R-Fe-B based rare earth magnetic material and method for producing the same

Abstract

This invention provides a high-performance R-Fe-B magnetic material that does not use heavy rare earth elements, and a method for producing the same. [Solution] An R-Fe-B rare earth magnetic material comprising a main phase, a two-particle grain boundary phase, and a triangular region grain boundary phase, wherein the main phase is R2T 14 The B phase is such that R is at least one of Nd and Pr, T is Fe, or Fe and Co, and the two-particle grain boundary phase includes a grain boundary Cu-rich phase and / or an acicular ZrBx phase and / or an acicular TiBx phase, and the atomic percentage ratio of the grain boundary Cu-rich phase with R within 20 nm near the interface with the main phase is 1.5

Description

[Technical Field]

[0001] The present invention belongs to the field of R-Fe-B magnetic materials, and more particularly to R-Fe-B rare earth magnetic materials and methods for producing the same. [Background technology]

[0002] In recent years, R-Fe-B rare earth magnetic materials have developed rapidly and are widely used in high-tech fields such as new energy vehicles, air conditioning compressors, and robots. However, high-performance magnetic materials containing heavy rare earth elements are expensive, making cost reduction a crucial issue. Therefore, the demand for inexpensive, high-performance R-Fe-B rare earth magnetic materials that do not contain heavy rare earth elements is growing ever stronger.

[0003] The heavy rare earth grain boundary diffusion method is the most effective way to improve the magnetic properties of R-Fe-B rare earth magnets. However, heavy rare earth elements are scarce and expensive, so manufacturing costs remain high. With the rapid increase in demand for high-performance R-Fe-B rare earth magnets and the soaring prices of heavy rare earth elements, reducing manufacturing costs without using heavy rare earth elements has become a crucial issue. Since the main manufacturing cost of R-Fe-B rare earth magnets lies in the ratio of heavy rare earth elements, the production of R-Fe-B rare earth magnets with low or no heavy rare earth element content has become a major research topic for researchers. Realizing high-performance R-Fe-B rare earth magnets that do not contain heavy rare earth elements is a crucial research theme.

[0004] Chinese Patent Publication CN110299237B describes R2T as an RTB-based sintered magnetic material. 14 An invention is disclosed that includes a plurality of main phase particles containing B phase crystals and a plurality of grain boundary multipoints as grain boundary phases surrounded by at least three of the main phase particles, wherein the plurality of grain boundary multipoints include at least two phases: a transition metal-rich phase and an R-rich phase. The R-rich phase is divided into at least two phases: a Cu-poor phase and a Cu-rich phase. This type of magnetic material exhibits excellent performance when it contains small amounts of heavy rare earth elements, but presents challenges in terms of manufacturing cost.

[0005] In addition, the Chinese Patent Publication No. CN111724959B discloses an invention of an R-T-B-based permanent magnet, which contains Ga, R is one or more rare earth elements, T is Fe or Fe and Co, and B is boron. The aforementioned magnet has a main phase composed of crystal particles having an R2T 14 B-type crystal structure, and includes grain boundaries formed by two or more adjacent main phase particles. The grain boundaries contain R6T 13 Ga phase. This type of magnet also has problems in terms of manufacturing cost because it has excellent performance even when it contains a small amount of heavy rare earth elements.

Prior Art Documents

Patent Documents

[0006]

Patent Document 1

Patent Document 2

Disclosure of the Invention

Problems to be Solved by the Invention

[0007] Different from both of the above two inventions, the present invention aims to provide a high-performance rare earth magnet with a new structure and a manufacturing method thereof while suppressing the manufacturing cost without containing any heavy rare earth elements.

Means for Solving the Problems

[0008] To achieve the above object, the R-Fe-B-based rare earth magnet according to the first invention of the present application includes a main phase, a two-particle grain boundary phase, and a triangular region grain boundary phase, The components of the R-Fe-B-based rare earth magnet are in mass percentage representation R: 29.5 to 33.5%, B: 0.85 to 1.05%, Al: 0.05 to 0.8%, Cu: 0.4 to 2.0%, Ga: 0.3 to 0.6%, Co: 0.5 to 2.0%, Zr or Ti: 0.15 to 0.5%, the balance being Fe and inevitable impurities, The main phase is R2T 14 which is a B phase, where R is at least one of the light rare earth elements, and T is Fe, or Fe and Co, The two-particle grain boundary phase includes a grain boundary Cu-rich phase, and / or a needle-like ZrBx phase, and / or a needle-like TiBx phase, In the grain boundary Cu-rich phase, the atomic percentage ratio of Cu to R within 20 nm near the interface with the main phase is 1.5 < Cu / R ≤ 2.0, and the atomic percentage ratio of Cu to R within 50 nm at the center of the grain boundary is 1.0 < Cu / R ≤ 1.5, The width of the needle-like ZrBx phase is 10 to 50 nm, The width of the needle-like TiBx phase is 10 to 50 nm, The triangular region grain boundary phase is an FCC-type NdOx phase, a RCu phase with an atomic percentage ratio of Cu to R of 1.0 < Cu / R ≤ 1.5, a RGa phase with an atomic percentage ratio of Ga to R of 0.1 < Ga / R ≤ 0.5, a 6:14 phase with an atomic percentage ratio of R to the transition metal element, a R-rich phase, and / or the needle-like ZrBx phase, and / or the needle-like TiBx phase, and comprises which is characterized in that.

[0009] Also in one embodiment, the R is at least one of Nd, Pr, La, Ce, Sm.

[0010] The second invention of the present application is a method for manufacturing the above R-Fe-B-based rare earth magnet, comprising Step 1: Performing hydrogen storage and dehydrogenation treatment on an alloy prepared by mixing each raw material based on mass percentage, and respectively preparing alloy thin flakes of the main phase alloy and related phases after the hydrogenation treatment, Step 2: The alloy flakes of the main phase alloy and the alloy flakes of the related phase obtained in Step 1 are mixed in a predetermined ratio, crushed by machine, the alloy of the related phase is dissolved in the main phase alloy, the mixed powder material is crushed into a powder, and a base material is prepared by magnetic field molding. Step 3: The material is placed in a sintering furnace and sintered, then kept warm to create a sintered blank, and then the sintered blank is subjected to a two-stage aging treatment. In the aforementioned two-stage prescription process, The primary aging temperature is 700-900°C, and the incubation time is 3-7 hours. The secondary aging temperature is 430-550°C, and the incubation time is 3-7 hours. It is characterized by the following:

[0011] In one embodiment, the hydrogen storage temperature in step 1 is 50 to 100°C, the dehydrogenation temperature is 500 to 650°C, and the dehydrogenation time is 3 to 7 hours.

[0012] In one embodiment, the particle size of the mechanical grinding in step 2 is 10 to 25 μm, the particle size D50 of the powder obtained by grinding is 2 to 7 μm, and the density of the base material is 4.0 to 4.5 g / cm³. 3 That is the case.

[0013] In one embodiment, the sintering temperature in step 3 is 950 to 1100°C, and the holding time is 4 to 15 hours. [Effects of the Invention]

[0014] The R-Fe-B rare earth magnetic material produced by the method according to the present invention is completely free of heavy rare earth elements, and possesses excellent performance while keeping manufacturing costs down. The uniformity and prismaticity of the magnetic material are both good, at 0.97 or higher.

[0015] Both the coercivity and remanent magnetic flux density of a magnetic material can be controlled by designing the grain boundary phase. The presence of a thick nonmagnetic phase at the grain boundary provides excellent magnetic isolation, improving the coercivity of the magnetic material.

