RTB series fused magnets

Abstract

This invention provides an RTB-type sintered magnet that has high residual magnetic flux density and high high-temperature stability. [Solution] R2Fe 14 An RTB-type sintered magnet comprising a main phase 2 of B and a grain boundary phase, containing 12.5 to 15.5 atomic percent of R (Nd, Pr, etc.), 4.7 to 5.5 atomic percent of B, 70 atomic percent or more of T (Fe, Co, etc.), 0.3 to 3.0 atomic percent of M1 (Al, Cu, Ga, etc.), 0.05 to 0.5 atomic percent of M2 (Sn, etc.), 0.1 to 1.0 atomic percent of M3 (Zr, etc.), and 0.8 atomic percent or less of O, with the remainder being C. The particles are N and unavoidable impurities, and the grain boundary phase contains RT-(M1,M2) phase 3, in which the R concentration, M1 concentration, and M2 concentration are higher than that of the main phase, and R-M2-C phase 1, in which the R concentration and M2 concentration are higher than that of this phase and the C concentration is higher than that of the main phase, and the relationship 0.6 < [M2] / [M1] < 3.0 is satisfied between the M1 concentration and the M2 concentration relative to the total amount of R, T, M1, and M2 in the RT-(M1,M2) phase.

Description

[Technical Field]

[0001] This invention relates to an RTB-type sintered magnet having high residual magnetic flux density and coercivity. [Background technology]

[0002] RTB-type sintered magnets (hereinafter sometimes referred to as Nd magnets) are functional materials indispensable for energy saving and high performance, and their range of applications and production volume are expanding year by year. For example, they are used in drive motors and electric power steering motors in hybrid and electric vehicles, and in air conditioner compressor motors. In these various applications, the high coercivity (hereinafter referred to as H) of RTB-type sintered magnets is required to withstand use in high-temperature environments. cJ This is called [name of the motor]. ) is a major advantage, but to operate the motor in harsher environments, further H cJ Improvement is needed.

[0003] H of Nd magnet cJ As a method to improve this, conventionally R2T 14 To improve the crystalline magnetic anisotropy of the B phase, methods have been employed to replace some of the R with heavy rare earth elements such as Dy and Tb. On the other hand, considering the resource risks associated with rare elements such as Dy and Tb, methods have been employed to avoid using heavy rare earth elements. cJ Active efforts are being made to develop methods to improve this, and various techniques have been proposed, such as refining the main phase crystal grains and controlling the structure of the grain boundary phase.

[0004] For example, Japanese Patent Publication No. 7-130522 (Patent Document 1) describes R6T containing Sn as M. 13 A method for manufacturing a permanent magnet having an M phase is described, and it is reported that permanent magnets manufactured by the prescribed method exhibit improved thermal stability of coercivity.

[0005] Furthermore, Japanese Patent Publication No. 2018-125445 (Patent Document 2) states that in a rare earth magnet containing Ga and Sn in a predetermined ratio, the addition of Sn suppresses the formation of the RT-Ga phase at the two-particle grain boundary, while promoting the formation of the R-Ga-Cu phase, thereby increasing H cJ It is stated that it is possible to improve this.

[0006] Furthermore, Japanese Patent Publication No. 2015-119132 (Patent Document 3) proposes a method for suppressing demagnetization of a magnet at high temperatures by forming a structure in which, in a rare earth magnet with a predetermined composition range including main phase particles and grain boundary phases, a first grain boundary phase (where R is a rare earth element, T is one or more iron group elements with Fe as an essential element, and M is at least one element selected from Al, Ge, Si, Sn, and Ga; the same applies hereinafter) has a composition ratio of R: 20-40 atomic%, T: 60-75 atomic%, and M: 1-10 atomic%, and a second grain boundary phase (where R is 50-70 atomic%, T: 10-30 atomic%, and M: 1-20 atomic%) has a composition ratio of R: 50-70 atomic%, T: 10-30 atomic%, and M: 1-20 atomic%, in predetermined proportions.

[0007] Furthermore, Japanese Patent Publication No. 2017-228771 (Patent Document 4) describes a grain boundary phase consisting of 25-35 atomic% R (R being two or more elements selected from rare earth elements including Y, and requiring Nd and Pr), 2-8 atomic% M1 (M1 being two or more elements selected from Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi), and 8 A magnet is disclosed that contains an R-Fe(Co)-M1 phase having a composition of less than atomic percent Co and the remainder Fe, wherein the R-Fe(Co)-M1 phase is a crystalline phase A in which crystallites with a particle size of 10 nm or more are formed at grain boundary triple points, and a microcrystalline phase B in which amorphous and / or crystallites with a particle size of less than 10 nm are formed at the interparticle grain boundaries or both the interparticle grain boundaries and grain boundary triple points, and has a different composition from phase A. It is also disclosed that in this sintered magnet, by adding Si, Ge, In, Sn, Pb, etc. as M1 to form two or more R-Fe(Co)-M1 phases with different peritectic temperatures, high coercivity can be obtained at high temperatures even without containing Dy, Tb, etc. [Prior art documents]

Patent Document

[0008]

Patent Document 1

Patent Document 2

Patent Document 3

Patent Document 4

Summary of the Invention

Problems to be Solved by the Invention

[0009] However, in the rare earth magnet described in Patent Document 1, although it is shown in the examples that the coercivity temperature coefficient is improved by adding Sn, that is, the high-temperature stability of the magnet is improved, the room-temperature coercivity is conversely decreased by adding Sn, and the effect of improving the high-temperature stability by adding Sn cannot be purely enjoyed.

