the magnetic change mainly of the exchange interaction.

  In addition, the dumbbell site in this invention refers to the Fe site which has Fe with the shortest distance with Fe of the adjacent Fe site among several Fe sites.

  In the above, the idea and idea of the present invention were described, but only dumbbell sites were specified. However, magnetism needs to be described as a strongly correlated many-body electron system considering the electronic state of the entire system, and the heat of the crystal lattice It should be noted that the stability greatly affects the electronic state of the substitution element.

  Therefore, the present invention is particularly effective in an iron-based alloy into which an F element is introduced, in which the distance between Fe elements is 0.245 nm or less, or the crystal structure has sites where the interaction between sites is antiferromagnetic. The substitution amount of the non-magnetic element is desirably as many as the number of dumbbell sites, but in practice, since the effect as a strongly correlated many-body electron system is also included, it is preferably performed within a range in which the magnetic characteristics are not deteriorated in actual measurement. The substitution amount is preferably such that the substitution amount of at least one of Al, Si, and Ga is 2/17 at% or more and 1/3 at% or less with respect to the total amount of Fe and T.

  Hereinafter, typical manufacturing methods and embodiments will be described.

<Manufacturing method 1> Preparation of parent phase alloy:
At least one of R, Fe, and Al, Ga, and Si is mixed and dissolved in a vacuum, an inert gas, or a reducing gas atmosphere to make the composition uniform (dissolution casting method). The obtained mother alloy is coarsely pulverized in an inert gas by various mills.

As an inexpensive method other than the above method, a reduction diffusion method in which an oxide or metal of a constituent element is mixed with granular calcium metal and heated and reacted in an inert gas atmosphere can be used. Sm ₂ Fe _{17 has} a peritectic temperature below the peritectic temperature, and by selecting the iron powder particle size, the Sm diffusion distance can be controlled to some extent, making it easy to produce a single-phase alloy with little residual α-Fe phase and dissolving it. The casting method does not require indispensable uniform heat treatment. Since the Sm ₂ Fe ₁₇ alloy is obtained directly as a powder, a coarse pulverization step is not necessary. Further, in order to obtain a raw material powder having a smaller particle size, a microcrystalline oxide can be obtained by precipitating a hydroxide from a sulfuric acid aqueous solution of Sm and Fe and firing in the air. Magnetic powder having a particle size of several μm that does not require fine pulverization can be produced.

  On the other hand, as magnetic powder for nanocomposite magnets, ultra-quenched ribbon powder having a grain size of several μm to several tens of μm composed of a polycrystal having a crystal grain size of several tens to several hundreds nm is required. The ribbon powder is obtained by a liquid superquenching method in which a molten mother alloy is jetted and cooled in an inert gas or reducing gas atmosphere such as argon gas on the surface of a rotating roll such as a single roll or a twin roll. A hydrogenation, decomposition, dehydrogenation, and recombination (HDR) method is also useful as a method for obtaining a fine alloy structure.

  A parent phase alloy can also be produced by a nanoparticle process or a thin film process. For example, the vapor phase method includes a thermal CVD method, a plasma CVD method, a molecular beam epitaxy method, a sputtering method, an EB vapor deposition method, a reactive vapor deposition method, a laser ablation method, a resistance heating vapor deposition method and the like. The liquid phase method includes coprecipitation method, microwave heating method, micelle method, reverse micelle method, hydrothermal synthesis method, sol-gel method and the like.

  The mother alloy or mother alloy powder obtained by the above method needs to be coarsely pulverized as necessary. For example, there is a method of mechanically pulverizing with a ball mill, brown mill, stamp mill or jet mill in an inert gas atmosphere or a reducing gas atmosphere. The HDDR method is also effective, and forming a fine structure by this method is effective in coercive force expression. Of course, the coercive force is also exhibited by pulverization after the treatment with hydrogen fluoride gas.

<Production method 2> Fluorination treatment:
As a method for introducing the F element, there are methods as described in Examples 1 to 5 below, all of which are solid-gas reactions by gas. In the solid-gas reaction process, in order to obtain a uniform composition, it is important to uniformly contact the surface of the powder with the gas. In particular, since a problem occurs when a large amount of powder is gas-treated, it is preferable to constantly flow and agitate the powder using a fluidized bed or the like. In addition, since the element intrusion reaction is diffusion-controlled, a powder having a uniform concentration distribution can be obtained more quickly as the particle size of the powder is smaller. However, if the particle size is too small, a stable flow becomes impossible, so there is a limit naturally, and a particle size of 10 μm to several 100 μm is appropriate.

  Because of the decomposition reaction of the parent phase, at least initially, the fluorination heat treatment must be carried out at a temperature lower than 320 ° C. However, based on the concept of idea and idea of the present invention, the decomposition temperature may increase with the increase in the amount of F element introduced, and is under intensive research. Although element F is extremely reactive, it forms a passive state and is characterized by a low diffusion rate in the iron-base alloy. By controlling the solid-gas reaction so as to cause slow fluorination, It is possible to suppress decomposition. Moreover, since it is diffusion-controlled, the reaction rate improvement by pressurization cannot be expected. However, it is conceivable to raise the reaction temperature somewhat by suppressing the decomposition by pressurization. Since microcracks enter by hydrogenation and assist in element entry, mixing with hydrogen gas is effective in terms of forming a fine structure. Since the diffusion rate of F element is significantly slower than other hydrogen and nitrogen elements, heat treatment in an inert gas atmosphere after fluorination is effective for the purpose of making the fluorine concentration distribution in the powder uniform. There are cases. As described above, the fluorination reaction is different from other element introduction reactions and requires a special device in its production method. According to the present invention, since the decomposition temperature of the F element introduction system is increased, at least the initial fluorination heat treatment temperature can be increased. large. Also in the magnetic characteristics, Al, Si, and Ga are accompanied by an increase in the Curie temperature of the parent phase, and therefore, it works preferentially by substituting within an appropriate range in which the magnetization and magnetic anisotropy of the parent phase do not decrease.

<Manufacturing method 3> Grinding step:
In order to develop the coercive force, the coarse powder obtained in the element intrusion step is preferably finely pulverized to a mean particle size of 2 to 3 μm using a jet mill or a ball mill. In principle, it is desirable to pulverize to a single domain critical particle size of about 0.3 μm. Since the particle size after pulverization is very fine, it may be necessary to inactivate the particle surface by various methods.

  The magnetic powder for nanocomposites and HDDR magnetic powder exhibit coercive force without being finely pulverized, but the particle size may be affected when molded as a magnet body, and pulverization may be necessary. In addition, since the particle size of the magnetic particles is related to the coercive force, adjusting the particle size distribution by fine pulverization and classification also leads to adjustment of the coercive force distribution, in the sense of increasing the performance and reliability of the magnetic particles. It is important.