[0016] Furthermore, the uniform presence of needle-shaped ZrBx or TiBx phases in the grain boundary phase provides excellent heat resistance. [Brief explanation of the drawing]

[0017] [Figure 1] This figure shows a cross-section of the R-Fe-B rare earth magnetic material of the present invention. [Figure 2] This figure shows the microstructure of the R-Fe-B rare-earth magnetic material of the present invention. Reference numeral 1 denotes the main phase particle (1), reference numeral 2 denotes the main phase particle (2), reference numeral 3 denotes the triangular region grain boundary phase (3), reference numeral 4 denotes the two-particle grain boundary phase, and reference numeral 5 denotes the acicular ZrBx phase and TiBx phase. [Figure 3] This figure shows the R-Fe-B rare-earth magnetic material of the present invention as captured by a scanning electron microscope. Reference numeral 1 denotes the main phase, reference numeral 2 denotes the HCP-type R-rich phase, reference numeral 3 denotes the RCu phase, reference numeral 4 denotes the RGa phase, reference numeral 5 denotes the R-rich phase, reference numeral 6 denotes the 6:14 phase, and reference numeral 7 denotes the acicular ZrBx phase and acicular TiBx phase. [Figure 4] This is the structural diffraction pattern of the FCC-type NdOx phase of the first embodiment of the present invention. [Figure 5] This is the structural diffraction pattern of the 6:14 phase of the first embodiment of the present invention. [Figure 6] This is the structural diffraction pattern of the R-rich phase of the first embodiment of the present invention. [Figure 7] This figure shows the two-particle grain boundary phase and its position according to the first embodiment of the present invention. [Modes for carrying out the invention]

[0018] The R-Fe-B sintered permanent magnetic material according to the present invention comprises a main phase, a two-particle grain boundary phase, and a triangular region grain boundary phase.

[0019] The main phase is R2T 14 It is phase B, R is a light rare earth element containing at least one of Nd, Pr, La, Ce, and Sm, T is Fe or Fe and Co, and R2T 14 Phase B consists of single-phase particles, and the average particle size of the main phase particles is in the range of 2 μm to 7 μm.

[0020] The two-particle grain boundary phase contains a grain boundary Cu-rich phase. In this grain boundary Cu-rich phase, the atomic percentage ratio of Cu to R within 20 nm near the interface of the main phase is 1.5 < Cu / R ≤ 2.0, and the atomic percentage ratio of Cu to R within 50 nm at the center of the grain boundary is 1.0 < Cu / R ≤ 1.5.

[0021] Also, the two-particle grain boundary phase contains a needle-like ZrBx phase with a width of 10 - 50 nm, and further contains a needle-like TiBx phase with a width of 10 - 50 nm.

[0022] The triangular region grain boundary phase contains an FCC-type NdOx phase, a RCu phase with an atomic percentage ratio of Cu to R of 1.0 < Cu / R ≤ 1.5, a RGa phase with an atomic percentage ratio of Ga to R of 0.1 < Ga / R ≤ 0.5, a R-rich phase of a 6:14 phase (hereinafter simply referred to as the 6:14 phase) with an atomic percentage ratio of R to the transition metal element of 6:14, and / or a needle-like ZrBx phase with a width of 10 - 50 nm, and / or a needle-like TiBx phase with a width of 10 - 50 nm.

[0023] Specifically, the needle-like ZrBx phase or TiBx phase refers to the ZrB phase or ZrB2 phase.

[0024] The main phase, the two-particle grain boundary phase, and the triangular region grain boundary phase exist in a cross-section perpendicular to the C-axis or parallel to the C-axis, and the phases in the triangular region grain boundary phase exist in the two-particle grain boundary phase.

[0025] Hereinafter, specific examples will be described. Example 1 The content of each component of the R-Fe-B-based rare earth magnetic material prepared as Example 1 is as follows in mass percentage representation.

[0026] PrNd: 29.5%, Al: 0.05%, B: 0.94%, Co: 0.5%, Cu: 0.4%, Ga: 0.3%, Zr: 0.3%, and the rest is Fe and impurities. Among the impurities, C was 800 ppm, N was 900 ppm, and O was 700 ppm.

[0027] The manufacturing method of the above R-Fe-B-based rare earth magnetic material is as follows. (Step 1) Hydrogen adsorption and dehydrogenation treatments were performed on the alloy made of each mixed raw material, and alloy flakes of the main phase alloy after the hydrogenation treatment and alloy flakes of the related phases were prepared respectively. The alloy flakes of the related phases are alloy flakes that become the Cu-rich phase and the Ga-rich phase. The alloy components are Pr 60 Nd 10 Cu 20 Ga 10 (atomic percentage ratio) melt-spun amorphous alloy flakes.

[0028] The hydrogen adsorption and dehydrogenation treatments may be low-temperature hydrogen adsorption treatment or high-temperature hydrogen adsorption treatment. The dehydrogenation treatment may be high-temperature dehydrogenation or low-temperature long-time dehydrogenation.

[0029] In this example, low-temperature hydrogen adsorption treatment was selected, the hydrogen adsorption temperature was 50 °C, the dehydrogenation temperature was 500 °C, and the dehydrogenation time was 7 hours.

[0030] (Step 2) Alloy flakes of the related phases were mixed with the alloy flakes of the main phase alloy created in Step 1 at a ratio of 2% by mass and crushed mechanically. By dissolving the amorphous alloy and the main alloy, the Cu-rich phase and the Ga-rich phase in the grain boundaries were drawn in, and the mixed powder material was crushed to make powder. The powder-making process may be ball-mill powder-making or jet-mill powder-making. In the powder-making process, a small amount of oxygen gas was introduced for magnetic field forming to create a substrate. Magnetic field forming is a forming method in which a magnetic field is oriented while performing press forming. After magnetic field forming, it may be subjected to hydrostatic pressing, followed by sintering and aging treatment, or directly subjected to sintering and aging treatment.

[0031] The alloy flakes of the main phase alloy and the alloy flakes of the related phases were mechanically crushed to a particle size of 20 μm, and further the particle size of the powder was made D50 = 3.5 μm using a jet mill. The density of the substrate after magnetic field forming was 4.5 g / cm 3 as.

[0032] (Step 3) The base material was placed in a sintering furnace, sintered, and then kept warm. The resulting blank was subjected to a two-stage aging treatment to create an R-Fe-B rare-earth magnetic material.

[0033] The sintering temperature was 1070°C and the holding time was 6 hours. For the two-stage aging treatment, the first aging temperature was 800°C and the holding time was 3 hours, and the second aging temperature was 430°C and the holding time was 7 hours.

[0034] The sintering process primarily increases density and improves contact characteristics between powder particles, thereby increasing the residual magnetic flux density and strength of the magnetic material and giving it microstructural characteristics of high permanent magnetic properties. This process is mainly liquid-phase sintering, and high-melting-point plates such as molybdenum plates or zirconium plates can be used as pads during the sintering process of the magnetic material. Mixed gases such as nitrogen, argon, and hydrogen can also be introduced to prevent oxidation of the magnetic material.

[0035] In the first stage of the aging process, the R-rich phase is melted to form a flowing liquid phase, which flows and permeates into the gaps between powder particles, melting and rounding the corners of the crystal grains and optimizing the distribution of the grain boundary phase of the magnetic material. Furthermore, in this process, alloy powder can be added to the surface of the blank, and the alloy powder may be PrNd, PrAl, PrAlCu, PrAlGa, etc.

[0036] In the second stage of the aging treatment, the grain boundary phase was liquefied and allowed to flow through the crystal grains, dispersing uniformly within the grain boundaries to obtain regular and smooth grain boundaries. This suppressed the generation of magnetization reversal nuclei and reduced the demagnetizing field in the main phase crystal grains. After heat treatment, the distribution of the R-rich phase at the grain boundaries became sufficiently uniform, better isolating the main phase crystal grains and eliminating the magnetic exchange coupling between crystal grains, thereby increasing the coercivity of the magnetic material. After the aging treatment, both Hcj(BH)m and angularity improved. Here, mainly by forming Cu-rich phases, Ga-rich phases, and 6:14 phases, the low-melting-point phase was further distributed more uniformly and continuously in the grain boundary phase, and the Cu-rich phases, Ga-rich phases, and 6:14 phases were better isolated from the main phase, resulting in a magnetic coupling removal effect and further increasing the coercivity and angularity of the magnetic material.