[0010] In addition, the invention of Patent Document 2 aims to obtain a high residual magnetic flux density (hereinafter referred to as Br) and a high H by adding Sn without using heavy rare earths such as Dy as much as possible. However, in recent years, a high H exceeding 20 kOe is required without using Dy, but the characteristics of the exemplified magnets are all insufficient. cJ In the invention of Patent Document 3, a magnet with a low high-temperature demagnetization rate, that is, excellent high-temperature stability, is obtained by controlling the first grain boundary phase and the second grain boundary phase to a predetermined ratio. However, judging from the exemplified characteristics, it is required that the cooling after the second aging treatment be at a rate of 100 °C / min or more, and it is presumed that it is difficult to achieve in a mass production process of heat-treating a large number of magnets simultaneously. cJ

[0011]

[0012] ​​On the other hand, the magnet described in Patent Document 4 improves the coercivity temperature coefficient and obtains high high-temperature coercivity by adding additive elements such as Si and Sn to form an R-Fe(Co)-M1 phase with a relatively high peritectic temperature. In particular, the R-Fe(Co)-M1 phase containing Sn has a peritectic temperature of 1080°C, which is equal to or higher than the sintering temperature. In the magnet described in Patent Document 4, the amount of R-Fe(Co)-M1 phase precipitated tends to increase and Br tends to decrease compared to the case where additive elements that raise the peritectic temperature of the R-Fe(Co)-M1 phase are not added.

[0013] This invention has been made in view of the above problems, and aims to provide an RTB-type sintered magnet having high Br and high high-temperature stability by optimizing its composition and forming a specific structure in an RTB-type sintered magnet. [Means for solving the problem]

[0014] The inventors of the present invention conducted diligent studies to solve the above problems and concluded that R (one or more elements selected from rare earth elements, with Nd being essential), B, T (Fe and Co, with Fe making up 90% or more of T in atomic ratio), M1 (M1 is one or more elements selected from Al, Mn, Ni, Cu, Zn, Ga, Pd, Ag, Cd, Sb, Pt, Au, Hg, and Bi), and M2 (M2 is one or more elements selected from Si, Ge, In, Sn, and Pb). In an RTB-type sintered magnet containing elements (T, V, Cr, Zr, Nb, Mo, Hf, Ta, and W), O, C, and N, it was discovered that by adjusting the composition to a predetermined range and including RT-(M1,M2) phases and R-M2-C phases having specific atomic concentrations in the grain boundary phase, an RTB-type sintered magnet with high Br and high high-temperature stability can be obtained, thus completing the present invention.

[0015] In other words, the present invention provides the following RTB-type sintered magnet. 1. R2Fe 14 An RTB-type sintered magnet comprising a main phase that is an intermetallic compound and a grain boundary phase, 12.5-15.5 atomic percent of R (R is one or more elements selected from rare earth elements, with Nd being mandatory), 4.7-5.5 atomic percent of B, 70 atomic percent or more of T (T is Fe and Co, with Fe making up 90% or more of T in atomic ratio), 0.3-3.0 atomic percent of M1 (M1 is selected from Al, Mn, Ni, Cu, Zn, Ga, Pd, Ag, Cd, Sb, Pt, Au, Hg, Bi) The composition contains one or more elements, 0.05 to 0.5 atomic percent of M2 (M2 is one or more elements selected from Si, Ge, In, Sn, and Pb), 0.1 to 1.0 atomic percent of M3 (M3 is one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta, and W), and 0.8 atomic percent or less of O, with the remainder being C, N, and unavoidable impurities. The grain boundary phase contains an RT-(M1,M2) phase with higher R, M1, and M2 concentrations than the main phase, and an R-M2-C phase with higher R and M2 concentrations than the RT-(M1,M2) phase and higher C concentration than the main phase. Furthermore, when the atomic concentration of M1 is [M1] and the atomic concentration of M2 is [M2] relative to the total amount of R, T, M1, and M2 in the RT-(M1,M2) phase, both are given by the following formula (1) 0.6 < [M2] / [M1] < 3.0 ···(1) An RTB-type sintered magnet characterized by satisfying the following relationship. 2. An RTB-type sintered magnet having a C content of 0.1 to 1.0 atomic percent. 3. R ​​in the above grain boundary phase 1.1 One or two RTB-type sintered magnets that do not contain the T4B4 compound phase or the M3 boride phase, but contain the M3 carbide phase. 4. An RTB-type sintered magnet of any of the above grain boundary phases, wherein the RT-(M1,M2) phase contains R, M1, and M2 in a range satisfying 25≦[R]≦35, 1≦[M1]≦7, and 0<[M2]≦5, where [R] is the atomic concentration (atomic %) of R, [M1] is the atomic concentration of M1, and [M2] is the atomic concentration of M2 relative to the total amount of R, T, M1, and M2 in the RT-(M1,M2) phase. 5. An RTB-type sintered magnet according to any of 1 to 4, characterized in that it contains Sn as M2, and the M2 content is 0.05 to 0.3 atomic percent. 6. An RTB-type sintered magnet according to any of 1 to 5, characterized in that it contains Sn as M2 and contains an R-Sn-C phase as the R-M2-C phase in the grain boundary phase. 7. An RTB-type sintered magnet according to any of 1 to 6, wherein the above-mentioned R-M2-C phase is an R-(M1)M2-C phase further containing the element M1, and the concentration of M1 in the R-(M1)M2-C phase is higher than the concentration of M1 in the above-mentioned main phase particles. 8. An RTB-type sintered magnet of any of types 1 to 7, wherein the average grain size D50 (μm), which is the area average of the equivalent circular diameter of the main phase crystal grains in a cross section parallel to the orientation direction of the RTB-type sintered magnet, is 1.2 to 4.0 μm. [Effects of the Invention]

[0016] According to the present invention, an RTB-type sintered magnet having high Br and high high-temperature stability can be obtained. [Brief explanation of the drawing]

[0017] [Figure 1] This is an electron microscope image (backscattered electron image) of the sintered body after low-temperature heat treatment in Example 1, observed with an electron microscope in a cross-section parallel to the magnetization direction. [Modes for carrying out the invention]

[0018] As described above, the RTB-type sintered magnet of the present invention is an RTB-type sintered magnet containing R (one or more elements selected from rare earth elements, with Nd being essential), B, T (Fe and Co, with 90% or more of T being Fe in atomic ratio), M1 (M1 is one or more elements selected from Al, Mn, Ni, Cu, Zn, Ga, Pd, Ag, Cd, Sb, Pt, Au, Hg, and Bi), M2 (M2 is one or more elements selected from Si, Ge, In, Sn, and Pb), M3 (M3 is one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta, and W), O, C, and N, and the grain boundary phase contains an RT-(M1,M2) phase and an R-M2-C phase having specific atomic concentrations.

