JP2017073544A - Method for manufacturing rare earth magnet - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a method for manufacturing a rare earth magnet having a ThMntype crystal structure, by which the residual magnetic flux density and anisotropy can be both increased at a time.SOLUTION: A method for manufacturing a rare earth magnet comprises the steps of: preparing a molten metal of an alloy having a composition expressed by the formula, (R1R2)(FeCo)TM; cooling the molten metal into flakes; pulverizing the flakes into powder; compacting the powder into a green compact; sintering the green compact into a sintered compact; and performing a hot working on the sintered compact at 700-1000°C.SELECTED DRAWING: Figure 3

Description

The present invention relates to a method for producing a rare earth magnet, and more particularly to a method for producing a rare earth magnet having a ThMn ₁₂ type crystal structure.

  The application of permanent magnets extends to a wide range of fields such as electronics, information communication, medical care, machine tool fields, industrial and automotive motors, and the demand for suppression of carbon dioxide emissions is increasing. In recent years, there are increasing expectations for the development of permanent magnets with even higher characteristics due to energy savings and improved power generation efficiency in the industrial field.

  At present, Nd—Fe—B magnets, which are dominating the market as high-performance magnets, are also used in drive motor magnets for HV / EHV. Recently, new permanent magnet materials are being developed in response to the pursuit of further miniaturization and higher output of motors (increase in residual magnetic flux density of magnets).

As one of the development of materials having performance exceeding that of Nd—Fe—B based magnets, research on rare earth magnets having a ThMn ₁₂ type crystal structure has been advanced.

For example, Patent Document 1 discloses a rare earth magnet made of an alloy including a hard magnetic phase having a ThMn ₁₂ type tetragonal structure and a nonmagnetic phase. Then, the composition of the alloy is _{[R 1-a (M1)} a] [T 1-b-c (M2) b (M3) c] d X α ( wherein, R at least one rare earth element is the (Y M1 is at least one selected from the group consisting of Zr and Hf, T is at least one selected from the group consisting of Fe, Co, and Ni, and M2 is Cu, Bi, Sn, Mg, In, and Pb M3 is at least one selected from the group consisting of Al, Ga, Ge, Zn, B, P, and S, and X is Si, Ti, V, Cr, Mn, Nb, At least one selected from the group consisting of Mo, Ta, and W, and 0 ≦ a ≦ 0.6, 0.01 ≦ b ≦ 0.20, 0 ≦ c ≦ 0.05, 7 ≦ d ≦ 11, And 0.5 ≦ α ≦ 2.0 (atomic ratio)) . It is further disclosed to hot work the alloy.

It is known that a rare earth magnet having a ThMn ₁₂ type crystal structure can orient the easy magnetization axis (c axis) in a direction perpendicular to the compression direction by compressing the rare earth magnet hot. .

For example, Non-Patent Document 1 discloses that an easily cast Nd—Fe—V alloy having a ThMn ₁₂ type crystal structure is compressed by hot forging so that the easy axis of magnetization is perpendicular to the compression direction. Orientation in the direction is disclosed.

JP 2001-189206 A

S. Rivoirard et al. , "Texture development in Nd-Fe-V and Nd-Fe-Balloy by hot forging permanent magnet properties", Solid State. 105 (2005), pp. 291-296, (CRETA-CNRS, FRANCE)

In the rare earth magnet disclosed in Patent Document 1, when the residual magnetic flux density is improved, a ThMn ₁₂ type crystal structure may not be stably obtained.

When the ratio of the portion having a ThMn ₁₂ type crystal structure in the rare earth magnet is reduced, the orientation as disclosed in Non-Patent Document 1 is weak even if the rare earth magnet is hot compressed.

Accordingly, the present inventors have found a problem that it is necessary to simultaneously increase both the residual magnetic flux density and the anisotropy in a rare earth magnet having a ThMn ₁₂ type crystal structure.

The present invention can improve both the residual magnetic flux density and anisotropic simultaneously, and to provide a method for producing a rare earth magnet having a crystal structure of _12-inch ThMn.