[0037] Under room temperature conditions of 20°C, the R-Fe-B rare-earth magnetic material obtained in Example 1 was measured using the NIM-2000 magnetic property measurement system. The measurement results for remanent magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) are shown in Table 1. The atomic percentages of each element in the two-particle grain boundary phase and triangular region grain boundary phase of the R-Fe-B rare-earth magnetic material are shown in Table 3.

[0038] Example 2 The content of each component of the R-Fe-B rare earth magnetic material prepared as Example 2 is as follows, expressed as a mass percentage.

[0039] The composition is as follows: PrNd: 32.0%, Al: 0.05%, B: 0.85%, Co: 2.0%, Cu: 2.0%, Ga: 0.4%, Zr: 0.5%, with the remainder being Fe and impurities. Among the impurities, C was present at 700 ppm, N at 600 ppm, and O at 800 ppm.

[0040] The differences between Example 2 and Example 1 lie in the manufacturing method, as well as slight differences in the ratio of each component. The manufacturing process for Example 2 is as follows.

[0041] (Step 1) Alloys prepared from each mixed raw material were subjected to hydrogen adsorption and dehydrogenation treatments, and alloy thin sheets of the main phase alloy and alloy thin sheets of the related phases were prepared after hydrogenation treatment. The alloy thin sheets of the related phases were Cu-rich and Ga-rich, and the alloy component was Pr 40 Nd 10 Cu 40 Ga 10 This is a melt-spun amorphous alloy thin film (atomic percentage ratio).

[0042] Hydrogen adsorption and dehydrogenation treatment may be performed using low-temperature hydrogen adsorption treatment or high-temperature hydrogen adsorption treatment. Dehydrogenation treatment may be performed using high-temperature dehydrogenation or low-temperature dehydrogenation for a long time.

[0043] In this example, room temperature hydrogen adsorption treatment was selected, with a dehydrogenation temperature of 600°C and a dehydrogenation time of 5 hours.

[0044] (Step 2) In Step 1, alloy thin sheets of the main phase alloy prepared were mixed with alloy thin sheets of the related phase at a mass ratio of 4%, and the mixture was crushed mechanically. By solid-solving the amorphous alloy with the main alloy, the Cu-rich and Ga-rich phases in the grain boundaries were drawn in, and the mixed powder material was crushed to produce flour. The flour milling process may be done by ball milling or jet milling. In the flour milling process, a small amount of oxygen gas was introduced and the base material was formed by magnetic field molding. Magnetic field molding is a molding method in which a magnetic field is oriented while press molding. After magnetic field molding, hydrostatic pressing may be performed followed by sintering and aging treatment, or sintering and aging treatment may be performed directly.

[0045] The alloy flakes of the main phase alloy and related phases were mechanically ground to a particle size of 15 μm. Further, a jet mill was used to obtain a powder with a particle size D50 of 3.5 μm, and the density of the substrate after magnetic field shaping was 4.2 g / cm³. 3 That's what I decided.

[0046] (Step 3) The base material was placed in a sintering furnace, sintered, and then kept warm. The resulting blank was subjected to a two-stage aging treatment to create an R-Fe-B rare-earth magnetic material.

[0047] The sintering temperature was 1100°C and the holding time was 4 hours. For the two-stage aging treatment, the first aging temperature was 850°C and the holding time was 3 hours, and the second aging temperature was 450°C and the holding time was 3 hours.

[0048] Under room temperature conditions of 20°C, the Rd-Fe-B rare-earth magnetic material obtained in Example 2 was measured using the NIM-2000 magnetic property measurement system. The measurement results for remanent magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) are shown in Table 1. The atomic percentages of each element in the two-particle grain boundary phase and triangular region grain boundary phase of the R-Fe-B rare-earth magnetic material are shown in Table 3.

[0049] Example 3 The content of each component of the R-Fe-B rare earth magnetic material prepared as Example 3 is as follows, expressed as a mass percentage.

[0050] The composition is as follows: PrNd: 30.8%, Al: 0.07%, B: 0.90%, Co: 0.50%, Cu: 0.70%, Ga: 0.60%, Zr: 0.30%, with the remainder being Fe and impurities. Among the impurities, C was present at 650 ppm, N at 700 ppm, and O at 700 ppm.

[0051] The differences between Example 3 and Examples 1 and 2 lie in the manufacturing method, as well as slight differences in the ratio of each component. The manufacturing process for Example 3 is as follows.

[0052] (Step 1) Alloys prepared from each mixed raw material were subjected to hydrogen adsorption and dehydrogenation treatments, and alloy thin sheets of the main phase alloy and alloy thin sheets of the related phases were prepared after hydrogenation treatment. The alloy thin sheets of the related phases were Cu-rich and Ga-rich, and the alloy component was Pr 50 Nd 10 Cu 20 Ga 10 This is a melt-spun amorphous alloy thin film (atomic percentage ratio).

[0053] Hydrogen adsorption and dehydrogenation treatment may be performed using low-temperature hydrogen adsorption treatment or high-temperature hydrogen adsorption treatment. Dehydrogenation treatment may be performed using high-temperature dehydrogenation or low-temperature dehydrogenation for a long time.

[0054] In this example, low-temperature hydrogen adsorption treatment was selected. The hydrogen adsorption temperature was 100°C, the dehydrogenation temperature was 550°C, and the dehydrogenation time was 4 hours.

[0055] (Step 2) In Step 1, alloy slabs of the main phase alloy prepared were mixed with alloy slabs of the related phase at a mass ratio of 3%, and the mixture was crushed mechanically. By solid-solving the amorphous alloy with the main alloy, the Cu-rich and Ga-rich phases in the grain boundaries were drawn in, and the mixed powder material was crushed to produce flour. The flour milling process may be done by ball milling or jet milling. In the flour milling process, a small amount of oxygen gas was introduced and the base material was formed by magnetic field molding. Magnetic field molding is a molding method in which a magnetic field is oriented while press molding. After magnetic field molding, hydrostatic pressing may be performed followed by sintering and aging treatment, or sintering and aging treatment may be performed directly.

[0056] The alloy flakes of the main phase alloy and related phases were mechanically ground to a particle size of 10 μm. Further, a jet mill was used to obtain a powder with a particle size D50 of 3.8 μm, and the density of the substrate after magnetic field molding was 4.0 g / cm³. 3 That's what I decided.

[0057] (Step 3) The base material was placed in a sintering furnace, sintered, and then kept warm. The resulting blank was subjected to a two-stage aging treatment to create an R-Fe-B rare-earth magnetic material.

[0058] The sintering temperature was 1050°C and the holding time was 6 hours. For the two-stage aging treatment, the first aging temperature was 750°C and the holding time was 5 hours, and the second aging temperature was 480°C and the holding time was 4 hours.

[0059] Under room temperature conditions of 20°C, the R-Fe-B rare-earth magnetic material obtained in Example 3 was measured using the NIM-2000 magnetic property measurement system. The measurement results for remanent magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) are shown in Table 1. The atomic percentages of each element in the two-particle grain boundary phase and triangular region grain boundary phase of the R-Fe-B rare-earth magnetic material are shown in Table 3.

[0060] Example 4 The content of each component of the R-Fe-B rare earth magnetic material prepared as Example 4 is as follows, expressed as a mass percentage.

[0061] The composition is as follows: PrNd: 31.5%, Al: 0.35%, B: 0.94%, Co: 0.8%, Cu: 1.2%, Ga: 0.3%, Ti: 0.15%, with the remainder being Fe and impurities. Among the impurities, C was present at 750 ppm, N at 500 ppm, and O at 700 ppm.