[0019] As described above, the element R constituting the RTB-type sintered magnet of the present invention is one or more elements selected from rare earth elements, with Nd being essential. Other rare earth elements besides Nd include Pr, La, Ce, Gd, Dy, Tb, and Ho, with Pr, Dy, and Tb being particularly preferred, and Pr being even more preferred. In addition, the R element may include elements introduced into the magnet after sintering by grain boundary diffusion.

[0020] The content of element R is preferably 12.5 atomic% or more, and more preferably 13.0 atomic% or more, from the viewpoint of suppressing the crystallization of α-Fe in the raw alloy during manufacturing and ensuring sufficient densification. Although it is difficult to eliminate α-Fe even with homogenization, within the above range, the H of RTB sintered magnets is cJ This can suppress a significant decrease in angularity. The same applies when producing the raw alloy by the strip casting method, which makes α-Fe crystallization less likely. In addition, it can prevent insufficient densification of RTB-type sintered magnets, as the amount of liquid phase, mainly composed of R components which plays a role in promoting densification during the sintering process described later, decreases, thus reducing sinterability. On the other hand, if the R content is too high, R2T in the sintered magnet 14 The proportion of phase B decreases, leading to a decrease in Br. Therefore, from the viewpoint of preventing this decrease in Br, the content of R is 17 atomic% or less, preferably 15.5 atomic% or less, and more preferably 15 atomic% or less.

[0021] The element T constituting the RTB-type sintered magnet of the present invention may contain Fe and may also contain Co. At least 90% of T is Fe in atomic ratio. From the viewpoint of obtaining a higher Br, the T content is 70 atomic percent or more, and preferably 75 atomic percent or more. Furthermore, the T content is not particularly limited, but R2T 17 Deterioration of angularity due to phase precipitation and H cJ From the viewpoint of suppressing a decrease in [the substance], it is preferable that the concentration be 82 atomic percent or less, and more preferably 80 atomic percent or less.

[0022] Furthermore, the above Co is R2T 14 It is possible to substitute a portion of the Fe contained in the T element of the B phase and RT-(M1,M2) phase. The Co content is preferably 0.1 atomic% or more of the total magnet, and more preferably 0.3 atomic% or more, from the viewpoint of obtaining an improved Curie temperature and corrosion resistance effect. cJ From the viewpoint of stably obtaining the product, the Co content is preferably 3.0 atomic% or less, and more preferably 2.0 atomic% or less.

[0023] The RTB-type sintered magnet of the present invention contains boron (B), and a portion of it may be substituted with carbon (C). The B content is 4.5 atomic% or more, preferably 4.7 atomic% or more, more preferably 4.8 atomic% or more, and 5.5 atomic% or less, preferably 5.3 atomic% or less, more preferably 5.2 atomic% or less. If the B content is less than 4.5 atomic%, the formed R2T 14 The proportion of phase B decreases, and Br drops significantly, along with R2T 17 Because a phase is formed, the hornbone structure deteriorates. On the other hand, if the B content exceeds 5.5 atomic percent, R 1.1 When the T4B4 compound phase is formed, it may not be possible to sufficiently form the RT-(M1,M2) phase, resulting in poor coercivity. In addition, the M3 boride phase is preferentially formed, inhibiting the precipitation of the M3 carbide phase. As will be explained later, this is due to the presence of excess C in the grain boundary phase. cJThis is undesirable because it induces a decrease. In other words, in the present invention, although not particularly limited, R in the grain boundary phase 1.1 Preferably, the T4B4 compound phase and the M3 boride phase are not included, but the M3 carbide phase is included.

[0024] The element M1 constituting the RTB-type sintered magnet of the present invention is one or more elements selected from Al, Mn, Ni, Cu, Zn, Ga, Pd, Ag, Cd, Sb, Pt, Au, Hg, and Bi. By adding M1 at a predetermined concentration, the RT-(M1,M2) phase can be stably formed. The M1 content is 0.1 atomic% or more, preferably 0.3 atomic% or more, and 3.0 atomic% or less, preferably 1.5 atomic% or less. If the M1 content is less than 0.1 atomic%, the amount of RT-(M1,M2) phase formed is insufficient, and sufficient coercivity cannot be obtained. On the other hand, if the M1 content exceeds 3.0 atomic%, the Br decreases, which is undesirable.

[0025] The element M2 constituting the RTB-type sintered magnet of the present invention is one or more elements selected from Si, Ge, In, Sn, and Pb. By adding M2 at a predetermined content, the RT-(M1,M2) phase and the R-M2-C phase can be stably formed. From the viewpoint of the stability of the R-M2-C phase, it is preferable to include Sn and In, and particularly preferable to include Sn. The M2 content is 0.01 atomic% or more, preferably 0.05 atomic% or more, and 0.5 atomic% or less, preferably 0.3 atomic% or less. If the amount of M2 added is less than 0.01 atomic%, the RT-(M1,M2) phase cannot be formed, and no improvement in the coercivity temperature coefficient is observed. On the other hand, if the amount of M2 added is greater than 0.5 atomic%, the volume ratio of the main phase decreases, leading to a significant decrease in Br.