<Embodiment 1> Magnet molded product As an embodiment of the present invention, a magnet molded product including a magnet material and a binder is exemplified.

[1] Binder There are inorganic binders and organic binders as binders for hardening the magnetic powder. For example, low melting point metals or resins can be used. The resin includes a thermosetting resin and a thermoplastic resin. EP (epoxy) resin can be used as the thermosetting resin, PA (polyamide, nylon) resin, PPS (polyphenylene sulfide) resin as thermoplastic resin, NBR (acrylonitrile butadiene rubber), CPE (chlorinated) as elastomer. Polyethylene) resin and EVA (ethylene vinyl acetate) resin can be used. Further, it can be hardened with an inorganic compound, and is a solution of SiO ₂ precursor, CH ₃ O— (Si (CH ₃ O) ₂ —O) _m —CH ₃ (m is 3 to 5, average is 4), water , Dehydrated methyl alcohol and dibutyltin dilaurate can be mixed, impregnated and solidified.

[2] Molding method When manufacturing a bonded magnet using isotropic magnetic powder, it is important how to increase the density with good productivity. Although it depends on the magnetization characteristics of the magnetic powder, any magnetization is possible in principle as long as a molded body is formed, and a necessary magnetization pattern can be realized.

  When manufacturing a bonded magnet using anisotropic magnetic powder, the problem is the orientation of the magnetic powder. It is important to orient the crystal axes of the particles in the desired direction during molding and to increase the density as in the case of an isotropic magnet.

  The magnetization reversal mechanism is roughly classified into a nucleation type and a pinning type. In the former case, the smaller the crystal grain size, the better the coercive force. In the latter case, the coercive force is determined by the shape and number of pinning sites. The magnet material of the present invention is expected to have both magnetization reversal mechanisms depending on the production method. It is considered that the coercive force is mainly determined by the nucleation type in the case of a crystal structure of several μm obtained by rapid cooling, and by the pinning type in the case of a crystal structure of several tens to several hundreds of nm obtained by liquid super rapid cooling. When the magnetic powder has shape anisotropy, the direction of the magnetic powder during molding is important for increasing the coercive force. In general, it is desirable to form a direction in which the demagnetizing factor is small (the permeance factor is large) as the magnetization direction.

  Since the rare earth iron fluorine system has low thermal stability, hot press molding is effective for increasing the density as a bulk body.

[3] Manufacturing process There are compression molding and injection molding in the usual bonded magnet manufacturing methods. In compression molding, since the density of magnetic powder can be increased, a high energy product can be realized. For example, a magnetic material for a bond magnet of a ferromagnetic fluorine compound as a raw material and an EP resin are kneaded together with an additive to form a compound, which is put into a mold and press-molded. Thereafter, after heat-curing, the excess powder is washed and dropped to coat the surface. Important factors during compound production are selection of powder particles, surface treatment of powder particles, selection of resin, selection of kneading conditions, and the like. It is possible to increase the density by optimizing the particle size distribution. In order to increase the density, it is effective to increase the slipping property between the magnetic particles by using a liquid resin. When using anisotropic magnetic powder, the orientation of the magnetic powder by applying a magnetic field is added. The degree of orientation varies depending on the type of resin used, and it is important that each magnetic powder can move freely overcoming the viscosity of the binder when a magnetic field is applied.

  Injection molding has a feature that a complicated shape can be formed without post-processing. As the binder, PA resin or PPS resin is used. For example, a ferromagnetic fluorine compound bond magnetic powder as a raw material and a PA resin or PPS resin are kneaded together with an additive with a kneader to form a compound in a pellet form. This compound is put into an injection molding machine, heated and melted in a cylinder, and then injected into a mold for molding. It is necessary to adjust the viscosity of the resin. When using anisotropic magnetic powder, the magnetic powder can be oriented by forming a magnetic circuit necessary for the mold. In order to produce a high-performance anisotropic injection-molded magnet, it is necessary to sufficiently orient the molten compound when it is injected into the mold.

<Embodiment 2> Rotating machine As an embodiment of the present invention, there is a rotating machine having a rotor containing a magnet material. For example, examples of the PC peripheral device include a spindle motor (for HDD, CD-ROM / DVD, and FDD) and a stepping motor (CD-ROM / DVD pickup, FDD head drive). Examples of the OA include a facsimile, a copy, a scanner, and a printer. Examples of the automobile include a fuel pump, an airbag sensor, an ABS sensor, an instrument, a position control motor, and an ignition device. There are also PC game machines equipped with HDDs and DVDs, and TV set boxes that are devices that download and use digital data from the Internet or cable TV. Examples of home appliances include mobile phones, digital cameras, video cameras, MP3 players, PDAs, stereos, and the like. Other examples include air conditioners, vacuum cleaners, and power tools.

  In addition, since the magnet material of the present invention has a large magnetovolume effect, a Villari effect is expected in which the strength of magnetization changes when a pressure is applied to the magnetic material. It belongs to the magnetostriction phenomenon in a broad sense and can be used industrially for sensors and actuators.

  EXAMPLES Hereinafter, although this invention is demonstrated in detail using an Example, the technical scope of this invention is not limited by this.

In the present embodiment, an example will be described in which a rare earth iron alloy containing at least one of Al, Si, and Ga is subjected to a fluorination heat treatment with F ₂ gas. In addition, from the viewpoint of a high-performance permanent magnet material, an example will be described that is implemented for, for example, the 2-17 system, the 3-29 system, and the 1-12 system.

  At the time of producing the alloy, a rare earth element having uniaxial magnetic anisotropy was selected by introducing F element. For example, Sm was selected as the representative of the cigar type orbit for the 2-17 series and the 3-29 series, and Nd was selected as the representative of the Anpan type orbit for the 1-12 series. The selection of these rare earth elements is determined by the relationship between the fine direction between the rare earth elements of the target system and the quantization axis direction. When the introduced element is a second periodic element, rare earth elements exhibiting uniaxial magnetic anisotropy are limited if the crystal structure is determined. Further, for example, Ti was selected as the structure stabilizing element. However, the present invention is not limited by the type of rare earth element or structure stabilizing element. Further, for example, Al was selected from Al, Si, and Ga disposed on the dumbbell site. The substitution amount was, for example, one element substitution in each composition formula.