[0062] The differences between Example 4 and Examples 1-3 lie in the manufacturing method, as well as slight differences in the ratio of each component. The manufacturing process for Example 4 is as follows.

[0063] (Step 1) Alloys prepared from each mixed raw material were subjected to hydrogen adsorption and dehydrogenation treatments, and alloy flakes of the main phase alloy and related phases were prepared after hydrogenation treatment. The related phase alloy flakes were Cu-rich phase, Ga-rich phase, etc., and the alloy composition was Pr 40 Nd 10 Cu 40 Ga 10 This is a melt-spun amorphous alloy thin film (atomic percentage ratio).

[0064] Hydrogen adsorption and dehydrogenation treatment may be performed using low-temperature hydrogen adsorption treatment or high-temperature hydrogen adsorption treatment. Dehydrogenation treatment may be performed using high-temperature dehydrogenation or low-temperature dehydrogenation for a long time.

[0065] In this example, room temperature hydrogen adsorption treatment was selected. The dehydrogenation temperature was 650°C and the dehydrogenation time was 3 hours.

[0066] (Step 2) In Step 1, alloy slabs of the main phase alloy prepared were mixed with alloy slabs of the related phase at a mass ratio of 3.5%, and the mixture was crushed mechanically. By solid-solving the amorphous alloy with the main alloy, the Cu-rich and Ga-rich phases in the grain boundaries were drawn in, and the mixed powder material was crushed to produce flour. The flour milling process may be done by ball milling or jet milling. In the flour milling process, a small amount of oxygen gas was introduced and the base material was formed by magnetic field molding. Magnetic field molding is a molding method in which a magnetic field is oriented while press molding. After magnetic field molding, hydrostatic pressing may be performed followed by sintering and aging treatment, or sintering and aging treatment may be performed directly.

[0067] The alloy flakes of the main phase alloy and related phases were mechanically ground to a particle size of 25 μm. Further, a jet mill was used to obtain a powder with a particle size D50 of 4.0 μm, and the density of the substrate after magnetic field molding was 4.4 g / cm³. 3 That's what I decided.

[0068] (Step 3) The base material was placed in a sintering furnace, sintered, and then kept warm. The resulting blank was subjected to a two-stage aging treatment to create an R-Fe-B rare-earth magnetic material.

[0069] The sintering temperature was 1060°C and the holding time was 8 hours. For the two-stage aging treatment, the first aging temperature was 700°C and the holding time was 7 hours, and the second aging temperature was 500°C and the holding time was 3 hours.

[0070] Under room temperature conditions of 20°C, the R-Fe-B rare-earth magnetic material obtained in Example 4 was measured using the NIM-2000 magnetic property measurement system. The measurement results for remanent magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) are shown in Table 1. The atomic percentages of each element in the two-particle grain boundary phase and triangular region boundary phase of the R-Fe-B rare-earth magnetic material are shown in Table 3.

[0071] Example 5 The content of each component of the R-Fe-B rare earth magnetic material prepared as Example 5 is as follows, expressed as a mass percentage.

[0072] The composition is as follows: PrNd: 33.5%, Al: 0.80%, B: 1.05%, Co: 1.50%, Cu: 0.8%, Ga: 0.20%, Ti: 0.10%, with the remainder being Fe and impurities. Among the impurities, C was present at 500 ppm, N at 400 ppm, and O at 400 ppm.

[0073] The differences between Example 5 and Examples 1-4 lie in the manufacturing method, as well as slight differences in the ratio of each component. The manufacturing process for Example 5 is as follows.

[0074] (Step 1) The alloys prepared from each of the blended raw materials were subjected to hydrogen adsorption and dehydrogenation treatments, and alloy flakes of the main phase alloy and related phases were prepared after the hydrogenation treatment. The alloy flakes of the related phases were alloy flakes of Cu-rich phase, Ga-rich phase, etc., and the alloy component was Pr 50 Nd 10 Cu 40 Melt-spun amorphous alloy thin sheets (atomic percentage ratio) were used.

[0075] Hydrogen adsorption and dehydrogenation treatment may be performed using low-temperature hydrogen adsorption treatment or high-temperature hydrogen adsorption treatment. Dehydrogenation treatment may be performed using high-temperature dehydrogenation or low-temperature dehydrogenation for a long time.

[0076] In this example, room temperature hydrogen adsorption treatment was selected. The dehydrogenation temperature was 500°C and the dehydrogenation time was 7 hours.

[0077] (Step 2) In Step 1, alloy slabs of the main phase alloy prepared were mixed with alloy slabs of the related phase at a mass ratio of 2.0%, and the mixture was crushed mechanically. By solid-solving the amorphous alloy with the main alloy, the Cu-rich phase in the grain boundaries was drawn in, and the mixed powder material was crushed to produce flour. The flour milling process may be done by ball milling or jet milling. In the flour milling process, a small amount of oxygen gas was introduced and the base material was formed by magnetic field molding. Magnetic field molding is a molding method in which a magnetic field is oriented while press molding. After magnetic field molding, hydrostatic pressing may be performed followed by sintering and aging treatment, or sintering and aging treatment may be performed directly.

[0078] The alloy flakes of the main phase alloy and related phases were mechanically ground to a particle size of 10 μm. Further, a jet mill was used to obtain a powder with a particle size D50 of 4.5 μm, and the density of the substrate after magnetic field shaping was 4.1 g / cm³. 3 That's what I decided.

[0079] (Step 3) The base material was placed in a sintering furnace, sintered, and then kept warm. The resulting blank was subjected to a two-stage aging treatment to create an R-Fe-B rare-earth magnetic material.

[0080] The sintering temperature was 950°C and the holding time was 15 hours. For the two-stage aging treatment, the first aging temperature was 900°C and the holding time was 4 hours, and the second aging temperature was 550°C and the holding time was 3 hours.

[0081] Under room temperature conditions of 20°C, the R-Fe-B rare-earth magnetic material obtained in Example 5 was measured using the NIM-2000 magnetic property measurement system. The measurement results for remanent magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) are shown in Table 1. The atomic percentages of each element in the two-particle grain boundary phase and triangular region grain boundary phase of the Nd-Fe-B rare-earth magnetic material are shown in Table 3.

[0082] Comparative Example To evaluate the characteristics of Examples 1 to 5 of the present invention described above, Comparative Examples 1 to 5 were prepared. Comparative Examples 1 to 5 were prepared using the same raw materials and methods as Examples 1 to 5, except that the second aging temperature in the two-stage aging treatment was different. The R-Fe-B rare earth magnetic materials of Comparative Examples 1 to 5 were measured using the NIM-2000 magnetic property measurement system. The measurement results for residual magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) are shown in Table 2.

[0083] Table 1: Data on remanent magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) JPEG2026100813000002.jpg91135

[0084] Table 2: Data on remanent magnetic flux density (Br), coercivity (Hcj), and aspect ratio (Hk / Hcj) JPEG2026100813000003.jpg79135

[0085] A comparison of the measured values ​​shown in Tables 1 and 2 revealed that the temperature during the aging process has a significant impact on diffusion performance. Specifically, the following is observed:

[0086] The manufacturing process parameters for Example 1 were a sintering time of 1070°C and 6 hours, a first aging temperature of 880°C and 3 hours, a second aging temperature of 430°C and 7 hours, and the magnetic properties of the magnetic material were Br=14.4kGs, Hcj=19.0kOe, and Hk / Hcj=0.98. Compared with Comparative Example 1, while the remanent magnetic flux density (Br) and square aspect ratio (Hk / Hcj) are basically the same, there is a significant difference of 2kOe in coercivity (Hcj). This suggests that the aging treatment temperature has a large influence on the coercivity performance of the magnetic material.