As a result of intensive studies, the inventors have arrived at the present invention. The gist of the present invention is as follows.
<1> formula _{(R1 (1-x) R2} x) p (Fe (1-y) Co y) q T r M s
(In the above formula,
R1 is one or more rare earth elements selected from the group consisting of Sm, Nd, Pm, Er, Tm, and Yb,
R2 is one or more elements selected from the group consisting of Zr, La, Ce, Pr, Eu, Gd, Tb, Dy, Ho, and Lu;
T is one or more elements selected from the group consisting of Ti, V, Mo and W,
M is an inevitable impurity element and one or more elements selected from the group consisting of Al, Cr, Cu, Ga, Ag and Au,
0 ≦ x ≦ 0.5,
0 ≦ y ≦ 0.7,
4 ≦ p ≦ 20,
q = 100−pr−s,
3.8 ≦ r ≦ 5.8,
0 ≦ s ≦ 3)
Preparing a molten metal having a composition represented by:
Cooling the molten metal to obtain flakes,
Crushing the flakes to obtain a powder;
Pressing the powder to obtain a green compact;
Sintering the green compact to obtain a sintered body; and
Hot-working the sintered body at 700 to 1000 ° C .;
A method for producing a rare earth magnet, comprising:
<2> The method according to <1>, wherein the hot working is hot upsetting.
<3> The method according to <1> or <2>, wherein the hot working is hot backward extrusion.

According to the present invention, it is possible to obtain a method for producing a rare earth magnet having a ThMn ₁₂ type crystal structure, which can simultaneously improve both the residual magnetic flux density and the anisotropy.

It is a graph showing the stable region of the T component of R1Fe _{12-x T} x in the alloy. It is the schematic of a strip cast apparatus. It is a longitudinal cross-sectional schematic diagram which shows an example of a hot upsetting process. It is a longitudinal cross-sectional schematic diagram which shows an example of a hot back extrusion process. It is a figure which shows the magnetic characteristic after the hot upsetting process of the rare earth magnet which concerns on Example 3. FIG. It is a figure which shows the X-ray-diffraction (XRD) analysis result before and behind the hot upsetting process of the rare earth magnet which concerns on Example 3. FIG.

  Hereinafter, embodiments of a method for producing a rare earth magnet according to the present invention will be described in detail. In addition, embodiment shown below does not limit this invention.

Method for producing a rare earth magnet according to the present invention, first, a melt of the formula _{(R1 (1-x) R2} x) p (Fe (1-y) Co y) q T r M s an alloy of composition expressed prepare.

(R1)
R1 is a rare earth element and is an essential component for the rare earth magnet obtained by the production method of the present invention to exhibit magnetic properties. R1 is one or more rare earth elements selected from the group consisting of Sm, Nd, Pm, Er, Tm, and Yb, and it is preferable to use Nd and Sm.

(R2)
R2 is one or more elements selected from the group consisting of Zr, La, Ce, Pr, Eu, Gd, Tb, Dy, Ho, and Lu, and a part of R1 is substituted to form a ThMn ₁₂ type crystal. Contributes to the stabilization of the structure. The R2 element substitutes for the R1 element in the ThMn ₁₂ type crystal and causes the crystal lattice to contract. Accordingly, when the alloy is raised to a high temperature or when nitrogen atoms or the like are penetrated into the crystal lattice, the ThMn ₁₂ type crystal structure is stably maintained. On the other hand, in terms of magnetic characteristics, strong magnetic anisotropy derived from the R1 element is reduced by R2 substitution.

The substitution rate x of R1 with R2 is determined from both aspects of crystal stability and magnetic properties. However, in the present invention, addition of R2 is not essential. When the substitution amount of R2 is 0, the ThMn ₁₂ type crystal structure can be stabilized by adjusting the alloy composition and heat treatment, and the anisotropic magnetic field is increased. On the other hand, if the substitution rate x is 0.5 or less, the anisotropic magnetic field will not be significantly reduced. The substitution rate x is preferably 0 ≦ x ≦ 0.3.