The R (Sm or Nd) -Fe-Al main phase alloy takes into account the evaporation of rare earth elements and mixes Fe and Al with Sm or Nd more frequently than the stoichiometric ratio, in a vacuum or in an inert gas. Alternatively, it was dissolved in a reducing gas atmosphere to make the composition uniform. It was prepared by quenching after heat treatment for phase formation. Since Sm ₂ Fe ₁₆ Al, Sm ₃ Fe ₂₇ AlTi, and NdFe ₁₀ AlTi are precipitated by the peritectic reaction of Fe, it is difficult to avoid a small amount of α-Fe in the obtained alloy. The obtained ingot of the alloy was pulverized in an inert gas using, for example, a jet mill or a ball mill so that the average particle size was 10 μm or less. For comparison, Sm ₂ Fe ₁₇ , Sm ₃ Fe ₂₈ Ti, and NdFe ₁₁ Ti containing no Al were also produced by the same method.

For the purpose of investigating the thermal stability of the crystal system by introducing F element alone, a fluorination heat treatment with F ₂ gas was performed. F ₂ gas was generated by electrolysis of hydrofluoric acid and was confirmed to have a purity of 99.5% or more. The reaction vessel was made of Ni, and the fluorination treatment was performed slowly and sufficiently before the reaction test to passivate the inside of the vessel. If the reaction vessel is about 600 ° C. or lower, a nickel-base alloy can be used, and if it is about 200 ° C. or lower, an iron-base alloy can be used. The F ₂ gas can be diluted to an appropriate concentration with Ar gas (5N) in the mixed gas tank.

FIG. 1 shows the temperature dependence of desorbed gas (mass-to-charge ratio m / z = 19) in Sm ₂ Fe ₁₇ magnetic powder and Sm ₂ Fe ₁₆ Al magnetic powder heat-treated in F ₂ gas at 190 ° C. for 3 hours. However, the heating rate is 5 ° C./min. The mass to charge ratio m / z = 19 corresponds to the F element. Sm ₂ Fe ₁₇ magnetic powder begins to release F element at 320 ° C., and the F element release amount becomes constant at 380 ° C. On the other hand, it was observed that the Sm ₂ Fe ₁₆ Al magnetic powder in which the parent phase alloy was replaced with Al was shifted to a high temperature side by about 40 ° C. compared to the Sm ₂ Fe ₁₇ magnetic powder. This means that the F element release temperature, that is, the decomposition temperature is increased by about 40 ° C. by Al substitution. It was found that the thermal stability of the crystal lattice system into which the F element was introduced was improved by substituting a part of the Fe element with Al. Similarly, an increase in the decomposition temperature was also observed in the Si and Ga substitution system for Sm ₂ Fe ₁₇ . Similarly, in the 3-29 system and the 1-12 system, an increase in the decomposition temperature was observed due to the substitution of Al, Si, and Ga with Fe.

Below, the example of the fluorination heat processing which utilized this knowledge is demonstrated. The maximum heat treatment temperature was set to 310 ° C. for Sm ₂ Fe ₁₇ magnetic powder, 350 ° C. for Sm ₂ Fe ₁₆ Al magnetic powder, and the other heat treatment conditions were set to the same heat treatment conditions. Intended for slow fluorination, the F ₂ gas concentration in the mixed gas was adjusted to 5%. After introducing the mixed gas into the reaction vessel at 120 ° C. up to 50 kPa in total pressure, the temperature was raised to the respective maximum heat treatment temperature at a rate of 1 ° C./1 minute. After maintaining at the maximum heat treatment temperature for 3 hours, the mixed gas was exhausted outside the furnace and cooled to room temperature in an Ar atmosphere. The crystal lattice constant and the Curie temperature of each magnetic powder subjected to the fluorination heat treatment in this way were evaluated. As a result, it was observed that the crystal lattice constant was expanded while maintaining the symmetry of the system, and an increase in the Curie temperature was also observed. Compared with the non-substituted system, the Al-substituted system had a large crystal lattice constant expansion rate and a Curie temperature increase rate. This is presumably because the maximum heat treatment temperature could be raised, so that the amount of F element introduced at the same time could be increased. The diffusion of the F element into the crystal grains follows the diffusion rule, which means shortening the fluorination heat treatment time. Similar results were observed in the Sm ₂ Fe ₁₇ Si and Ga partial substitution systems, and also in the 3-29 and 1-12 Al, Si and Ga partial substitution systems.

In the present embodiment, an example will be described in which a rare earth iron alloy containing at least one of Al, Si, and Ga is subjected to a fluorination heat treatment with HF gas. Since HF gas undergoes a fluorination reaction more easily than F ₂ gas, the use of HF gas is extremely important in mass production in terms of shortening reaction time and uniform fluorination. In some cases, the combined use of HF gas and F ₂ gas can be expected to double the fluorination promoting effect.

  The parent phase alloy was produced by the same method as in Example 1, for example. However, the present invention is not limited by the type of rare earth element or structure stabilizing element.

The fluorination heat treatment was performed in the form of HF gas flow in a tubular furnace. As a method for generating HF gas, hydrogen fluoride generated by thermal decomposition and sublimation of inorganic salt of hydrofluoric acid such as NH ₄ F, NH ₄ F · HF, (NH ₄ F) ₂ SiF ₆ , NH ₄ BF _4, etc. , A method using hydrogen fluoride generated by a chemical reaction between calcium fluoride and sulfuric acid, and a method using hydrogen fluoride which is vaporized by heating hydrofluoric acid. For example, hydrogen fluoride generated by a controlled chemical reaction between calcium fluoride and sulfuric acid was used. The fluorination treatment apparatus was composed of a reaction furnace for controlling the hydrogen fluoride reaction and a tubular furnace for heat-treating the magnetic powder. Quartz could be used for the reaction furnace and the tubular furnace on the condition that the amount of HF gas generated was made small and disposable for each test. The entire apparatus was constructed in a draft with a scrubber, and an extra HF gas trapping mechanism was provided downstream of the tubular furnace. Each magnetic powder of Sm ₂ Fe ₁₆ Al, Sm ₃ Fe ₂₇ AlTi, and NdFe ₁₀ AlTi was spread thinly on a glassy carbon (GC) so as to have a large surface area and placed in a tubular furnace. After evacuating the inside of the tubular furnace, HF gas was induced from the reaction furnace into the tubular furnace with Ar gas (5N), and a fluorination heat treatment was performed. As a result of intensive studies, it has been found that the outer peripheral portion of the magnetic powder is mainly decomposed into rare earth fluoride, rare earth acid fluoride, iron fluoride, iron, etc. by the redox diffusion reaction of HF gas which is too strong. Therefore, the amount of HF gas generated needs to be appropriately selected depending on the particle size of the magnetic powder to be used, and heat treatment was performed to suppress decomposition by generally slowing the fluorination treatment.