[0026] The element M3 constituting the RTB-type sintered magnet of the present invention is one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta, and W, and its content is 0.05 atomic% or more, preferably 0.1 atomic% or more, and 1.0 atomic% or less, preferably 0.5 atomic% or less. If it is less than 0.05 atomic%, the effect of suppressing abnormal grain growth of crystal grains during the sintering process cannot be obtained. Also, if it is more than 1.0 atomic%, the M3 boride phase and M3 carbide phase are formed in excess, which reduces the amount of B and C needed to form the main phase, leading to a decrease in the Br due to the decrease in the main phase ratio, and consequently to R2Fe 17 The formation of this phase worsens the angularity. Furthermore, the ratio of constituent elements in M3 boride is M3:B=1:2, which is higher than the ratio of constituent elements in M3 carbide (M3:C=1:1). Therefore, it is preferable that the M3 boride phase is not present in the grain boundary phase, as it significantly reduces the ratio of the main phase. In addition, M3 carbide has a high melting point and suppresses abnormal grain growth by segregating at the grain boundary triple point, and also fixes C in the grain boundary phase, thereby promoting H cJ An improvement effect can be expected.

[0027] The oxygen (O) content contained in the RTB-type sintered magnet of the present invention is high at room temperature H cJ , and high temperature H cJ From the viewpoint of obtaining R2T, the content is 0.8 atomic% or less, preferably 0.5 atomic% or less, and more preferably 0.3 atomic% or less. If the content exceeds 0.8 atomic%, the amount of R-OCN phase formation increases, which reduces the amount of C that can substitute for part of the main phase, and R2T 17 The phase precipitates, worsening the angularity.

[0028] Furthermore, the RTB-type sintered magnet of the present invention may contain any element other than the above-mentioned R, T, B, M1, M2, M3, and O, and may contain C, N, etc. as such arbitrary elements.

[0029] The carbon (C) content of the RTB-type sintered magnet of the present invention is not particularly limited, but is preferably 0.1 atomic% or more, more preferably 0.4 atomic% or more, and even more preferably 0.5 atomic%. The preferred upper limit is 1.0 atomic% or less, more preferably 0.8 atomic% or less, and even more preferably 0.7 atomic% or less. The C originates from the raw materials and lubricants added to improve the orientation of the fine powder during magnetic field molding. By adding lubricants in an amount such that the C content is 0.1 atomic% or more, sufficient orientation can be obtained in the molding process, the Br can be increased, and the R-M2-C phase described later can be formed well. On the other hand, if the C content is 1.0 atomic% or less, excess C is generated, resulting in a low temperature of H cJ This can suppress the decline.

[0030] The above N content is good H cJ From the viewpoint of obtaining the desired result, it is preferably 1.0 atomic% or less, more preferably 0.5 atomic% or less, and even more preferably 0.2 atomic% or less.

[0031] The microstructure of the RTB-type sintered magnet of the present invention includes R2T 14 The main phase is an intermetallic compound (B). The grain boundary phase includes an RT-(M1,M2) phase and an R-M2-C phase. In addition to the above, the grain boundary phase may also include an RT-M1 phase that does not contain M2, or an M3 carbide phase. By segregating the M3 carbide phase at the grain boundary triple point, excess carbon (excess C) can be fixed, and the decrease in room temperature coercivity can be suppressed. Furthermore, the RTB-type sintered magnet of the present invention may contain an R-rich phase in the grain boundary phase, and may also contain phases of compounds that are unavoidable impurities introduced during the manufacturing process, such as R carbides, R oxides, R nitrides, R halides, and R acid halides, but may also contain Br and H cJ From the perspective of suppressing the decline, it is preferable to keep it to the minimum necessary.

[0032] The RT-(M1,M2) phase described above is a phase in which the R concentration, M1 concentration, and M2 concentration are higher than those of the main phase. Furthermore, when the atomic concentration (atomic %) of R relative to the total amount of R, T, M1, and M2 in the RT-(M1,M2) phase is [R], the atomic concentration of M1 is [M1], and the atomic concentration of M2 is [M2], it is preferable that 25≦[R]≦35, 1≦[M1]≦7, 0<[M2]≦5, and 0.6<[M2] / [M1]<3.0, and more preferably 27≦[R]≦33, 2≦[M1]≦5, 1≦[M2]≦4, and 0.8<[M2] / [M1]<2.0. Having these ranges allows for good high-temperature H cJ This method achieves the desired result while suppressing the decrease in Br associated with the precipitation of the RT-(M1,M2) phase. In particular, as the B content increases, the [M2] / [M1] ratio decreases, and when it is 0.6 or less, the high-temperature coercivity relative to the room-temperature coercivity decreases, and the amount of RT-(M1,M2) phase formation increases, which may cause a decrease in Br. On the other hand, when [M2] / [M1] is 3.0 or more, the amount of RT-(M1,M2) phase formation may be insufficient, and the effect of improving the high-temperature coercivity relative to the room-temperature coercivity may not be fully obtained. Furthermore, the grain boundary phase may contain an RT-M1 phase that does not contain M2.

[0033] Furthermore, the RTB-type sintered magnet of the present invention has a high room temperature H cJ From the viewpoint of obtaining the desired result, the grain boundaries contain an R-M2-C phase with higher R, M2, and C concentrations than the RT-(M1,M2) phase, and from the viewpoint of the stability of the R-M2-C phase, it is preferable that M2 contains Sn or In, and particularly preferable that it contains Sn. Furthermore, the R-M2-C phase may contain M1 at a higher concentration than the M1 concentration in the main phase particles. Specifically, when the concentration of R atoms in the R-(M1)M2-C phase is [R'], the concentration of M1 atoms is [M1'], the concentration of M2 atoms is [M2'], and the concentration of C atoms is [C], relative to the total amount of R, M1, M2, and C, preferably the values ​​are 35≦[R']≦55, 0≦[M1']≦10, 5≦[M2']≦25, 25≦[C]≦45, more preferably 40≦[R']≦50, 0≦[M1']≦5, 10≦[M2']≦20, and 30≦[C]≦40. Within this range, the R-(M1)M2-C phase is stably formed, and H is formed by immobilizing C in the liquid phase. cJAn improvement effect can be obtained.