  The blending amount p of R1 and R2 is the sum of R1 and R2, and is 4 to 20 atomic%. If the blending amount p is 4 atomic% or more, the precipitation of the α-Fe phase is not remarkable, the volume fraction of the Fe phase can be lowered after the heat treatment, and the performance as a rare earth magnet can be sufficiently exhibited. it can. The blending amount p is preferably 7 atomic% or more. On the other hand, if the blending amount p is 20 atomic% or less, the grain boundary phase becomes excessive, and as a result, the residual magnetic flux density does not decrease. The blending amount p is preferably 15 atomic% or less.

(T)
T is one or more elements selected from the group consisting of Ti, V, Mo and W. FIG. 1 shows the stabilization region of the T element in the R1Fe _12-x T _x compound (source: KHJ Buschuw, Rep. Prog. Phys. 54, 1123 (1991)). It is known that the addition of Ti, V, Mo, and W as the third element to the binary alloy stabilizes the ThMn ₁₂ type crystal structure and exhibits excellent magnetic properties.

Conventionally, in order to stabilize the ThMn ₁₂ type crystal structure, it has been said that a relatively large amount of T component needs to be added to the alloy. Thereby, the content rate of the Fe component which comprises an alloy became low, and the occupation site of the Fe atom which has the most influence on the residual magnetic flux density was replaced with, for example, T atoms, and the overall residual magnetic flux density was lowered. In order to improve the residual magnetic flux density, the amount of T should be reduced. However, in this case, stabilization of the ThMn ₁₂ type crystal structure is impaired. This is because if the T that stabilizes the ThMn ₁₂ type crystal structure is reduced, the stabilization of the ThMn ₁₂ type crystal structure is impaired, and α- (Fe, Co, which becomes an obstacle to an anisotropic magnetic field or coercive force. ) Will precipitate.

Therefore, the blending amount of T is set to the minimum necessary for stabilizing the ThMn ₁₂ type crystal structure. That is, the blending amount r of T is 3.8 to 5.8 atomic%. When the blending amount r is 3.8 atomic% or more, the stability of the ThMn ₁₂ type crystal structure is not impaired. . On the other hand, if the blending amount r is 5.8 atomic% or less, the Fe content in the alloy will not be too low, and the residual magnetic flux density of the obtained rare earth magnet will not be reduced.

(M)
M is an inevitable impurity element and one or more elements selected from the group consisting of Al, Cr, Cu, Ga, Ag, and Au. These elements are elements that are inevitably mixed in the raw material and / or the manufacturing process.

  The smaller the amount M of s, the better, and it may be 0 atomic%. However, in order not to use an excessively high purity raw material, it is preferable that the compounding amount s of M is 0.1 atomic% or more. On the other hand, when the compounding amount s of M is 3 atomic% or less, the performance of the obtained rare earth magnet is not substantially deteriorated. The compounding amount s of M is preferably 1 atomic% or less.

(Fe and Co)
The alloy used in the present invention contains Fe and Co in addition to R1, R2, T, and M described so far. Co substitutes a part of Fe to increase the spontaneous magnetization and improve both the anisotropic magnetic field and saturation magnetization characteristics by the Slater poling rule. The effect of the Slater polling rule is exhibited when the substitution rate of Fe by Co is 0.5 or less. In addition, by substituting part of Fe with Co, the Curie point of the compound is raised, so that a decrease in residual magnetic flux density at a high temperature is suppressed.

  Fe and Co are the balance of R1, R2, T, and M in the alloy. Therefore, the compounding quantity q (atomic%) of Fe and Co is represented by 100-pr-s.