The maximum heat treatment temperature was set to 310 ° C. for Sm ₂ Fe ₁₇ magnetic powder, 350 ° C. for Sm ₂ Fe ₁₆ Al magnetic powder, and the other heat treatment conditions were set to the same heat treatment conditions. After sufficiently evacuating the inside of the tubular furnace at room temperature, HF gas was introduced into the tubular furnace at an Ar flow rate of 200 ml / min. The temperature was raised to each maximum heat treatment temperature at 5 ° C./min, and held for 1 hour. Thereafter, the supply of HF gas was stopped and the furnace was cooled to room temperature in an Ar stream. The crystal lattice constant and the Curie temperature of each magnetic powder subjected to the fluorination heat treatment in this way were evaluated. As a result, it was observed that the crystal lattice constant was expanded while maintaining the symmetry of the system, and an increase in the Curie temperature was also observed. However, in the unsubstituted system, the release of H element was insufficient, and it was concluded that the F element introduction reaction was not uniform as judged from the lattice shift and peak half width in the X-ray diffraction pattern. On the other hand, in the Al substitution system, since the maximum heat treatment temperature could be raised, H element was released sufficiently, the F element introduction reaction proceeded uniformly, and only one phase showing the same symmetry as the parent phase was observed. It was done.

As a result, the heat treatment temperature for H element emission can be improved by improving the decomposition temperature of the F element introduction system by Al substitution. Thereby, it is possible to provide a highly fluorinated and uniform rare earth iron fluorine-based magnet material. We observed dramatic improvements in the basic physical properties of magnets such as Curie temperature, magnetization, and magnetic anisotropy due to high fluorination. Similar results were observed in the Sm ₂ Fe ₁₇ Si and Ga partial substitution systems, and also in the 3-29 and 1-12 Al, Si and Ga partial substitution systems.

In this embodiment, an example will be described in which a rare earth iron alloy containing at least one of Al, Si, and Ga and an H element is subjected to a fluorination heat treatment with F ₂ gas. In addition, from the viewpoint of a high-performance permanent magnet material, an example will be described that is implemented for, for example, the 2-17 system, the 3-29 system, and the 1-12 system.

  At the time of producing the alloy, a rare earth element having uniaxial magnetic anisotropy was selected by introducing F element. For example, Sm was selected as the representative of the cigar type orbit for the 2-17 series and the 3-29 series, and Nd was selected as the representative of the Anpan type orbit for the 1-12 series. The selection of these rare earth elements is determined by the relationship between the fine direction between the rare earth elements of the target system and the quantization axis direction. When the introduced element is a second periodic element, rare earth elements exhibiting uniaxial magnetic anisotropy are limited if the crystal structure is determined. Further, for example, Ti was selected as the structure stabilizing element. However, the present invention is not limited by the type of rare earth element or structure stabilizing element. Further, for example, Al was selected from Al, Si, and Ga disposed on the dumbbell site. The substitution amount was, for example, one element substitution in each composition formula.

The R (Sm or Nd) -Fe-Al main phase alloy takes into account the evaporation of rare earth elements and mixes Fe and Al with Sm or Nd more frequently than the stoichiometric ratio, in a vacuum or in an inert gas. Alternatively, it was dissolved in a reducing gas atmosphere to make the composition uniform. After the heat treatment for phase formation, room temperature Ar was introduced and rapidly cooled to 800 ° C. In order to perform the HDDR process, an appropriate time was maintained in the H ₂ gas flow at 800 ° C., and then vacuuming was performed for an appropriate time. Further, in order to introduce H element into the intrusion position of the parent phase alloy, H ₂ gas flow was carried out at 200 ° C. for an appropriate time. The obtained ingot of the alloy was pulverized in an inert gas using, for example, a jet mill or a ball mill so that the average particle size was 10 μm or less. For comparison, a system not containing Al was also produced in the same manner. However, the HDDR process is not necessarily required in this embodiment.

F ₂ gas was generated by electrolysis of hydrofluoric acid and was confirmed to have a purity of 99.5% or more. The reaction vessel was made of Ni, and the fluorination treatment was performed slowly and sufficiently before the reaction test to passivate the inside of the vessel. If the reaction vessel is about 600 ° C. or lower, a nickel-base alloy can be used, and if it is about 200 ° C. or lower, an iron-base alloy can be used. The F ₂ gas can be diluted to an appropriate concentration with Ar gas (5N) in the mixed gas tank. The maximum heat treatment temperature was set to 310 ° C. for Sm ₂ Fe ₁₇ magnetic powder, 350 ° C. for Sm ₂ Fe ₁₆ Al magnetic powder, and the other heat treatment conditions were set to the same heat treatment conditions. Intended for slow fluorination, the F ₂ gas concentration in the mixed gas was adjusted to 5%. After introducing the mixed gas into the reaction vessel at 120 ° C. up to 50 kPa in total pressure, the temperature was raised to the respective maximum heat treatment temperature at a rate of 1 ° C./1 minute. After maintaining at the maximum heat treatment temperature for 3 hours, the mixed gas was exhausted outside the furnace and cooled to room temperature in an Ar atmosphere. In the reaction at 300 ° C. or higher, it is assumed that H element released from the alloy reacts with F ₂ gas to become HF gas. Since HF gas also corrodes Ni reaction vessels, it cannot be used after an appropriate number of tests. As a result, it was observed that the crystal lattice constant was expanded while maintaining the symmetry of the system, and an increase in the Curie temperature was also observed. However, in the unsubstituted system, the release of H element was insufficient, and it was concluded that the F element introduction reaction was not uniform as judged from the lattice shift and peak half width in the X-ray diffraction pattern. On the other hand, in the Al substitution system, since the maximum heat treatment temperature could be raised, H element was released sufficiently, the F element introduction reaction proceeded uniformly, and only one phase showing the same symmetry as the parent phase was observed. It was done.

As a result, the heat treatment temperature for H element emission can be improved by improving the decomposition temperature of the F element introduction system by Al substitution. Thereby, it is possible to provide a highly fluorinated and uniform rare earth iron fluorine-based magnet material. We observed dramatic improvements in the basic physical properties of magnets such as Curie temperature, magnetization, and magnetic anisotropy due to high fluorination. Similar results were observed in the Sm ₂ Fe ₁₇ Si and Ga partial substitution systems, and also in the 3-29 and 1-12 Al, Si and Ga partial substitution systems.

  In this embodiment, an example in which a rare earth iron alloy containing at least one of Al, Si, and Ga is fluorinated, nitrided, and carbonized will be described. In addition, from the viewpoint of a high-performance permanent magnet material, an example will be described that is implemented for, for example, the 2-17 system, 3-29 system, and 1-12 system.