[0034] Here, the composition of the RT-(M1,M2) phase and R-M2-C phase in the grain boundary phase can be confirmed using EDS (energy-dispersive X-ray spectrometer) and WDS (wavelength-dispersive X-ray spectrometer). It is generally known that when C analysis is performed using an EDS device attached to a SEM (scanning electron microscope), contamination is superimposed on the analytical value. Therefore, when analyzing the composition of the R-M2-C phase, it is preferable to reduce contamination and obtain a clean surface by milling the magnet surface to be used for analysis with an ion milling or FIB (focused ion beam) device, removing the effects of oxidation on the outermost surface, and then measuring with an EDS device. Furthermore, considering that the influence of C contamination cannot be completely eliminated in analysis by EDS or WDS, and that it is difficult to discuss the absolute value of C concentration, when the composition is calculated using only R, M1, and M2 in the R-(M1)M2-C phase, it is preferable that 65≦[R']≦85, 0≦[M1']≦10, and 15≦[M2']≦35, and more preferably 70≦[R']≦80, 0≦[M1']≦5, and 20≦[M2']≦30.

[0035] Furthermore, to identify the RT-(M1,M2) phase and the R-M2-C phase, it is preferable to confirm them by obtaining ED (electron diffraction) patterns or the like. When the RT-(M1,M2) phase is tetragonal and M2 is Sn or In, the R-M2-C phase is CaTiO3 type cubic.

[0036] Furthermore, the grain size D50 (μm), defined as the median of the equivalent circular diameter of the main phase crystal grains in a plane parallel to the magnetization direction of an RTB-type sintered magnet, is sufficient H cJ From the viewpoint of obtaining the desired result, the particle size is preferably 4.0 μm or less, more preferably 3.5 μm or less, and from the viewpoint of obtaining a sufficient degree of orientation within an appropriate range of lubricant addition amounts, it is preferably 1.2 μm or more, more preferably 1.8 μm or more.

[0037] Conventional RTB-type sintered magnets use materials such as Sn and Si, and R6T 13While attempts had been made to improve high-temperature coercivity by adding elements that increase the peritectic temperature of the M phase, R6T 13 The M phase is so R6T that it can be seen even immediately after sintering and rapid cooling. 13 The problem is that the M phase is actively formed, which significantly reduces Br, and in particular, in the magnet exemplified in Patent Document 4, the addition of Sn reduces Br by 200G. On the other hand, as described above, the RTB-type sintered magnet of the present invention forms an R-M2-C phase within a predetermined range of oxygen concentration and M2 element addition, thereby suppressing the reduction in Br due to the addition of M2 element, while achieving both high room temperature coercivity and high-temperature stability. The reason for this is not entirely clear, but it can be inferred as follows.

[0038] First, regarding the effect of improving coercivity by lowering the oxygen concentration in the magnet to 0.1-0.8 atomic percent compared to conventional methods, it is thought that the coercivity improves because the amount of R in the liquid phase increases by reducing the oxygen content that forms the R oxide phase and R-OCN phase compared to conventional methods. On the other hand, due to the reduced oxygen content, excess C (excess C) present in the grain boundary phase becomes H at room temperature. cJ While it is known that adding M2 or M3 elements reduces Br, the formation of excess C can be suppressed by creating an R-M2-C phase or an M3 carbide phase in the sintered magnet. On the other hand, the RT-(M1,M2) phase has a higher decomposition temperature than the RT-M1 phase and is formed at the grain boundary triple point at a relatively high temperature during the cooling process after sintering. The interface with the main phase exhibits a rounded shape, which not only suppresses the generation of reverse domains but also reduces the local demagnetizing field near the grain boundary triple point, which is thought to be effective in suppressing the decrease in coercivity at high temperatures. Previously, controlling the amount of precipitation of the RT-(M1,M2) phase was difficult due to its high peritectic temperature, resulting in a significant decrease in Br compared to cases without M2 elements. However, in this invention, by appropriately forming the R-M2-C phase, the volume fraction of the RT-(M1,M2) phase is reduced, and the effect of C on the decrease in coercivity is further reduced. As a result, it is thought that a higher high-temperature coercivity can be obtained while suppressing the decrease in Br due to the addition of M2 elements compared to conventional methods.

[0039] Next, the method for manufacturing the R-Fe-B sintered magnet of the present invention will be described below. The steps for manufacturing the R-Fe-B sintered magnet of the present invention are basically the same as those of a conventional powder metallurgy method and are not particularly limited, but typically include a melting step of melting raw materials to obtain a raw material alloy, a grinding step of grinding a raw material alloy having a predetermined composition to prepare alloy fine powder, a molding step of compacting the alloy fine powder in a magnetic field to obtain a molded body, and a sintering step of heat-treating the molded body to obtain a sintered body.

[0040] First, in the melting process described above, the metals or alloys that will be the raw materials for each element are weighed to obtain the predetermined composition described above in the present invention. For example, the raw materials are melted by high-frequency induction melting and then cooled to produce the raw material alloy. For casting the raw material alloy, melt casting methods such as pouring into flat molds or book molds, or strip casting methods are generally employed. In addition, R2T is the main phase of RTB alloys. 14 The present invention can also be applied to the so-called two-alloy method, in which an alloy close to the composition of compound B and an R-rich alloy that acts as a liquid phase additive at the sintering temperature are prepared separately, and then weighed and mixed after coarse grinding. However, since the α-Fe phase tends to crystallize in the alloy close to the composition of the main phase depending on the cooling rate during casting and the alloy composition, it is preferable to perform a homogenization treatment at 700 to 1200°C for 1 hour or more in a vacuum or Ar atmosphere as needed to homogenize the structure and eliminate the α-Fe phase. Homogenization can be omitted if the alloy close to the composition of the main phase is prepared by the strip casting method. For the R-rich alloy that acts as a liquid phase additive, in addition to the above casting method, the so-called liquid quenching method can also be used.