(Preparation of molten metal)
An alloy having the composition described so far is melted to form a molten metal. The melting method is not particularly limited, but a melting method in which the composition of the alloy does not change during melting is preferable. For example, high frequency dissolution is preferable. During melting, if a specific component is depleted due to evaporation or the like, or if a specific component is formed as an oxide and discharged as slag, the above-mentioned amount is taken into account during the cooling described below. The raw materials are weighed so that the alloy composition is obtained, and the raw materials are melted to form a molten metal.

(Cooling of molten metal)
In the manufacturing method of the present invention, since the minimum amount of T is ensured in order to maintain the stability of the ThMn ₁₂ type crystal structure, the cooling rate of the molten metal does not need to be specifically specified. From the viewpoint of further improving the stability, it is preferable to quench the molten metal.

The temperature at which the ThMn ₁₂ type crystal precipitates is lower than the temperature at which α- (Fe, Co) precipitates. Therefore, if the α- (Fe, Co) precipitation is further reduced by not staying in the vicinity of the temperature at which α- (Fe, Co) precipitates, and as many ThMn ₁₂ type crystals as possible are produced, The stability of the ThMn ₁₂ type crystal structure can be further improved.

That is, the cooling rate of the molten metal is preferably 1 × 10 ² to 1 × 10 ⁷ K / sec. . If the speed is 1 × 10 ² K / sec or more, the precipitation of α- (Fe, Co) can be further reduced in the flakes, and more ThMn ₁₂ type crystals can be generated. The speed is more preferably 1 × 10 ³ K / second or more. On the other hand, if the speed is 1 × 10 ⁷ K / second or less, the apparatus for cooling does not become too complicated. The speed is more preferably 1 × 10 ⁶ K / second or less.

When the molten metal is cooled at a rate of 1 × 10 ² to 1 × 10 ⁷ K / sec, the average crystal grain size of the flakes is 10 to 10,000 nm. In addition, the average particle diameter in this case is an average value of the equivalent circle diameters when the diameter of each crystal is the equivalent circle diameter.

As the cooling device, for example, a strip cast device 10 as shown in FIG. 2 can be used. In the strip cast apparatus 10, the alloy raw material is melted in the melting furnace 11, and has a composition represented by the formula (R _{1 (1-x)} R 2 _x ) _p (Fe _(1-y) Co _y ) _q T _r M _S. A molten alloy 12 is prepared. The molten metal 12 is supplied to the tundish 13 at a constant supply amount. The molten metal 12 supplied to the tundish 13 is supplied from the end of the tundish 13 to the cooling roll 14 by its own weight or Ar differential pressure.

  Here, the tundish 13 is made of ceramics or the like, temporarily stores the molten metal 12 continuously supplied from the melting furnace 11 at a predetermined flow rate, and can rectify the flow of the molten metal 12 to the cooling roll 14. . The tundish 13 also has a function of adjusting the temperature of the molten metal 12 immediately before reaching the cooling roll 14.

  The cooling roll 14 is made of a material having high thermal conductivity such as copper or chromium, and the roll surface is subjected to chromium plating or the like in order to prevent erosion with a high-temperature molten metal. This roll can be rotated in the direction of the arrow at a predetermined rotational speed by a driving device (not shown). By controlling this rotation speed, the cooling rate of the molten metal can be controlled.

  The molten alloy 12 is cooled on the outer periphery of the cooling roll 14. As a result, the molten alloy 12 is solidified into thin pieces 15. The thin piece 15 is peeled off from the cooling roll 14. The peeled thin piece 15 is recovered using a recovery device (not shown).

(Fracture of flakes)
The collected flakes 15 are pulverized to obtain a powder. The pulverization method may be a conventional method, for example, a method using a cutter mill. Although there will be no restriction | limiting in particular if the particle size of powder will be the particle size which can sinter later mentioned, 1-1000 micrometers is preferable. If the particle size is 1 μm or more, excessive man-hours are not required for grinding. The particle size is more preferably 10 μm or more. On the other hand, if the particle size is 1000 μm or less, there is no problem in sintering. The particle size is more preferably 200 μm or less.

(Powder compaction)
The powder is compacted to obtain a compact. The compacting method may be a conventional method. For example, the powder may be charged into a mold and compressed using a press.