  At the time of producing the alloy, a rare earth element having uniaxial magnetic anisotropy was selected by introducing F element. For example, Sm was selected as the representative of the cigar type orbit for the 2-17 series and the 3-29 series, and Nd was selected as the representative of the Anpan type orbit for the 1-12 series. The selection of these rare earth elements is determined by the relationship between the fine direction between the rare earth elements of the target system and the quantization axis direction. When the introduced element is a second periodic element, rare earth elements exhibiting uniaxial magnetic anisotropy are limited if the crystal structure is determined. Further, for example, Ti was selected as the structure stabilizing element. However, the present invention is not limited by the type of rare earth element or structure stabilizing element. Further, for example, Al was selected from Al, Si, and Ga disposed on the dumbbell site. The substitution amount was, for example, one element substitution in each composition formula.

The R (Sm or Nd) -Fe-Al main phase alloy takes into account the evaporation of rare earth elements and mixes Fe and Al with Sm or Nd more frequently than the stoichiometric ratio, in a vacuum or in an inert gas. Alternatively, it was dissolved in a reducing gas atmosphere to make the composition uniform. It was prepared by quenching after heat treatment for phase formation. Since Sm ₂ Fe ₁₆ Al, Sm ₃ Fe ₂₇ AlTi, and NdFe ₁₀ AlTi are precipitated by the peritectic reaction of Fe, it is difficult to avoid a small amount of α-Fe in the obtained alloy. The obtained ingot of the alloy was pulverized in an inert gas using, for example, a jet mill or a ball mill so that the average particle size was 10 μm or less. For comparison, Sm ₂ Fe ₁₇ , Sm ₃ Fe ₂₈ Ti, and NdFe ₁₁ Ti containing no Al were also produced by the same method.

For nitriding and fluorination, for example, nitrogen trifluoride (NF ₃ , purity 99.99% or more) gas was used. Instead may be a mixed gas of N ₂ gas and F ₂ gas, also may be a mixed gas of NF ₃ gas, N ₂ gas and F ₂ gas. N ₂ gas and F ₂ gas do not react with each other in the temperature range of this embodiment. As the reaction apparatus, a Ni reaction vessel having a sufficient wall thickness was used. In addition, Ni-based alloys such as Inconel and Monel can be used. Ni was also used for the sample container. In advance, a sample container was placed in the reaction container, and NF ₃ gas and F ₂ gas were slowly introduced at a temperature higher than the working temperature, and a sufficient amount of time was taken for passivation treatment.

For example, in this example, the following procedure was used. The carbon introduction system has a crystal lattice decomposition temperature as high as about 700 ° C., and the fluorine introduction system exhibits the highest magnet characteristics at room temperature. When performing at least two of nitriding, fluorination and carbonization, it is desirable to perform nitriding first. Fluorination and carbonization constitute a stable fluoride and carbon film around the magnetic powder, respectively, and hinder the introduction of other elements. The priority of fluorination and carbonization depends on the case. Further, since the decomposition temperature of the fluorine-introduced crystal lattice is relatively low, it is preferable to perform carbonization in advance. However, nitriding, carbonizing, and fluorination change the thermal stability of the crystal lattice in a complicated manner depending on the amount of substitution elements, and therefore the processing order or processes such as simultaneous execution depend on the case. In the following, the temperature of the heat treatment is continuously changed for nitriding, fluorination, and carbonization, but details of the chemical reaction actually proceeding at that temperature are not clear. Is not limited. The essence of the present embodiment is that at least one of C element and N element is arranged at the intrusion position in the F element-introduced rare earth iron alloy containing at least one kind of Al, Si, and Ga. The working process temperature is increased when at least two of fluorination, carbonization, and nitridation of at least one substituted rare earth iron alloy of Si and Ga are performed. In particular, by applying the present technology, a high-performance permanent magnet material can be uniformly produced in a short time. A sample 5 g was thinly spread and placed in a Ni container (5 cm × 7 cm). The operation of evacuation (0.4 kPa) and Ar (purity 99.9995% or more) replacement was performed three times. After raising the temperature to 120 ° C. in dilute Ar, vacuuming and Ar replacement were performed again for the purpose of removing moisture. Thereafter, the temperature was raised to 450 ° C. for nitrogenation, and NF ₃ gas was introduced. The introduced NF ₃ gas is mixed with Ar in advance, and the concentration is about 10% in terms of ideal gas. The treated pressure was 90 kPa. NF ₃ gas was introduced and held for 1 hour, cooled to 300 ° C. for fluorination, and held for 4 hours. The selection of the processing temperature for nitriding and fluorination and the selection of each holding time govern the penetration rate occupancy and the element diffusion. When nitridation has progressed completely, the intrusion position of the fluorine element disappears and fluorination does not occur. Therefore, appropriate adjustment is necessary. After holding, sufficient exhaust treatment was performed, and the furnace was cooled in an Ar atmosphere and released to the atmosphere at 40 ° C. or lower.

Separately, an example of performing carbonization and fluorination will be described. In Sm ₂ Fe ₁₆ Al, it has been intensively studied and clarified that fluorine element is discharged out of the crystal lattice at a temperature higher than 360 ° C. and carbon element is taken into the crystal lattice from around 300 ° C. Therefore, carbonization and fluorination can be performed in a temperature range of 300 ° C. to 360 ° C. However, in the system in which the substitution amount is changed, the absorption and emission temperatures of the fluorine element and the carbon element are different, so it cannot be generally stated. There is no problem if it is a substance that can be carbonized or fluorinated, such as hydrocarbon gas and fluorine gas (regardless of solid, gas, liquid, etc.). However, since hydrocarbon gas and fluorine gas react violently, it is desirable not to carry out mixing of both gases for safety, and there are few advantages in terms of process and characteristics. It is more efficient and effective to use the fluorocarbon from the beginning. In this embodiment, for example, carbon tetrafluoride (CF ₄ ) gas is used. Among the fluorocarbons, fluoromethane (CH ₃ F), difluoromethane (CH ₂ F ₂ ), trifluoromethane (CHF ₃ ) and the like containing hydrogen are effective in improving reactivity and increasing fluorination. A sample 5 g was thinly spread and placed in a Ni container (5 cm × 7 cm). The operation of evacuation (0.4 kPa) and Ar (purity 99.9995% or more) replacement was performed three times. After raising the temperature to 120 ° C. in dilute Ar, vacuuming and Ar replacement were performed again for the purpose of removing moisture. Thereafter, the temperature was raised to 350 ° C. for fluorination and carbonization, and CF ₄ gas was introduced. The introduced CF ₄ gas is mixed with Ar in advance, and the concentration is about 10% in terms of ideal gas. The treated pressure was 90 kPa. CF ₄ gas was introduced and held for 4 hours, and sufficient exhaust treatment was performed. The furnace was cooled in an Ar atmosphere and released to the atmosphere at 40 ° C. or lower. The selection of the treatment temperature for carbonization and fluorination, and the selection of each holding time dominate the penetration rate and element diffusion. When nitridation has progressed completely, the intrusion position of the fluorine element disappears and fluorination does not occur. Therefore, appropriate adjustment is necessary. The above results were separately performed for Sm ₃ Fe ₂₇ TiAl and NdFe ₁₀ TiAl in the same manner as Sm ₂ Fe ₁₆ Al, and the same results were obtained.