[0087] In Example 2, the process parameters were sintering time at 1100°C for 4 hours, first aging temperature at 850°C for 3 hours, and second aging temperature at 450°C for 3 hours. The magnetic properties of the magnetic material were Br = 13.5 kGs, Hcj = 24.5 kOe, and Hk / Hcj = 0.98. Compared with Comparative Example 2, while the remanent magnetic flux density (Br) and square aspect ratio (Hk / Hcj) are basically the same, there is a significant difference of 2 kOe in coercivity (Hcj). This suggests that the aging temperature has a large influence on the coercivity performance of the magnetic material.

[0088] In Example 3, the process parameters were sintering time at 1050°C for 5 hours, first aging temperature at 750°C for 5 hours, second aging temperature at 480°C for 4 hours, and the magnetic properties of the magnetic material were Br=14.0kGs, Hcj=22.0kOe, and Hk / Hcj=0.98. Compared with Comparative Example 3, while the residual magnetic flux density (Br) and square aspect ratio (Hk / Hcj) are basically the same, there is a significant difference of 3kOe in coercivity (Hcj). This suggests that the aging temperature has a large influence on the coercivity performance of the magnetic material.

[0089] The process parameters for Example 4 were a sintering time of 1060°C and 8 hours, a first aging temperature of 700°C and 7 hours, a second aging temperature of 500°C and 3 hours, and the magnetic properties of the magnetic material were Br=13.2kGs, Hcj=23.5kOe, and Hk / Hcj=0.98. Compared with Comparative Example 4, while the remanent magnetic flux density (Br) and square aspect ratio (Hk / Hcj) are basically the same, there is a significant difference of 3kOe in coercivity (Hcj). This suggests that the aging temperature has a large influence on the coercivity performance of the magnetic material.

[0090] The process parameters for Example 5 were a sintering time of 950°C and 15 hours, a first aging temperature of 900°C and 4 hours, a second aging temperature of 550°C and 3 hours, and the magnetic properties of the magnetic material were Br=13.2kGs, Hcj=26.0kOe, and Hk / Hcj=0.98. Compared with Comparative Example 5, although the residual magnetic flux density (Br) and square aspect ratio (Hk / Hcj) are basically the same, there is a significant difference of 4kOe in coercivity (Hcj). This suggests that the aging temperature has a large influence on the coercivity performance of the magnetic material.

[0091] Table 3: Atomic percentages of each element in R-Fe-B rare earth magnetic materials JPEG2026100813000004.jpg65135

[0092] As shown in Table 3, the R-Fe-B rare-earth magnetic materials of Examples 1 to 5 are mainly composed of a main phase, a two-particle grain boundary phase, and a triangular region grain boundary phase. Here, the main phase is basically the same except for slight differences due to differences in the mixing ratio of PrNd or pure Nd in R. As shown in the observation results for Example 1 in Figures 2 to 7, the positions and ranges of the two-particle grain boundary phase and the triangular region grain boundary phase are clear, and the main phase, two-particle grain boundary phase, and triangular region grain boundary phase can be defined as follows.

[0093] The average atomic percentage ratios of Cu / R within 20 nm of the main phase of the two-particle grain boundary phase of the R-Fe-B rare-earth magnetic material in Example 1 were 1.6, 2.0, 1.7, 1.8, and 1.6, respectively. The average atomic percentage ratios of Cu / R within 50 nm of the grain boundary center of the two-particle grain boundary phase were 1.1, 1.5, 1.2, 1.2, and 1.2, respectively. Needle-shaped high-melting-point phases were present in all two-particle grain boundary phases of Example 1, with widths of 15, 50, 20, 15, and 10 nm, respectively.

[0094] In the R-Fe-B rare-earth magnetic material of Example 1, both the two-particle grain boundary phase and the triangular region grain boundary phase contained R-rich phase, FCC-type Nd-rich phase, and 6:14 phase. The average atomic percentage ratios of Cu / R in the RCu phase were 1.05, 1.50, 1.10, 1.30, and 1.10, and the average atomic percentage ratios of Ga / R in the RGa phase were 0.15, 0.20, 0.50, 0.15, and 0.15. In the triangular region grain boundary phase of Example 1, needle-shaped high-melting-point phases were present in all cases, with widths of 15, 50, 20, 15, and 10 nm, respectively. Although the measurements using the NIM-2000 magnetic property measurement system were performed on the magnetic material of Example 1, it can be theoretically assumed that Examples 2 to 5 also have similar crystal structures.

[0095] As described above, the d-Fe-B rare earth magnetic material obtained by the manufacturing method of the present invention does not contain heavy rare earth elements, has good uniformity, and exhibits an excellent square-to-size ratio. The coercivity and remanent magnetic flux density of the Nd-Fe-B rare earth magnetic material can both be controlled by the design of the grain boundary phase. The presence of a thick non-magnetic phase in the grain boundary phase provides excellent magnetic isolation, improving the coercivity of the magnetic material. The uniform presence of needle-shaped ZrBx or TiBx phase in the grain boundary phase provides excellent heat resistance.

[0096] Of the magnetic materials prepared in each embodiment of the present invention, the observation results and characteristics of the magnetic material according to Example 1 will be explained again below. As shown in Figure 1, arbitrary cross-sections of the interface of the R-Fe-B magnetic material according to Example 1 were observed and measured. The specific observation and measurement methods included observation of the microstructure using a scanning electron microscope (SEM) and observation of the microstructure of grain boundaries and the distribution of trace elements using a transmission electron microscope (TEM). Although the cross-section shown in Figure 1 is rectangular, the actual shape is not limited to a rectangle and may be square, tile-shaped, bun-shaped, cylindrical, arc-shaped, etc.

[0097] As shown in Figure 2, the R-Fe-B magnetic material has main phase particles (1) and main phase particles (2), and the components of main phase particles (1) and main phase particles (2) are the same. In addition to those shown in Examples 1 to 5, the main phase is La2Fe 14 B, Ce2Fe 14 It may also be B, etc. In addition to the elements mentioned above, the main phase particles may also contain elements such as Al, Ga, Cu, etc. Note that the main phase corresponds to reference numeral 1 in Figure 3.

[0098] The triangular region grain boundary phase (3) in Figure 2 is the phase excluding the main phase and the two-particle grain boundary phase. The triangular region grain boundary phase (3) is an FCC phase containing ROx, and the triangular region grain boundary phase contains ROx, which is an FCC phase, in a predetermined ratio. The misfit at the interface between two adjacent phases in the crystal orientation

[0011] and crystal plane (111) of the main phase and the FCC-type ROx phase is small at 4.2%, which increases the magnetic isolation coefficient of the grain boundaries of the R-Fe-B rare earth magnetic material, and as a result the coercivity of the magnetic material is improved. The oxygen content of this phase is usually 10 at% to 45 at%. The R-Fe triangular region grain boundary phase may further contain a DHCP-type R-rich phase and an HCP-type R-rich phase. The DHCP-type R-rich phase is abundant in the two-particle grain boundary phase or the triangular region grain boundary phase. The reason for the abundance of the HCP-type R-rich phase is that it is generated due to inadequate control of the oxygen content or an increase in oxygen content due to long-term storage. The HCP-type R-rich phase content is very low, accounting for less than 15% of the triangular grain boundary phase. This phase corresponds to reference numeral 2 in Figure 3.

[0099] The triangular region grain boundary phase (3) contains an RCu phase, and the RCu phase may be a PrCu phase, a NdCu phase, or a NdPrCu phase, etc. The atomic percentage ratio of the RCu phase may be 1.0 < Cu / R ≦ 1.5. Also, the R-Fe triangular region grain boundary phase may further contain phases such as LaCu and CeCu. The RCu phase occupies 15% - 50% of the triangular region grain boundary phase. The RCu phase corresponds to reference numeral 3 in Figure 3.