(Sintering of compacts)
The green compact is sintered to obtain a sintered body. The sintering method may be a conventional method, and examples thereof include a spark plasma sintering (SPS) method, a hot press using high frequency heating, and a hot press using condensing heating. The spark plasma sintering method, hot pressing by high-frequency heating, and hot pressing by condensing heating can rapidly raise the green compact to a desired temperature, and the crystal grains are reduced before the green compact reaches the desired temperature. It is preferable at the point which can prevent coarsening.

  What is necessary is just to select a sintering temperature suitably with an alloy composition, and 700-800 degreeC is preferable. If the sintering temperature is 700 ° C. or higher, the green compact can be sintered in a relatively short time. The sintering temperature is more preferably 720 ° C. or higher. On the other hand, if the sintering temperature is 800 ° C. or lower, the crystal grains will not be coarsened during sintering. The sintering temperature is more preferably 780 ° C. or lower.

  During sintering, in order to prevent the green compact from being deformed due to expansion, it is preferable to pressurize the lid of the mold containing the green compact at 100 to 500 MPa. If the applied pressure is 100 MPa or more, the expansion force can be countered. On the other hand, if the applied pressure is 500 MPa or less, the green compact is not plastically deformed during sintering.

  What is necessary is just to determine a sintering temperature suitably according to the mass of a sintered compact. The sintering atmosphere is not particularly limited, but an argon gas atmosphere is preferable in order to prevent the sintered body from being oxidized.

(Hot processing of sintered body)
Hot processing which compresses the obtained sintered compact at 700-1000 degreeC is performed. If the temperature of the hot working is 700 ° C. or higher, the sintered body is softened and the sintered body can be compressed relatively easily. As a result, it is possible to obtain a rare earth magnet having a high anisotropy in which the easy axis is oriented in a direction perpendicular to the compression direction. From this viewpoint, the hot working temperature is preferably 800 ° C. or higher.

  On the other hand, if the temperature of the hot working is 1000 ° C. or less, the crystal grains are not coarsened and the magnetic properties are not deteriorated. Further, the sintered body is excessively softened, and the sintered body is not damaged when compressing and deforming the sintered body. From these viewpoints, the hot working temperature is preferably 900 ° C. or lower.

The processing rate at the time of hot processing is defined as follows.
Processing rate (%) =
(Compression direction thickness before compression-Compression direction thickness after compression) / (Compression direction thickness before compression) × 100

  About a processing rate, 30 to 90% is preferable. If the processing rate is 30% or more, the orientation effect of the easy axis described above can be obtained. On the other hand, if the processing rate is 90% or less, the sintered body will not be damaged due to excessive compression.

  About the method of hot processing, if the temperature at the time of hot processing is the range mentioned above, the method will not be specifically limited. In the hot working in the present invention, even if one direction of the sintered body is enlarged within the above-described temperature range, it is only necessary to compress at least another direction of the sintered body. Examples of hot working methods include hot upsetting and hot backward extrusion. Each will be described next.

(Hot upsetting)
FIG. 3 is a schematic longitudinal sectional view showing an example of hot upsetting. FIG. 3A shows a state immediately before hot upsetting the sintered body, and FIG. 3B shows a state when the sintered body is hot upsetting.

  The hot upsetting mold 50 includes a movable mold 52 and a fixed mold 54. The sintered body 56 is installed between the movable mold 52 and the fixed mold 54. The sintered body 56 is deformed by moving the movable mold 52 in a direction approaching the fixed mold 54. That is, the sintered body 56 is deformed by moving the movable mold 52 in the direction of the arrow in FIG.