  In the present embodiment, an example relating to the element substitution amount of an F element-introduced rare earth iron alloy containing at least one of nonmagnetic elements Al, Si, and Ga will be described. In addition, from the viewpoint of a high-performance permanent magnet material, an example will be described that is implemented for, for example, the 2-17 system, 3-29 system, and 1-12 system.

  The substitution amount of at least one of Al, Si, and Ga and the Fe element changes not only the thermal stability of the system but also the magnetic property value of the parent phase. Therefore, the substitution amount must be determined according to the purpose. Don't be. The thermal stability of the system stabilizes as the substitution amount increases. However, the Curie temperature reaches the maximum value for the 2-17 system at about 2 to 3 element substitution. This is because these nonmagnetic elements are preferentially arranged on the dumbbell sites, so that the antiferromagnetic exchange interaction is weakened and the ferromagnetic interaction is strengthened. It is for thinning. In addition, the saturation magnetization at room temperature becomes maximum at about one element substitution to about two element substitution. Since this is determined by the balance between the dilution of the magnetic element and the increase in the Curie temperature, it varies depending on the type of the nonmagnetic element. Therefore, the substitution amount must be determined according to the amount of F element introduced and the amount of carbon element and nitrogen element introduced. Even if the magnetism becomes weak due to oversubstitution, the saturation magnetization at the Curie temperature and room temperature can be increased by substituting a part of Fe element with Co element. The magnetic anisotropy changes from in-plane to uniaxial according to the amount of nonmagnetic element substitution and Co substitution.

The above means that it is necessary to maximize and optimize each magnetic property value with the element substitution amount (nonmagnetic element and Co element), reactivity due to thermal stability, and the amount of F element introduced as main parameters. To do. These main parameters are correlated with each other, and the correlation is extremely complex. However, performing nonmagnetic element substitution more than the number of dumbbellsites has a large negative effect on magnetism, and substitution less than the number of dumbbellsites is desirable. The nonmagnetic element substitution amount is desirably 2 element substitution or less in the 2-17 system, 3 element substitution or less in the 3-29 system, and 4 element substitution or less in the 1-12 system. For a magnet material characterized in that the substitution amount of at least one of Al, Si and Ga is 2/17 at% or more and 1/3 at% or less with respect to the total amount of Fe and T, for example, in this embodiment, An example of introducing an F element into Sm ₂ Fe ₁₄ Co ₂ Al and Sm ₂ Fe ₁₄ Co ₂ Si will be described. Substitution and effects associated with the 3-29 system and the 1-12 system also occur, and the same applies to a crystal system configured by changing the three-dimensional stacking method of the CaCu ₅ structure. However, quantitativeness is different. In addition, the nonmagnetic element substitution amount dependence on the thermal stability of the matrix crystal lattice does not depend much on the type of rare earth element, but the magnetism changes depending on the type of rare earth element, so that Absolute temperature varies. The fluorination treatment and the nitridation and carbonization treatment used in combination with the fluorination can also be carried out by the methods described in Examples 1 to 4.

  The R (Sm or Nd) -Fe-Al-Co main phase alloy takes into account the evaporation of rare earth elements and mixes Fe, Al, Co with Sm or Nd more frequently than the stoichiometric ratio, and the vacuum or It was dissolved in an inert gas or a reducing gas atmosphere to make the composition uniform. It was prepared by quenching after heat treatment for phase formation. The obtained ingot of the alloy was pulverized in an inert gas using, for example, a jet mill or a ball mill so that the average particle size was 10 μm or less.

As the fluorination method, for example, a fluorination heat treatment with F ₂ gas was performed. F ₂ gas was generated by electrolysis of hydrofluoric acid and was confirmed to have a purity of 99.5% or more. The reaction vessel was made of Ni, and the fluorination treatment was performed slowly and sufficiently before the reaction test to passivate the inside of the vessel. If the reaction vessel is about 600 ° C. or lower, a nickel-base alloy can be used, and if it is about 200 ° C. or lower, an iron-base alloy can be used. The F ₂ gas can be diluted to an appropriate concentration with Ar gas (5N) in the mixed gas tank.

A sample 5 g was thinly spread and placed in a Ni container (5 cm × 7 cm). The operation of evacuation (0.4 kPa) and Ar (purity 99.9995% or more) replacement was performed three times. After raising the temperature to 120 ° C. in dilute Ar, vacuuming and Ar replacement were performed again for the purpose of removing moisture. Thereafter, F ₂ gas was introduced to 40 kPa over 2 hours. The introduced F ₂ gas is mixed with Ar in advance and the concentration is about 5% in terms of ideal gas. After heating to 400 ° C. at a rate of 1 ° C./1 minute and holding for 3 hours, sufficient exhaust treatment was performed, furnace cooling was performed in an Ar atmosphere, and the atmosphere was released to 40 ° C. or less. Since the thermal stability was improved by substitution of Co and Al, the fluorination heat treatment temperature could be increased, and the fluorination could be performed in a short time. In addition, high fluorination may be possible for the first time by increasing the heat treatment temperature. It is added that the selection of the heat treatment temperature and the selection of each holding time dominate the occupation ratio of the intrusion position and the diffusion of elements.

The magnetic characteristics of the magnetic powder thus produced are Sm ₂ Fe ₁₄ Co ₂ AlF _x magnetic powder with a saturation magnetization of 159 emu / g or higher, a Curie temperature of 579 K or higher, and Sm ₂ Fe ₁₄ Co ₂ SiF _x magnetic powder with a saturated magnetization of It was observed that it was 142 emu / g or more and the Curie temperature was 641 K or more. The magnetic anisotropy showed uniaxial magnetic anisotropy at room temperature, and the occurrence of magnetic hysteresis was confirmed by pulverization.