[0041] The above grinding process can be a multi-stage process including, for example, a coarse grinding process and a fine grinding process. In the coarse grinding process, for example, a jaw crusher, a brown mill, a pin mill, or hydrogenation grinding is used. In the case of alloys made by strip casting, hydrogenation grinding is usually applied to obtain coarse powder with a particle size of, for example, 0.05 to 3 mm, and especially 0.05 to 1.5 mm. In the fine grinding process, the coarse powder obtained in the coarse grinding process is finely ground to an average particle size of preferably 0.5 to 5 μm, more preferably 1 to 3.5 μm, using a method such as jet mill grinding. In one or both of the coarse grinding and fine grinding processes of the raw alloy, it is preferable to add 0.08 to 0.30 mass% of lubricant, and more preferably 0.1 to 0.2 mass% of lubricant, for the purpose of improving the degree of orientation.

[0042] In this case, there are no particular restrictions on the lubricant, but examples include fatty acids such as stearic acid, alcohols, esters, and metal soaps. Furthermore, when adjusting the carbon content, if the amount of lubricant added exceeds the lower limit of the above range, a portion of it may be added as a carbon source other than the lubricant, such as hydrocarbons like carbon black, paraffin, or polyvinyl alcohol, and carbon black may also be added in the melting process. On the other hand, when adjusting the oxygen content to be within a predetermined range, the coarse grinding process and fine grinding process of the raw alloy are preferably carried out in a gas atmosphere such as nitrogen gas or Ar gas, but oxygen may also be introduced into the gas atmosphere and its concentration controlled.

[0043] In the above molding process, a magnetic field of 400 to 1600 kA / m is applied, and the alloy powder is oriented in the direction of its easy magnetization axis while being compacted using a compression molding machine. At this time, the density of the molded body should be 2.8 to 4.2 g / cm³. 3 It is preferable to have a molded article density of 2.8 g / cm³, in other words, from the viewpoint of ensuring the strength of the molded article and obtaining good handling properties. 3 It is preferable to have the above, while on the other hand, from the viewpoint of obtaining suitable Br by ensuring good particle orientation during pressurization while obtaining sufficient molded body strength, the molded body density should be 4.2 g / cm³. 3The following is preferable. Furthermore, molding is preferably carried out in a gas atmosphere such as nitrogen gas or Ar gas in order to suppress oxidation of the alloy fine powder.

[0044] In the above sintering process, the molded body obtained in the molding process is sintered in a high vacuum or in a non-oxidizing atmosphere such as Ar gas. Generally, it is preferable to hold the sintering at a temperature range of 950 to 1200°C for 0.5 to 15 hours. The sintered body after sintering is preferably cooled to a temperature of 400°C or lower, more preferably 300°C or lower, and even more preferably 200°C or lower. The cooling rate is not particularly limited, but until the upper limit of the above range is reached, it is preferably 5°C / min or higher, more preferably 15°C / min or higher, and preferably 100°C / min or lower, and even more preferably 50°C / min or lower.

[0045] A heat treatment step may be performed on the obtained sintered body. This heat treatment step preferably consists of two stages: a high-temperature heat treatment step in which the sintered body, cooled to a temperature of 400°C or lower, is heated to a temperature of preferably 700°C or higher, more preferably 800°C or higher, and preferably 1100°C or lower, more preferably 1050°C or lower, and then cooled again to 400°C or lower; and a low-temperature heat treatment step in which, after the high-temperature heat treatment step, the body is heated to a temperature in the range of 400 to 600°C and then cooled to 300°C or lower, preferably 200°C or lower. Furthermore, the heat treatment atmosphere at this time is preferably a vacuum or an inert gas atmosphere such as Ar gas.

[0046] The heating rate in the high-temperature heat treatment process is not particularly limited, but is preferably 1°C / min or more, more preferably 2°C / min or more, and more preferably 20°C / min or less, and more preferably 10°C / min or less. The holding time after heating to the high-temperature heat treatment temperature is preferably 1 hour or more, more preferably 10 hours or less, and more preferably 5 hours or less. After heating, the temperature is preferably cooled to 400°C or less, more preferably 300°C or less, and even more preferably 200°C or less. The cooling rate at this time is not particularly limited, but until the upper limit of the above range is reached, it is preferably 1°C / min or more, more preferably 5°C / min or more, and more preferably 100°C / min or less, and more preferably 50°C / min or less.

[0047] In the low-temperature heat treatment step following the high-temperature heat treatment step, the cooled sintered body is heated to a temperature preferably of 400°C or higher, more preferably of 430°C or higher, and preferably of 600°C or lower, and more preferably of 550°C or lower. The heating rate in the low-temperature heat treatment step is not particularly limited, but is preferably 1°C / min or higher, more preferably 2°C / min or higher, and preferably 20°C / min or lower, and even more preferably 10°C / min or lower. The holding time after heating to the low-temperature heat treatment temperature is preferably 0.5 hours or higher, more preferably 1 hour or higher, and preferably 50 hours or lower, and even more preferably 20 hours or lower. The cooling rate after heating is not particularly limited, but until it reaches the upper limit of the above range, it is preferably 1°C / min or higher, more preferably 5°C / min or higher, and preferably 100°C / min or lower, more preferably 80°C / min or lower, and even more preferably 50°C / min or lower. The sintered body after heat treatment is then usually cooled to room temperature.