  In FIG. 3B, regarding the sintered body 56, the vertical direction on the paper surface is compressed, but the horizontal direction on the paper surface is enlarged. That is, the compression direction in the hot upsetting process is the vertical direction on the paper surface in FIG. Therefore, by the hot upsetting process, the easy magnetization axis is oriented in the direction perpendicular to the compression direction, that is, in the left-right direction on the paper surface in FIG. In other words, the easy axis is oriented in a direction perpendicular to the upsetting direction. The thickness in the compression direction of the sintered body before deformation refers to the distance between the movable mold 52 and the fixed mold 54 before deformation, and the thickness in the compression direction of the sintered body after deformation is the movable thickness at the end of deformation. This is the distance between the mold 52 and the fixed mold 54.

(Hot rear extrusion)
FIG. 4 is a schematic longitudinal sectional view showing an example of hot backward extrusion. FIG. 4A shows a state immediately before the sintered body is subjected to hot rear extrusion, and FIG. 4B shows a state when the sintered body is backward extruded.

  The rear extrusion mold 60 includes a fixed mold 62 and a movable mold 64. The fixed mold 62 includes an insertion hole 63 that accommodates the movable mold 64 so as to be slidable.

  Only one of the insertion holes 63 is open, and the movable mold 64 is inserted through the opening. The movable mold 64 slides in the axial direction of the insertion hole 63 while the outer wall of the movable mold 64 is in contact with the peripheral wall of the insertion hole 63. As described above, as long as the movable mold 64 can slide, the shapes of the outer wall of the movable mold 64 and the peripheral wall of the insertion hole 63 are not particularly limited, and are, for example, a cylindrical shape or a quadrangular prism shape.

  The movable mold 64 includes a through hole 65. The shape of the through-hole 65 is not limited as long as it can extrude a deformed sintered body 66 described later. For example, the through-hole 65 has a columnar shape or a quadrangular prism shape.

  The sintered body 66 is installed on the opposite side of the opening of the insertion hole 63, that is, on the bottom of the insertion hole 63 in FIG. Then, when the movable die 64 is inserted into the insertion hole 63 and the sintered body 66 is crushed by the end face of the movable die 64 as shown in FIG. Extruded. Extrusion is performed until the end surface of the movable mold 64 reaches the end surface of the insertion hole 63.

  In FIG. 4B, regarding the sintered body 66, the horizontal direction on the paper surface is compressed, but the vertical direction on the paper surface is enlarged. That is, the compression direction in the backward extrusion process is the left-right direction in FIG. 4B. Therefore, the easy magnetization axis is oriented in the direction perpendicular to the compression direction, that is, in the vertical direction on the paper surface in FIG. In other words, the easy axis of magnetization is oriented in a direction parallel to the extrusion direction. When the cross section perpendicular to the axis of the insertion hole 63 and the through hole 65 is a circle, the thickness in the compression direction of the sintered body before deformation is the cross section perpendicular to the axis of the insertion hole 63. The thickness in the compression direction of the sintered body after deformation refers to the diameter of the cross section perpendicular to the axis of the through-hole 65. When the cross section perpendicular to the axis of the insertion hole 63 and the through hole 65 is not a circle, the thickness in the compression direction of the sintered body before deformation is the equivalent circle diameter of the cross section perpendicular to the axis of the insertion hole 63. The thickness in the compression direction of the sintered body after deformation refers to the equivalent circle diameter of the cross section perpendicular to the axis of the through hole 63.

  Hereinafter, the present invention will be described more specifically with reference to examples. The present invention is not limited to the conditions used in the following examples.

(Examples 1-4)
A molten alloy having a composition represented by Sm _7.7 (Fe _0.75 Co _0.25 ) _88.5 Ti _3.8 was prepared, and the apparatus shown in FIG. 2 was used to prepare 2 × 10 ⁶ K / The molten metal was cooled at a rate of seconds to produce a flake.

  The flakes were pulverized using a cutter mill to obtain powder. The powder was inserted into a mold to produce a green compact. The obtained green compact was sintered by a hot press method using high frequency heating. The sintering temperature was 750 ° C., and the sintering atmosphere was an argon gas atmosphere. During the sintering, the green compact was pressurized at 400 MPa.