[0100] The triangular region grain boundary phase (3) in Figure 2 contains an RGa phase. The RGa phase may be a PrGa phase, a NdGa phase, or a NdPrGa phase, etc. The atomic percentage ratio of the RGa phase may be 0.1 < Ga / R ≦ 0.5. Further, the triangular region grain boundary phase may contain phases such as LaGa and CeGa. Further, the RGa phase may contain a composite phase of Co, Al, Cu, Fe, Mg, Zn, Mo, Ti, Zr, Sb, Bi, etc. The RGa phase is 10% or less of the triangular region, and the RGa phase corresponds to reference numeral 4 in Figure 3.

[0101] The triangular region grain boundary phase (3) in Figure 2 contains an R-rich phase. The R-rich phase is mainly dominated by the occupancy ratio of R. Further, the R element in this triangular region grain boundary phase may contain rare earth elements such as La, Ce, Sm, etc. At the same time, the R-rich phase may contain a certain amount of Fe, and its content is 0 at% - 20 at%. Also, the R-rich phase may contain elements such as Co, Al, Cu, Fe, Mg, Zn, Mo, Ti, Zr, Sb, Bi, etc., but the content of these elements is not constant and may be zero. The R-rich phase may exist in an amorphous state, and the occupancy ratio of the R-rich phase in the triangular region grain boundary phase is 35% or less. The R-rich phase corresponds to reference numeral 5 in Figure 3.

[0102] The triangular region grain boundary phase (3) in Figure 2 contains a phase with an atomic percentage ratio of R to transition metal element of 6:14. The 6:14 phase satisfies the proportional relationship of 6:14, which is mainly R6T 13It exists as the (GaCu)1 phase. The 6:14 phase may contain at least one of Al, Mg, Zn, Ti, Zr, Mo, Sb, and Bi. The T element refers to either Fe, or Fe and Co. R may include rare earth elements such as La, CE, and SM in addition to Pr and Nd, and at least four elements are present in the 6:14 phase. The 6:14 phase accounts for 20% to 50% of the triangular grain boundary phase. The 6:14 phase corresponds to the symbol 6 in Figure 3.

[0103] The triangular region grain boundary phase (3) in Figure 2 includes acicular ZrBx phases with a width of 10-50 nm and / or acicular TiBx phases with a width of 10-50 nm. The acicular ZrBx and acicular TiBx phases are high-melting-point phases, specifically ZrB2 and TiB2 phases. The high-melting-point phases may further include other high-melting-point phases such as Mo and W. The high-melting-point phases may sometimes exhibit point-like or linear shapes. The presence of these high-melting-point phases creates a pinning effect within the magnetic domain, resulting in high coercivity and heat resistance in the magnetic material. The acicular ZrBx and acicular TiBx phases are mainly distributed within the triangular region grain boundary phase or the two-particle grain boundary phase, along the grain boundaries or at the edges or center of the triangular region grain boundary phase. The high-melting-point phases are formed in the immediate vicinity of the R-rich phase, 6:14 phase, RCu phase, RGa phase, or ROx. The needle-shaped ZrBx and TiBx phases are diffusely distributed and are not fixed to any particular phase. The distribution locations of the needle-shaped ZrBx and TiBx phases are indicated by reference numeral 5 (black needle-shaped morphology) in Figure 2, which corresponds to reference numeral 7 in Figure 3.

[0104] The triangular region grain boundary phase (3) includes the FCC-type ROx phase, RCu phase, RGa phase, 6:14 phase, or acicular ZrBx phase or TiBx phase. These phases do not necessarily exist individually, but may also exist intersecting within a single triangular region. Examples of two-phase intersections include RCu phase and FCC-type ROx phase, RGa phase and FCC-type ROx phase, R-rich phase and FCC-type ROx phase, R-rich phase and 6:14 phase, 6:14 phase and acicular ZrBx phase, and 6:14 phase and acicular TiBx phase. Examples of three-phase intersections include the R-rich phase, 6:14 phase and acicular ZrBx phase or acicular TiBx phase. A single triangular region may also exist, such as the FCC-type ROx phase or 6:14 phase. Figure 4 shows the diffraction pattern of the FCC-type ROx phase structure measured by transmission electron microscopy, confirming the presence of the FCC-type ROx phase. Figure 5 shows the diffraction pattern of the 6:14 phase structure, confirming the presence of the 6:14 phase. Figure 6 shows the diffraction pattern of the R-rich phase structure, confirming the presence of the R-rich phase.

[0105] FIG. 7 is a diagram showing a two-particle grain boundary and its position photographed with a transmission electron microscope. At the position a in FIG. 7, the grain boundary Cu-rich phase is located within 20 nm from the vicinity of the main-phase interface, the atomic percentage ratio of Cu to R is 1.5 < Cu / R ≤ 2.0, and a is equal to 20 nm. At the position b in FIG. 7, the grain boundary Cu-rich phase is located at a position 50 nm from the center line of the grain boundary, and the atomic percentage ratio of Cu to R satisfies 1.0 < Cu / R ≤ 1.5, and b is equal to 50 nm. The two-particle grain boundary refers to the grain boundary between two main-phase particles. The width of the grain boundary ranges from 2 nm to 500 nm, and the average width is 100 nm. There is an obvious regularity in the composition distribution of the grain boundary. At the edge part, the content of Cu increases significantly, and the content of Cu is higher than that at other positions of the grain boundary. At a position within 20 nm from the vicinity of the main-phase interface where the content of Cu increases, the content of R decreases significantly. On the other hand, it cannot be denied that there may be the same regularity at a position further away from the vicinity of the main-phase interface, for example, at 30 nm from the interface. The contents of Cu and R are much higher than those of other elements in the grain boundary. Furthermore, it is speculated that rare earth elements such as La, Ce, and Sm other than Pr and Nd used in the examples also have the same regularity. Furthermore, it is speculated that elements such as Al, Cu, Mg, and Zn also have the same regularity as Cu. That is, the contents of these elements at the edge part increase significantly compared with other parts of the grain boundary. Within the range of the a value near the main-phase interface, the content of R elements such as Pr and Nd decreases significantly. On the other hand, it cannot be denied that there may be the same regularity at a position further away from the vicinity of the main-phase interface, for example, at a position within the range of a from the interface. The Cu-rich phase in the grain boundary is within the center line b of the grain boundary, and the atomic percentage ratio of Cu to R satisfies 1.0 < Cu / R ≤ 1.5. Similarly, at this position, the length of the center line may be longer, and the contents of Cu and R are much higher than the distribution of other elements. The content of Cu is more than that of R, and the atomic percentage ratio of Cu to R satisfies 1.0 < Cu / R ≤ 1.5. R is mainly Pr and Nd, but it may also be La, Ce, Sm, etc. in addition. Furthermore, elements such as Al, Cu, Mg, and Zn may also have the same regularity as Cu.

[0106] The needle-shaped ZrBx and needle-shaped TiBx phases have widths of 10 to 50 nm. These phases are distributed within the two-particle grain boundary phase and may be high-melting-point phases distributed along the grain boundaries, or they may be needle-shaped structures distributed perpendicular to the grain boundaries. Since Zr and Ti are high-melting-point elements, the high-melting-point phase may further contain other high-melting-point phases such as Mo and W. Furthermore, the high-melting-point phase may have shapes such as points or lines. The presence of this high-melting-point phase provides a pinning effect on the magnetic domains, giving the magnetic material high coercivity and heat resistance. This phase is located in the two-particle grain boundary phase and corresponds to reference numeral 7 in Figure 3.

[0107] Nd-Fe-B rare-earth magnetic materials possess high magnetic properties such as high coercivity, high remanent magnetic flux density, and prismatic shape. The composition of this magnetic material is as follows:

[0108] The two-particle grain boundary phase is located within 20 nm of the interface of the main phase and satisfies one of the following conditions (all of which represent the atomic percentage ratio of Cu to R). 1.5 <Cu / R≦1.6、 1.5 <Cu / R≦1.7、 1.5 <Cu / R≦1.8、 1.5 <Cu / R≦1.9、 1.5 <Cu / R≦2.0、 1.6 ≤ Cu / R ≤ 1.7, 1.6 ≤ Cu / R ≤ 1.8, 1.6 ≤ Cu / R ≤ 1.9, 1.6 ≤ Cu / R ≤ 2.0, 1.7 ≤ Cu / R ≤ 1.8, 1.7 ≤ Cu / R ≤ 1.9, 1.7 ≤ Cu / R ≤ 2.0.