[0048] Furthermore, the conditions in the high-temperature and low-temperature heat treatment processes can be appropriately adjusted within the above-mentioned range, depending on the composition, including the type and content of M1 element and M3 element, impurities, particularly the concentration of impurities caused by the atmospheric gas during manufacturing, and sintering conditions, as well as other variations caused by manufacturing processes other than the high-temperature and low-temperature heat treatment. [Examples]

[0049] The present invention will be described in more detail below with reference to examples and comparative examples, but the present invention is not limited to the following examples.

[0050] [Examples 1-2, Comparative Examples 1-2] A strip of alloy was prepared by melting the raw metal or alloy in a high-frequency induction furnace in an Ar gas atmosphere to the composition shown in Table 1 below, and then cooling the molten alloy on a water-cooled copper roll using a strip casting method. Next, the prepared alloy strip was coarsely ground by hydrogenation to obtain a coarse powder, and then 0.15% by mass of stearic acid was added to the obtained coarse powder as a lubricant and mixed. Next, the mixture of coarse powder and lubricant was finely ground in a jet mill under a nitrogen atmosphere to an average particle size of 3.0 μm. At this time, the oxygen content was adjusted by setting the oxygen concentration in the jet mill system to 10 ppm or less in Example 1 and Comparative Example 2, 50 ppm in Example 2, and 100 ppm in Comparative Example 1.

[0051] Next, the obtained fine powder was filled into a mold of a molding apparatus equipped with an electromagnet in a nitrogen atmosphere, and pressure molded by applying pressure perpendicular to the magnetic field while oriented in a magnetic field of 15 kOe (1.19 MA / m). Next, the obtained molded body was sintered in a vacuum at 1080°C for 5 hours, cooled to below 200°C at a rate of 20°C / min, then subjected to high-temperature heat treatment at 900°C for 2 hours, and cooled again to below 200°C at a rate of 20°C / min. After that, low-temperature heat treatment was performed at 450°C for 3 hours, and cooled to below 200°C at a rate of 20°C / min to obtain a sintered body. The metal elements in the composition of the sintered magnets in Table 1 were measured by ICP analysis, the oxygen concentration by inert gas fusion infrared absorption spectroscopy, the nitrogen concentration by inert gas fusion thermal conduction spectroscopy, and the carbon concentration by combustion infrared absorption spectroscopy.

[0052] [Table 1]

[0053] Next, the center of each obtained sintered body was cut into a rectangular parallelepiped shape measuring 18 mm × 15 mm × 12 mm to obtain sintered magnets, and the magnetic properties of each sintered magnet were measured using a BH tracer (manufactured by Toei Kogyo). In addition, the average crystal grain size D50 (μm) was measured by the following method: The cross-section parallel to the magnetization direction of the sintered magnet was polished to a mirror finish, and the grain boundary phase of the cross-section was selectively etched by immersion in a mixed solution of glycerin:ethylene glycol:nitric acid:hydrochloric acid = 4:4:1:1. Twenty-five cross-sectional images in the range of 85 × 85 μm were obtained using a laser microscope, and the cross-sectional area of ​​each particle was measured by image analysis based on the obtained cross-sectional images. The average area of ​​the diameter of each particle, calculated as the equivalent diameter of a circle, was then determined.

[0054] Table 2 shows the Br and H levels obtained from the above measurements at room temperature (approximately 23°C). cJ H at 140℃ cJ The value of H at 140℃ cJ Room temperature H cJ Ratio to (H cJ (140℃) / H cJ (23℃)) was observed. Furthermore, the surface of the cross-section of the obtained sintered body was scraped with a FIB device to remove the effects of oxidation on the outermost surface, and then the presence or absence of the RT-(M1,M2) phase, the concentration ratio of M2 to M1 in the RT-(M1,M2) phase, the value of [M2] / [M1], and the presence or absence of the R-M2-C phase, M3 boride phase, and M3 carbide phase were examined using an EDS device attached to the SEM. These are shown in Table 3.

[0055] [Table 2]

[0056] [Table 3]

[0057] As shown in Tables 1 and 2, among magnets with different oxygen concentrations, the sintered magnets of Examples 1 and 2, which were produced by a method satisfying the conditions of the present invention, had higher coercivity at 140°C compared to Comparative Example 1. Furthermore, while the high-temperature coercivity values ​​were equivalent to the room-temperature coercivity in Examples 1 and 2 and Comparative Example 1, the room-temperature coercivity was higher as the oxygen concentration decreased. In addition, as shown in Table 3, the R-M2-C phase was observed in the magnet structure of Examples 1 and 2, which had high room-temperature and high-temperature coercivity, whereas the R-M2-C phase was not observed in Comparative Example 1. Regarding the M3 compound phase, in Examples 1 and 2, the M3 element formed only carbides, whereas in Comparative Example 1, it formed both borides and carbides. Next, comparing Example 1 and Comparative Example 2, which have equivalent oxygen concentrations but differ in the presence or absence of Sn addition, Example 1, which has Sn added, shows superior coercivity at room temperature and high temperature compared to Comparative Example 2. Furthermore, the decrease in Br due to Sn addition is less than 100G, indicating that the magnet of the present invention was able to suppress the decrease in Br due to Sn addition compared to conventional magnets.

[0058] Furthermore, the sintered body of Example 1 after low-temperature heat treatment was observed using an electron microscope in a cross-section parallel to the magnetization direction. The electron microscope image (backscattered electron image) is shown in Figure 1. As shown in Figure 1, the magnet of Example 1 showed RT-(M1,M2) phase (3 in the figure) and R-M2-C phase (1 in the figure). Measurements were taken using an EDS device for the intraparticles (10 points) of the main phase particles (2 in the figure), the RT-(M1,M2) phase (10 points), and the R-M2-C phase (10 points), and the average composition was determined. The atomic percentages for predetermined elements were then calculated. The results are shown in Table 4. Note that 4 in Figure 1 is the RT-M1 phase and 5 is the M3 carbide phase.