  The obtained sintered body was hot upset processed. The hot upsetting temperature was 700 ° C. in Example 1, 800 ° C. in Example 2, and 900 ° C. in Example 3. For all Examples 1 to 3, the processing rate was 75%.

(Examples 4 and 5)
A molten alloy having a composition represented by Sm _7.7 (Fe _0.75 Co _0.25 ) _88.5 Ti _3.8 was prepared, and the apparatus shown in FIG. 2 was used to prepare 2 × 10 ⁶ K / The molten metal was cooled at a rate of seconds to produce a flake.

  The flakes were pulverized using a cutter mill to obtain powder. The powder was inserted into a mold to produce a green compact. The obtained green compact was sintered by a hot press method using high frequency heating. The sintering temperature was 750 ° C., and the sintering atmosphere was an argon gas atmosphere. During the sintering, the green compact was pressurized at 200 MPa.

  The obtained sintered body was hot back extruded. About the temperature of hot back extrusion, Example 4 was 1000 degreeC and Example 5 was 900 degreeC. In both Examples 4 and 5, the processing rate was 70%.

(Example 6)
(Sm _0.8 Zr _0.2 ) _7.7 Fe _0.75 Co _0.25 ) The same as Example 3 except that a molten alloy having a composition represented by _86.5 Ti _5.8 was prepared. In addition, a rare earth magnet of Example 7 was manufactured.

(Example 7)
(Nd _0.9 Zr _0.1 ) _7.7 Fe _0.75 Co _0.25 ) The same as Example 3 except that a molten alloy having a composition represented by _86.5 Ti _5.8 was prepared. In addition, a rare earth magnet of Example 7 was manufactured.

(Reference Example 1)
The rare earth magnet of Reference Example 1 was prepared in the same manner as in Example 2 except that a molten alloy having a composition represented by Nd _13.8 Fe _75.7 Co _4.5 B _5.5 Ga _0.5 was prepared. Produced.

(Reference Example 2)
Other than having prepared a molten metal having a composition represented by Nd _13.8 Fe _75.7 Co _4.5 B _5.5 Ga _0.5 and having set the temperature of hot backward extrusion to 800 ° C. Produced a rare earth magnet of Reference Example 2 in the same manner as in Example 4.

(Reference Example 3)
A rare earth magnet of Reference Example 3 was produced in the same manner as in Example 1 except that the hot upsetting process was not performed on the sintered body.

(Reference Example 4)
A rare earth magnet of Reference Example 4 was produced in the same manner as in Example 7 except that the hot upsetting process was not performed on the sintered body.

(Evaluation)
Each of the obtained rare earth magnets was cut into 2 mm squares, and the magnetic properties were evaluated using a vibrating sample magnetometer (VSM). In addition, each of the obtained rare earth magnets was analyzed by X-ray diffraction (XRD: X Ray Diffraction).

  The evaluation results are shown in FIGS.

  FIG. 5 is a diagram illustrating the magnetic characteristics of the rare earth magnet according to Example 3 after hot upsetting. A solid line indicates magnetic characteristics in a direction perpendicular to the compression direction, and a broken line indicates magnetic characteristics in a direction parallel to the compression direction. As can be seen from FIG. 5, the magnetic properties in the direction perpendicular to the compression direction are better than the magnetic properties in the direction parallel to the compression direction. Therefore, it can be said that the easy axis of the rare earth magnet of Example 3 is oriented in a direction perpendicular to the compression direction.

FIG. 6 is a diagram showing the results of X-ray diffraction (XRD) analysis before and after hot upsetting of the rare earth magnet according to Example 3. The solid line shows the analysis result after hot upsetting, and the broken line shows the analysis result before hot upsetting. In FIG. 6, the peak position of the NdFe ₁₁ Ti phase having a ThMn ₁₂ type crystal structure is indicated by a black square, and the peak position of the α-Fe phase is indicated by a black triangle.