[0109] At the center of the two-particle grain boundary phase, the Cu-rich phase of the grain boundary is within 100 nm of the center of the grain boundary, and the RCu phase within the triangular region satisfies one of the following conditions. (Both indicate the atomic percentage ratio of Cu to R.) 1.0 <Cu / R≦1.2、 1.0 <Cu / R≦1.3、 1.0 <Cu / R≦1.4、 1.0 <Cu / R≦1.5、 1.1 <Cu / R≦1.2、 1.1 <Cu / R≦1.3、 1.1 <Cu / R≦1.4、 1.1 <Cu / R≦1.5、 1.2 <Cu / R≦1.3、 1.2 <Cu / R≦1.4、 1.2 <Cu / R≦1.5、 1.3 <Cu / R≦1.4、 1.3 <Cu / R≦1.5、 1.4 <Cu / R≦1.5。

[0110] The RGa phase within the triangular region satisfies one of the following conditions (all of which represent the atomic percentage ratio of Ga to R): 0.1 <Ga / R≦0.2、 0.1 <Ga / R≦0.3、 0.1 <Ga / R≦0.4、 0.1 <Ga / R≦0.5、 0.2 <Ga / R≦0.3、 0.2 <Ga / R≦0.4、 0.2 <Ga / R≦0.5、 0.3 <Ga / R≦0.4、 0.3 <Ga / R≦0.5、 0.4 <Ga / R≦0.5。

[0111] The remanent magnetic flux density of R-Fe-B rare-earth magnetic materials is mainly influenced by factors such as the volume fraction of positive domains, the proportion of non-magnetic phases, relative density, and degree of orientation of magnetic domains. Therefore, these parameters can be compared in the process of designing magnetic materials with different remanent magnetic flux densities. Factors influencing the coercivity of R-Fe-B rare-earth magnetic materials include the anisotropic magnetic field, structural elements, demagnetizing elements, and saturation magnetic polarization intensity, while the angularity is mainly influenced by the uniformity of the magnetic material. Based on the above, the mechanism for improving the performance of R-Fe-B rare-earth magnetic materials can be analyzed as follows.

[0112] The triangular region grain boundary phase includes FCC-type ROx phase, RCu phase, RGa phase, 6:14 phase, or acicular ZrBx phase or acicular TiBx phase, while the two-particle grain boundary phase includes Cu-rich phase, R-rich phase, 6:14 phase, etc.

[0113] The triangular region grain boundary phase contains an FCC-type ROx phase, resulting in small misfits at the interface between two adjacent phases of this phase and the main phase. This improves the microstructure of the R-Fe-B rare-earth magnetic material, reduces the stray magnetic field formed after the misfit at the interface between two adjacent phases is reduced, and lowers the defect rate, thus improving the coercivity of the magnetic material. Since the triangular region grain boundary phase contains the non-magnetic phases RCu and RGa, it can magnetically isolate two or more adjacent main phase particles, thus improving the coercivity of the magnetic material. The triangular region grain boundary phase contains an FCC-type ROx phase, RCu phase, RGa phase, 6:14 phase, or acicular ZrBx phase or acicular TiBx phase. The weakly magnetic 6:14 phase included in the two-particle grain boundary phase and the Cu-rich phase of the two-particle grain boundary phase can both form a continuous and good grain boundary phase. Furthermore, a grain boundary phase formed to a predetermined thickness can effectively isolate the magnetism between main phase particles, thus improving the coercivity of the magnetic material. The needle-shaped ZrBx or TiBx phases in the triangular region grain boundary phase and the two-particle grain boundary phase prevent magnetic deflection of the positive particle magnetic domains and exert a pinning effect, thereby improving the coercivity of the magnetic material. The ratio of the triangular region grain boundary phase, the two-particle grain boundary phase, and the main phase is one of the factors that determine the magnitude of the remanent magnetic flux density. Furthermore, at the same magnetization strength, the FCC-type ROx phase, RCu phase, RGa phase, 6:14 phase, or needle-shaped ZrBx or needle-shaped TiBx phase contained in the triangular region grain boundary phase, and the Cu-rich phase in the two-particle grain boundary phase, all improve the remanent magnetic flux density of the magnetic material by effectively increasing the volume fraction of the positive magnetic domains. In addition, if no bright spots occur in the magnetic material, the remanent magnetic flux density of the magnetic material can be increased by raising the sintering temperature to increase the relative density of the magnetic material. This is consistent with the principle of improving the remanent magnetic flux density of a magnetic material.

[0114] The mechanism for improving the coercivity of R-Fe-B rare-earth magnetic materials is not limited to the mechanism described above.

[0115] The R content in R-Fe-B rare earth magnetic materials is 29.5 to 33.5% by mass percentage, with the most preferred range being 30.0 to 32.0%. To give a magnetic material excellent magnetic properties, it is necessary to design it to have a special microstructure with structural features such as those of the present invention. R is mainly involved in the construction of the main phase and grain boundary phase, and R is in the main phase, for example, Nd2Fe 14 B and PR2Fe 14 B and other R2T 14 R is involved in the construction of B. Furthermore, R is also involved in the construction of grain boundary phases, including triangular region grain boundary phases and two-particle grain boundary phases. The grain boundaries of a magnetic material are constructed as a continuous structure of the magnetic material and the grain boundary phase of the corresponding component. As the proportion of R element increases, the magnetic isolation effect of the microstructure improves, thereby significantly increasing the coercivity of the magnetic material. The grain boundaries are designed based on the construction ratio of the main R phase, and as the proportion of grain boundaries increases, the residual magnetic flux density of the magnetic material decreases accordingly. Magnetic materials with various components are designed according to the various microstructures required, thereby obtaining R-Fe-B rare-earth magnetic materials with various magnetic properties.

[0116] The B content in R-Fe-B rare earth magnetic materials is 0.85 to 1.05% by mass percentage, with the most preferred range being 0.85 to 0.95%. B is, for example, Nd2Fe. 14 B is essential for the formation of the tetragonal phase of B, and B increases the interatomic distance between Fe atoms, reducing the number of atoms closest to Fe. Therefore, adding B can significantly improve the magnetic properties of a magnetic material. As the B content increases from 0, the α-Fe and Nd2Fe atoms in the magnetic material... 17 The phase gradually disappears, and Nd2Fe 14The tetragonal phase of B is gradually formed. Therefore, by controlling the B content in the magnetic material, the proportion of the main phase, the proportion and composition of the grain boundary phase can be controlled, thereby controlling the residual magnetic flux density and coercivity of the magnetic material. Furthermore, by adding other elements, the composition of the grain boundary phase of the magnetic material can be further designed. By adding a predetermined amount of Cu, a Cu-rich phase can be included in the triangular region grain boundary phase or the two-particle grain boundary phase of the magnetic material. Cu is a non-magnetic element, and including a predetermined amount of Cu increases the coercivity of the magnetic material. By including a ZrB phase or TiB phase in the grain boundary phase and setting the phase ratio to 1:1 or 1:2, it is also possible to form alloy phases containing other B elements.

[0117] The Al content in R-Fe-B rare earth magnetic materials is 0.05 to 0.80% by mass percentage, with the most preferred range being 0.10 to 0.60%. The addition of Al refines the crystal grains of the main phase, reduces the size of the R-rich and B-rich phases, and improves the wettability between the R-rich and main phases. Al penetrates the main phase, for example, Nd2Fe 14-x BAl x Al is involved in phase formation or forms an Al-containing 6:14 phase in the grain boundary phase. The Al-containing 6:14 phase can exist in the triangular region grain boundary phase and also in the two-particle grain boundary phase. As an interstitial atom, Al can exist in various phases, including the 6:14 phase, Cu-rich phase, R-rich phase, etc.