  Since the magnetic powders produced by the methods described in Examples 1 to 5 do not have chemical thermal stability in the sintering temperature range, it is not possible to increase the density of the bulk body by the sintering process. Therefore, it is necessary to increase the density by pressure molding or hot molding, and an example thereof will be described below. However, the sintering temperature zone generally refers to 1000 ° C. or higher.

  There are various types of molding methods such as shock wave pressurization and pressure shearing. In this embodiment, the hot molding method will be described as an example. The high thermal stability exhibited by the F element-introduced rare earth iron alloy substituted with at least one of Al, Si, and Ga is particularly effective in a molding method using heating, and is a technical scope of the present invention. It is not intended to limit.

For example, the Sm ₂ Fe ₁₆ AlF _x magnetic powder produced in Example 1 is used as the magnetic powder. Of course, the pressure molding temperature differs depending on the type of substitution element and the presence or absence of Co addition, and this embodiment is not limited by the pressure molding temperature. In this example, an example in which a high-density bulk body is provided by pressure molding at a high temperature due to high thermal stability will be described.

The Sm ₂ Fe ₁₆ AlF _x magnetic powder was put in a nonmagnetic mold and pressed at 30 MPa in a 2T magnetic field to prepare a temporary molded body. At this time, as the slurry magnetic powder in the volatile nonpolar organic solvent, it is effective to perform the temporary molding because the orientation can be improved. The temporary molded body was put into a WC superhard metal mold (10 mm × 10 mm), and hot forming was performed at 350 ° C., 2 t, for 1 minute under a reduced pressure of 1 × 10 ⁻⁴ Pa. At that time, by coating the mold with a release agent such as BN, there are advantages in reducing the friction of the mold and releasing the molded article. For comparison, Sm ₂ Fe ₁₇ F _x magnetic powder was pressure-molded at 310 ° C. in the same manner. As a result, the bulk body of Sm ₂ Fe ₁₆ AlF _x magnetic powder showed a higher molding density than the bulk body of Sm ₂ Fe ₁₇ F _x . This difference is considered to be due to the difference in pressure molding temperature. Moreover, both magnetic properties were evaluated, and the saturation magnetic flux density of Sm ₂ Fe ₁₆ AlF _x magnetic powder was higher than that of Sm ₂ Fe ₁₇ F _x . Such an effect is similarly observed in the 3-29 system and the 1-12 system, and is also observed in a crystal structure formed by changing the three-dimensional stacking method of the CaCu ₅ structure. In addition, the use of a low melting point metal / alloy at the time of heat forming as described above may be effective in increasing the density. In particular, Zn, Sn, and their alloys form an alloy with Fe, which is unavoidably generated during the F element introduction reaction, by heating at least about 200 ° C., and this alloy exhibits paramagnetism at room temperature and therefore contributes as a reverse domain inversion site. In some cases, it is also effective in increasing the coercive force.

In this example, an example of an introduced rare earth iron alloy containing at least one of Al, Si, Ga, Cr, and Mn and containing at least one of Ti, Mo, and W will be described. This rare earth iron alloy is based on the empirical generation principle of R _mn T _{5m + 2n} (where R is a 4f transition element or Y, and T is Fe and Al, Si, Ga, Cr, Mn, Ti, Mo, W). This is a crystal structure based on a three-dimensional stacking method of CaCu5 structure. The composition ratio is RT19 composition, T cannot be stabilized only by Fe element, and must contain at least one of Al, Si, Ga, Cr, and Mn and at least one of Ti, Mo, and W. Don't be. The proportion of at least one of Al, Si, Ga, Cr, and Mn in T is 5% or more, and the proportion of at least one of Al, Si, Ga, Cr, and Mn in T is 5% or more, and the remaining 90 % Or less is Fe. With this composition ratio structure, the H, B, C, N, and F elements can be arranged at the intrusion position, and accordingly, the magnetic characteristics can be dramatically improved.

  The present invention can provide a hexagonal oxyfluoride ferrite magnet in which the magnetization of the entire crystal structure is increased.

Five crystal types have been reported to date for hexagonal oxyfluoride ferrite magnets, abbreviated as M-type, U-type, W-type, Y-type and Z-type, respectively. Among them, those having large uniaxial magnetic anisotropy are M type and W type, and these two types are hard magnetic materials. 1A shows an M-type crystal structure, and FIG. 1B shows a W-type crystal structure. The M-type crystal structure is represented by the general formula MFe ₁₂ O ₁₉ (M is an alkaline earth metal, a rare earth element and a divalent element), and the unit cell is composed of a unit twice the general formula. The unit cell is composed of two S blocks in the spinel layer and two R blocks in the bowl-shaped layer including M, and has mirror symmetry with respect to the c-axis direction on the plane including M. On the other hand, the W-type crystal structure is represented by the general formula MFe ₁₈ O ₂₇ , and the unit cell is composed of twice the general formula as before. In the unit cell, the R block is separated by two S blocks in one SR block of the M-type crystal structure.

  From the coordination of oxygen ions, tetracoordinate and hexacoordinate Fe sites are included in the S block, and pentacoordinate and hexacoordinate Fe sites are included in the R block. Since trivalent Fe ions have zero synthetic orbital angular momentum, there is no magnetic anisotropy at any coordination number. However, it has been pointed out from the first-principles calculation that the valence of pentacoordinated Fe ions deviates from the trivalence, and therefore the orbital angular momentum does not become zero. In addition, it is known that the synthetic orbital angular momentum does not completely freeze in Fe and Co due to the spread of the wave function. Therefore, the expected value of the combined orbital angular momentum does not become zero, and magnetic anisotropy occurs. The easy magnetization axis direction is the c-axis direction, and since it has uniaxial anisotropy, it functions as a hard magnetic material.

  Magnetization is ferrimagnetic, and Fe ions at the 2a site, 4e site, and 12k site are up-spinned, and Fe ions at the 4f1 site and 4f2 site have a down-spin magnetic structure. The magnetic interaction between Fe ions is dominated by super-exchange interaction via oxygen ions, and is determined to be ferromagnetic or antiferromagnetic depending on the angle formed by Fe ions-oxygen ions-Fe ions. From the Kanamori-Goodenough rule, it is concluded that in the case of trivalent Fe ions, the 180 ° superexchange interaction is simply antiferromagnetic and the 90 ° superexchange interaction is ferromagnetic. (However, P.W. Anderson's conclusion is reversed, but the Kanamori-Goodenough rule is appropriate for hexagonal ferrite). In the hexagonal ferrite, there is an interaction between each Fe site, and among these, the antiferromagnetic interaction is the largest and dominates the magnetism of the system.