As can be seen from FIG. 6, the peak intensities of the (220), (320), (410), and (330) planes derived from the a direction of the ThMn ₁₂ type crystal and the peak intensities such as (301) close thereto. However, it has become stronger after hot installation. Therefore, it can be said that the c-plane direction, which is the easy axis of magnetization of the rare earth magnet of Example 3, is oriented perpendicular to the compression direction. Note that when such an orientation is strengthened, the peak position of the NdFe ₁₁ Ti phase apparently shifts, but the actual peak position is the same before and after hot upsetting.

  In Table 1, Br (parallel) represents the residual magnetic flux density in the direction parallel to the compression direction, and Br (vertical) represents the residual magnetic flux density in the direction perpendicular to the compression direction. The orientation degree ω was defined as ω = {Br (vertical) −Br (parallel)} / Br (vertical). When ω is greater than 0 and less than or equal to 1, the easy magnetization axis of the rare earth magnet is oriented in a direction perpendicular to its compression direction.

  For reference, for each rare earth magnet, hexagons A, B, and C are represented by A: a six-membered ring composed of Fe (8i) and Fe (8j) sites centered on rare earth atoms, and B: Fe. 6-membered ring composed of Fe (8i) and Fe (8j) sites centered on (8i) -Fe (8i) dumbbells, Fe (8j) centered on the line of C: Fe (8i) -rare earth atoms ) And a Fe (8f) site, the length in the a-axis direction of Hexagon A: Hex (A) is also shown.

  As is clear from Table 1, in all of Examples 1 to 7, it was confirmed that ω was greater than 0 and 1 or less, and the easy axis of the rare earth magnet was oriented in a direction perpendicular to the compression direction. . Moreover, about Examples 1-7, it confirmed that Br (vertical) was higher than the reference examples 3 and 4 which did not perform hot processing. On the other hand, in Reference Examples 1 and 2, it was confirmed that ω was not larger than 0 and did not satisfy 1 or less, and the easy magnetization axis of the rare earth magnet was not oriented in the direction perpendicular to the compression direction. In Reference Examples 3 and 4, ω = 0 because no hot working was performed, and no orientation (anisotropy) was observed.

According to the present invention, it is possible to obtain a method for producing a rare earth magnet having a ThMn ₁₂ type crystal structure, which can simultaneously improve both the residual magnetic flux density and the anisotropy. Therefore, the present invention has great industrial applicability.

DESCRIPTION OF SYMBOLS 10 Strip cast apparatus 11 Melting furnace 12 Molten metal 13 Tundish 14 Cooling roll 15 Thin piece 50 Mold for hot upsetting 52 Movable type 54 Fixed mold 56 Sintered body 60 Mold for back extrusion process 62 Fixed mold 63 Insertion hole 64 Movable type 65 Through hole 66 Sintered body

Claims (3)

Formula _{(R1 (1-x) R2} x) p (Fe (1-y) Co y) q T r M s
(In the above formula,
R1 is one or more rare earth elements selected from the group consisting of Sm, Nd, Pm, Er, Tm, and Yb,
R2 is one or more elements selected from the group consisting of Zr, La, Ce, Pr, Eu, Gd, Tb, Dy, Ho, and Lu;
T is one or more elements selected from the group consisting of Ti, V, Mo and W,
M is an inevitable impurity element and one or more elements selected from the group consisting of Al, Cr, Cu, Ga, Ag and Au,
0 ≦ x ≦ 0.5,
0 ≦ y ≦ 0.7,
4 ≦ p ≦ 20,
q = 100−pr−s,
3.8 ≦ r ≦ 5.8,
0 ≦ s ≦ 3)
Preparing a molten metal having a composition represented by:
Cooling the molten metal to obtain flakes,
Crushing the flakes to obtain a powder;
Pressing the powder to obtain a green compact;
Sintering the green compact to obtain a sintered body; and
Hot-working the sintered body at 700 to 1000 ° C .;
A method for producing a rare earth magnet, comprising:   The method according to claim 1, wherein the hot working is hot upsetting.   The method of claim 1, wherein the hot working is a hot back extrusion process.

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