[0059] [Table 4]

[0060] [Examples 3-4, Comparative Examples 3-4] The raw metals or alloys were melted in a high-frequency induction furnace in an Ar gas atmosphere to obtain the compositions shown in Table 5 below, and alloy strips were prepared by a strip casting method in which the molten alloy was cooled on a water-cooled copper roll. Next, the prepared alloy strips were coarsely ground by hydrogenation to obtain coarse powder, and then stearic acid was added as a lubricant to the obtained coarse powder at a concentration of 0.15 mass% in Examples 3 and 4, and 0.09 mass% in Comparative Example 4, and mixed. Subsequently, the mixture of coarse powder and lubricant was finely ground in a jet mill in a nitrogen stream with an oxygen concentration of 10 ppm or less to an average particle size of approximately 3.0 μm.

[0061] The following molding and heat treatment were performed using the same method as in Example 1, and the magnetic properties and average grain size were measured. The results are shown in Table 6. In addition, the presence or absence of the RT-(M1,M2) phase, the value of [M2] / [M1] (the concentration ratio of M2 to M1 in the RT-(M1,M2) phase), and the presence or absence of the R-M2-C phase, M3 boride phase, and M3 carbide phase were confirmed using the same method as in Example 1. The results are shown in Table 7.

[0062] [Table 5]

[0063] [Table 6]

[0064] [Table 7]

[0065] As shown in Tables 5-7, Examples 3 and 4, in which the amount of Sn added satisfied the range of the present invention, showed similar room-temperature coercivity to Comparative Example 2 (Table 2), which did not contain Sn, but obtained high high-temperature coercivity. On the other hand, Comparative Example 3, in which an excess of Sn was added, showed a decrease in both Br, room-temperature coercivity, and high-temperature coercivity compared to Examples 3 and 4. Furthermore, in Comparative Example 4, in which the amount of B added exceeded the range of the present invention, although the R-M2-C phase was formed, the RT-(M1,M2) phase was not formed, and the value of high-temperature coercivity relative to room-temperature coercivity was lower than that of Examples 2 and 3. [Explanation of symbols]

[0066] 1 R-M2-C phase 2 main phase 3 RT-(M1,M2) phase 4 RT-M1 phase 5 M3 carbide phase

Claims

1. R 2 Fe 14 A B intermetallic compound is used as the main phase and the grain boundary phase in an R-T-B sintered magnet, 12.5–15.5 atomic percent of R (where R is one or more elements selected from rare earth elements, with Nd being a must), 4.7–5.5 atomic percent of B, 70 atomic percent or more of T (where T is Fe and Co, with Fe making up 90% or more of T in atomic ratio), and 0.3–3.0 atomic percent of M 1 (M 1 (This is one or more elements selected from Al, Mn, Ni, Cu, Zn, Ga, Pd, Ag, Cd, Sb, Pt, Au, Hg, Bi), 0.05 to 0.5 atomic percent of M 2 (M 2 (One or more elements selected from Si, Ge, In, Sn, and Pb), 0.1 to 1.0 atomic percent of M 3 (M 3 It has a composition containing one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta, and W, 0.8 atomic% or less of O, and the remainder being C, N, and unavoidable impurities. In the grain boundary phase, the R concentration, M 1 concentration, M 2 concentration is higher than that of the main phase, and an R-T-(M 1 , M 2 ) phase, and an R-M 1 , M 2 ) phase having a higher R concentration and M 2 concentration than the R-T-(M 2 , M 1 , M 2 ) phase and a higher C concentration than the main phase are contained. Further, when the atomic concentration of M 1 , M 2 in the total amount of R, T, M 1 in the R-T-(M 1 , M 2 , M 2 ) phase is [M 0.6<[M 2 ] / [M 1 ]<3.0・・・(1) An R-T-B type sintered magnet characterized by satisfying the following relationship.

2. The R-T-B sintered magnet according to claim 1, wherein the content of C is 0.1 to 1.0 atomic percent.

3. R in the above grain boundary phase 1.1 T 4 B 4 Compound phase and M 3 It does not contain a boride phase, M 3 An R-T-B sintered magnet according to claim 1 or 2, containing a carbide phase.

4. R-T-(M) in the above grain boundary phase 1 , M 2 ) phase, the R-T-(M 1 , M 2 ) R, T, M 1 M 2 [R] is the atomic concentration (atomic %) of R relative to the total amount of M. 1 Atomic concentration [M 1 ), M 2 Atomic concentration [M 2 When ] is set, 25 ≤ [R] ≤ 35, 1 ≤ [M 1 ] ≤ 7, 0 < [M 2 Within the range that satisfies ] ≤ 5, R, M 1 M 2 An R-T-B sintered magnet according to claim 1 or 2, containing the following:

5. The above M 2 It contains Sn as M 2 The R-T-B sintered magnet according to claim 1 or 2, characterized in that the content of is 0.05 to 0.3 atomic percent.

6. The above M 2 It contains Sn as such, and the above R-M in the grain boundary phase 2 The R-T-B sintered magnet according to claim 1 or 2, characterized in that it contains an R-Sn-C phase as the -C phase.

7. The above R-M 2 - Phase C is further M 1 R-(M) containing the element 1 ) M 2 -C phase, and R-(M 1 ) M 2 - M in phase C 1 The concentration of M within the main phase particles is 1 An R-T-B sintered magnet according to claim 1 or 2, with a concentration higher than that of [the specified value].

8. The R-T-B sintered magnet according to claim 1 or 2, wherein the average grain size D50 (μm), which is the area average of the equivalent circular diameter of the main phase crystal grains in a cross section parallel to the orientation direction of the R-T-B sintered magnet, is 1.2 to 4.0 μm.

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