  In the present invention, based on the above, attention is focused on oxygen ions (6h) that surround a five-coordinate Fe site and exist in a mirror surface containing M atoms. By replacing at least one 6h oxygen ion with a fluorine ion, the Fe ion is reduced. Therefore, the Fe ion at the 4e site or the 4f2 site has a value close to trivalent to divalent. By becoming divalent, the magnetic moment is reduced to 4 μB, and at the same time, orbital angular momentum is generated. By reducing the magnetic moment of the antiferromagnetic site, the magnetization of the entire crystal is improved.

  Furthermore, orbital angular momentum is generated by taking a value close to bivalence. An orbital angular momentum is generated between a pentacoordinate Fe ion and an adjacent six-coordinate Fe ion, thereby newly generating an anisotropic exchange interaction and a pseudodipole interaction. The former is generally called Dzyaloshinskii-Moriya antisymmetric interaction. From the above, the two pairs of spin Hamiltonians H of the system are described as follows.

  Where J is the superexchange interaction between atom 1 and atom 2, S1 and S2 are the spin magnetic moments of atom 1 and atom 2, respectively, and d is the secondary perturbation of the spin orbit interaction, given by . Γ represents a pseudodipole interaction.

  The first term of the spin Hamiltonian indicates the superexchange interaction between the spins, the second term indicates the Dzyaloshinskii-Moriya antisymmetric interaction, and the third term indicates the quasi-dipole interaction between the spins. The second term and the third term are caused by the survival of the orbital angular momentum. The order of the size of the second term and the third term is

It is. In general, the Dzyaloshinskii-Moriya antisymmetric interaction is the origin of weak ferromagnetism (parasitic ferromagnetism), and the antiferromagnetic interaction is six-coordinated Fe ²⁺ and Co ²⁺ with a second term size of d˜ The size of J cannot be ignored.

  Therefore, the Fe ions at the 4f2 site are nearly divalent, the orbital angular momentum survives, the Dzyaloshinskii-Moriya antisymmetric interaction is induced, and the 4f2 downspin deviates from the c-axis direction. As a result, the magnetization of the entire crystal increases.

The present invention was verified as follows using, for example, the M type. SrCO ₃ , SrO, α-Fe ₂ O ₃ and SrF ₂ were used as raw materials. The purity is 99% or more. From the viewpoint of promoting chemical reaction, the raw material is preferably fine powder, desirably 1 μm or less. Each raw material was pretreated to remove moisture. In addition, all the steps must be carried out in a dry atmosphere in which the humidity is controlled, and the dew point temperature is preferably 0 ° C. or lower, desirably −10 ° C. or lower. In this example, the dew point was −11 ° C.

The raw materials were weighed and mixed so as to be SrFe ₁₂ O _19-x F _x (X = 0 to 6). Since slight evaporation of Sr occurs during firing, it is necessary to weigh the raw material so that Sr is slightly larger than the stoichiometric ratio of the target composition. For mixing, for example, the mixture was performed using a mortar or ball mill. Then, it pressure-molded using the super steel metal mold | die, and produced the pellet. The higher the pressure, the better. The pressure of 20 MPa or more is necessary. In this example, the pressure was 50 MPa. The pressure-molded pellets were temporarily fired in a temperature range of 1100 ° C. to 1300 ° C. for 1 minute to 10 hours in a dry atmosphere. The oxygen partial pressure was set to 1 Pa or more. The temporarily fired molded body was coarsely pulverized to 10 μm or less using, for example, a mortar and a ball mill.

  In order to evaluate the obtained fine powder, the structure was evaluated by X-ray diffraction, and the composition was evaluated by a time-of-flight secondary ion mass spectrometer (TOF-SIMS). FIG. 2 (a) shows an X-ray diffraction pattern obtained at room temperature. It was confirmed that the same crystal structure as the M type ferrite was formed. Unreacted and heterogeneous phases were not observed by X-ray diffraction. From the composition analysis by TOF-SIMS, F element was detected in the matrix.

  In order to evaluate the magnetic properties, the fine powder was weighed and packed into a capsule, and hardened in a non-oriented state using stearic acid. This is because a vibrating sample magnetometer (VSM) is used to evaluate the intrinsic magnetic moment of a substance without being affected by density, grain shape, orientation, and the like. The saturation magnetization was evaluated by applying a magnetic field up to 6T. Compared with M-type Sr ferrite (x = 0), the magnitude of magnetization increased by 10%. It was found that the coercive force has almost the same value. In addition, as compared with M-type Sr ferrite (x = 0), the magnetic moment was increased by 10% in the whole crystal by neutron diffraction measurement.

In addition, Zn or Cu as an element that substitutes for the Fe element may increase the saturation magnetization. By substituting Zn for the Fe element at the down spin site and Cu ²⁺ being Yang-Teller ions, the crystal system induces a strain tilt magnetic structure, thereby increasing the saturation magnetization of the entire crystal system.

Claims (10)

A magnet material made of a binary or ternary alloy containing Fe, R [R is a 4f transition element or Y],
F is arranged at the penetration position of the Fe crystal lattice,
A magnet material, wherein a part of Fe site is substituted with at least one element of Al, Si and Ga. The magnet material according to claim 1,
The ternary alloy is R-Fe-T [R is a 4f transition element or Y, and T is a 3d transition element excluding Fe, Mo or Nb]. The magnet material according to claim 1,
The ternary alloy is R-Fe-T [R is a 4f transition element excluding W or Y, and T is W]. The magnet material according to claim 1,
The alloy material includes H, B, N, or C. The magnet material according to claim 1,
The alloy material has a crystal structure that can be formed by changing a three-dimensional stacking method of a CaCu ₅ structure. The magnet material according to claim 1,
The alloy has a Th ₂ Zn ₁₇ type structure, a Th ₂ Ni ₁₇ type structure, a CaCu ₅ type structure, an Nd ₃ (Fe, Ti) ₂₉ type structure, or a ThMn ₁₂ type structure. The magnet material according to claim 1,
A magnet material, wherein a distance between Fe existing before substitution at the substituted Fe site and Fe at another Fe site adjacent to the substituted Fe site is 0.245 nm or less. The magnet material according to claim 1,
The magnet material, wherein the substituted Fe site is a Fe site having Fe having a shortest distance from Fe of an adjacent Fe site among a plurality of Fe sites. The magnet material according to claim 1,
The magnet material, wherein the substitution amount of the substituted element is 2/17 at% or more and 1/3 at% or less with respect to the total amount of the alloy. The magnet material according to claim 1,
The alloy is produced by subjecting a rare earth iron alloy containing at least one element of Al, Si, and Ga to a fluorination heat treatment.