[0118] The Cu content in R-Fe-B rare-earth magnetic materials is 0.40 to 2.00% by mass percentage, and may range from 1.50 to 2.00%, 0.40 to 1.00%, or 1.00 to 1.50%. The Cu element can improve irreversible losses in magnetic materials, and by uniformly distributing it in the R-rich phase, it can improve the grain boundary phase by exhibiting effects such as corrosion resistance and oxidation resistance. By substituting the Cu element for Fe atoms in the grain boundary phase near the main phase, for example, Nd2Fe can be used. 14-x BCu xIt forms a phase or a Cu-containing 6:14 phase within the grain boundary phase. Cu can form a Cu-rich phase, which can exist in the triangular region grain boundary phase or the two-particle grain boundary phase. In other words, Cu can exist as an interstitial atom in various phases, including the 6:14 phase, Cu-rich phase, R-rich phase, etc.

[0119] The Ga content in R-Fe-B rare-earth magnetic materials is 0.30-0.60% by mass percentage. Ga lowers the melting point of grain boundaries and distributes them evenly within the grain boundary phase, thereby improving the structure of the grain boundary phase and repairing defects in the magnetic material. By substituting Ga for Fe atoms in the grain boundary phase near the main phase, for example, Nd2Fe can be produced. 14-x BGa x It forms a phase or a Ga-containing 6:14 phase within the grain boundary phase. The Ga element can form a Ga-rich phase, which can exist in the triangular region grain boundary phase or the two-particle grain boundary phase. In other words, the Ga element can exist as an interstitial atom in various phases, including the 6:14 phase, Cu-rich phase, R-rich phase, etc.

[0120] The Co content in R-Fe-B rare-earth magnetic materials is 0.50 to 2.00% by mass percentage. The Co element primarily improves the thermal stability of the magnetic material, forms stable structures such as Nd3Co at grain boundaries, and can also improve the corrosion resistance of the magnetic material. The Co element content may also be 0.50 to 1.00%. Similarly, the Co element can exist as an interstitial atom in various phases, including the 6:14 phase, Cu-rich phase, R-rich phase, and main phase.

[0121] The content of Zr or Ti in R-Fe-B rare-earth magnetic materials is 0.15 to 0.50% by mass. As high-melting-point elements, Zr or Ti can refine crystal grains and form high-melting-point B or C oxides. Similarly, Zr or Ti can exist as interstitial atoms in various phases, including 6:14 phase, Cu-rich phase, and R-rich phase.

[0122] The carbon (C) impurity content in R-Fe-B rare-earth magnetic materials is 1200 ppm or less, with a more preferable content of 300-800 ppm. High C content leads to the destruction of the main phase, impaired continuity of the grain boundary phase, and reduced magnetic properties.

[0123] The content of oxygen (O) impurities in R-Fe-B rare-earth magnetic materials is 900 ppm or less, with a more preferable content being 200-700 ppm. If the O content is high, a large amount of HCP-type R oxide is formed, the HCP-type crystal lattice is destroyed, its thickness increases, and the number of defects increases. As a result, the misfit at the interface between the affected phase and the main phase increases, the stray magnetic field increases, and the magnetic properties of the magnetic material deteriorate rapidly.

[0124] The nitrogen (N) impurity content in R-Fe-B rare-earth magnetic materials is 1500 ppm or less, with a more preferable content of 500-1000 ppm. If the N content is too high, R nitrides will form, reducing the magnetic properties of the magnetic material. [Explanation of Symbols]

[0125] 1. Main phase particle (I) (corresponding to the main phase in Figure 3) 2. Main phase particles (II) (corresponding to the HCP-type R-rich phase in Figure 3) 3. Triangular region grain boundary phase (III) (corresponding to the RCu phase in Figure 3) 4. Two-particle grain boundary phase (corresponding to the RGa phase in Figure 3) 5. Needle-shaped ZrBx phase and TiBx phase (corresponding to the R-rich phase in Figure 3) 6 6:14 phase 7. Needle-shaped ZrBx phase and needle-shaped TiBx phase

Claims

1. R-Fe-B type rare earth magnetic material, It includes the main phase, the two-particle grain boundary phase, and the triangular region grain boundary phase. The components of the R-Fe-B rare earth magnetic material are, in terms of mass percentage, R: 29.5–33.5%, B: 0.85–1.05%, Al: 0.05–0.8%, Cu: 0.4–2.0%, Ga: 0.3–0.6%, Co: 0.5–2.0%, Zr or Ti: 0.15–0.5%, the remainder being Fe and unavoidable impurities. The main phase is R ２ T １４ It is phase B, where R is at least one light rare earth element, and T is Fe, or Fe and Co. The two-particle grain boundary phase comprises a grain boundary Cu-rich phase and / or a needle-shaped ZrBx phase and / or a needle-shaped TiBx phase, The grain boundary Cu-rich phase has an atomic percentage ratio of Cu to R within 20 nm near the interface with the main phase of 1.5 < Cu / R ≤ 2.0, and an atomic percentage ratio of Cu to R within 50 nm of the grain boundary center of 1.0 < Cu / R ≤ 1.

5. The width of the needle-shaped ZrBx phase is 10 to 50 nm. The width of the needle-shaped TiBx phase is 10 to 50 nm. The aforementioned triangular region grain boundary phase is FCC-type NdOx phase, An RCu phase in which the atomic percentage ratio of Cu to R is 1.0 < Cu / R ≤ 1.5, An RGa phase in which the atomic percentage ratio of Cu to R is 0.1 < Ga / R ≤ 0.5, The atomic percentage ratio of Cu, R, and transition metal elements is 6:14 in the phase, R-rich phase, and / or the needle-shaped ZrBx phase, and / or the needle-shaped TiBx phase, A rare earth magnetic material characterized by the following features:

2. The aforementioned R is at least one of Nd, Pr, La, Ce, and Sm. The R-Fe-B rare earth magnetic material according to feature 1.

3. A method for producing an R-Fe-B type rare earth magnetic material according to claim 1, Step 1: The alloy prepared by mixing each raw material based on mass percentage is subjected to hydrogen absorption and dehydrogenation treatment, and alloy thin sheets of the main phase alloy and related phases are prepared after the hydrogenation treatment. Step 2: The alloy flakes of the main phase alloy and the alloy flakes of the related phase obtained in Step 1 are mixed in a predetermined ratio, crushed by machine, the alloy of the related phase is dissolved in the main phase alloy, the mixed powder material is crushed into a powder, and a base material is prepared by magnetic field molding. Step 3: The material is placed in a sintering furnace and sintered, then kept warm to create a sintered blank, and then the sintered blank is subjected to a two-stage aging treatment. In the aforementioned two-stage prescription process, The primary aging temperature is 700-900°C, and the incubation time is 3-7 hours. The secondary aging temperature is 430–550°C, and the incubation time is 3–7 hours. A method for producing an R-Fe-B type rare earth magnetic material, characterized by the following:

4. In step 1, the hydrogen storage temperature is 50 to 100°C, the dehydrogenation temperature is 500 to 650°C, and the dehydrogenation time is 3 to 7 hours. The method for producing an R-Fe-B type rare earth magnetic material according to feature 3.

5. In step 2, the particle size of the mechanical grinding was 10 to 25 μm, the particle size D50 of the powder obtained by grinding was 2 to 7 μm, and the density of the base material was 4.0 to 4.5 g / cm³. ３ That is, The method for producing an R-Fe-B type rare earth magnetic material according to feature 3.

6. The sintering temperature in step 3 is 950 to 1100°C, and the holding time is 4 to 15 hours. The method for producing an R-Fe-B type rare earth magnetic material according to feature 4.

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