WO2015149685A1 - 一种含W的R‐Fe‐B‐Cu系烧结磁铁及急冷合金 - Google Patents

一种含W的R‐Fe‐B‐Cu系烧结磁铁及急冷合金 Download PDF

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WO2015149685A1
WO2015149685A1 PCT/CN2015/075512 CN2015075512W WO2015149685A1 WO 2015149685 A1 WO2015149685 A1 WO 2015149685A1 CN 2015075512 W CN2015075512 W CN 2015075512W WO 2015149685 A1 WO2015149685 A1 WO 2015149685A1
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Prior art keywords
sintered magnet
content
less
powder
magnet
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French (fr)
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永田浩
喻荣
蓝琴
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Xiamen Tungsten Co Ltd
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Xiamen Tungsten Co Ltd
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Priority to DK15772705.8T priority Critical patent/DK3128521T3/da
Priority to ES15772705T priority patent/ES2742188T3/es
Priority to JP2016560501A priority patent/JP6528046B2/ja
Priority to EP15772705.8A priority patent/EP3128521B8/en
Priority to CN201580002027.7A priority patent/CN105659336B/zh
Priority to BR112016013421A priority patent/BR112016013421B8/pt
Publication of WO2015149685A1 publication Critical patent/WO2015149685A1/zh
Priority to US15/185,430 priority patent/US10381139B2/en
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Priority to US16/410,090 priority patent/US10614938B2/en
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    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01F—MAGNETS; INDUCTANCES; TRANSFORMERS; SELECTION OF MATERIALS FOR THEIR MAGNETIC PROPERTIES
    • H01F1/00—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties
    • H01F1/01—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials
    • H01F1/03—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity
    • H01F1/032—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials
    • H01F1/04—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials metals or alloys
    • H01F1/047—Alloys characterised by their composition
    • H01F1/053—Alloys characterised by their composition containing rare earth metals
    • H01F1/0536—Alloys characterised by their composition containing rare earth metals sintered
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/10—Ferrous alloys, e.g. steel alloys containing cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01F—MAGNETS; INDUCTANCES; TRANSFORMERS; SELECTION OF MATERIALS FOR THEIR MAGNETIC PROPERTIES
    • H01F1/00—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties
    • H01F1/01—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials
    • H01F1/03—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity
    • H01F1/032—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials
    • H01F1/04—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials metals or alloys
    • H01F1/047—Alloys characterised by their composition
    • H01F1/053—Alloys characterised by their composition containing rare earth metals
    • H01F1/055—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5
    • H01F1/057—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01F—MAGNETS; INDUCTANCES; TRANSFORMERS; SELECTION OF MATERIALS FOR THEIR MAGNETIC PROPERTIES
    • H01F1/00—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties
    • H01F1/01—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials
    • H01F1/03—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity
    • H01F1/032—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials
    • H01F1/04—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials metals or alloys
    • H01F1/047—Alloys characterised by their composition
    • H01F1/053—Alloys characterised by their composition containing rare earth metals
    • H01F1/055—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5
    • H01F1/057—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B
    • H01F1/0571—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B in the form of particles, e.g. rapid quenched powders or ribbon flakes
    • H01F1/0575—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B in the form of particles, e.g. rapid quenched powders or ribbon flakes pressed, sintered or bonded together
    • H01F1/0577—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B in the form of particles, e.g. rapid quenched powders or ribbon flakes pressed, sintered or bonded together sintered
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01F—MAGNETS; INDUCTANCES; TRANSFORMERS; SELECTION OF MATERIALS FOR THEIR MAGNETIC PROPERTIES
    • H01F41/00—Apparatus or processes specially adapted for manufacturing or assembling magnets, inductances or transformers; Apparatus or processes specially adapted for manufacturing materials characterised by their magnetic properties
    • H01F41/02—Apparatus or processes specially adapted for manufacturing or assembling magnets, inductances or transformers; Apparatus or processes specially adapted for manufacturing materials characterised by their magnetic properties for manufacturing cores, coils, or magnets
    • H01F41/0253—Apparatus or processes specially adapted for manufacturing or assembling magnets, inductances or transformers; Apparatus or processes specially adapted for manufacturing materials characterised by their magnetic properties for manufacturing cores, coils, or magnets for manufacturing permanent magnets
    • H01F41/0293—Apparatus or processes specially adapted for manufacturing or assembling magnets, inductances or transformers; Apparatus or processes specially adapted for manufacturing materials characterised by their magnetic properties for manufacturing cores, coils, or magnets for manufacturing permanent magnets diffusion of rare earth elements, e.g. Tb, Dy or Ho, into permanent magnets

Definitions

  • the present invention relates to the technical field of manufacturing magnets, and more particularly to a low-oxygen rare earth sintered magnet and a quenched alloy containing a trace amount of W in a crystal grain boundary phase.
  • Low oxygen content magnet manufacturing process reduce the oxygen content in the magnet to deteriorate the sintering performance and deteriorate the coercive force as much as possible;
  • the raw material manufacturing process the raw material alloy represented by the entrainment method, at least a part of which is manufactured by quenching method;
  • the object of the present invention is to overcome the deficiencies of the prior art and to provide a W-containing R-Fe-B-Cu based sintered magnet which is uniformly segregated and pinned in a crystal grain boundary by a trace amount of W-pinned crystals ( Pinning effect)
  • the migration of grain boundaries can effectively prevent the occurrence of abnormal grain growth (AGG) and achieve significant improvement.
  • a W-containing R-Fe-B-Cu based sintered magnet comprising a R 2 Fe 14 B-type main phase, wherein R is at least one rare earth element comprising Nd or Pr, characterized in that:
  • the crystal grain boundary of the rare earth magnet has a W-rich region having a W content of 0.004 at% or more and 0.26 at% or less, and the W-rich region has a uniformly dispersed distribution in the crystal grain boundary phase, and accounts for 5.0% by volume to 11.0% by volume of the sintered magnet.
  • the crystal grain boundary is a portion other than the main phase (R 2 Fe 14 B) in the sintered magnet.
  • the sintered magnet is made from a raw material comprising the following components:
  • X 5.0 at% or less, X is at least one element selected from the group consisting of Al, Si, Ga, Sn, Ge, Ag, Au, Bi, Mn, Nb, Zr or Cr, and includes Nb and/or Zr at X. When the total content of Nb and Zr is below 0.20 at%,
  • the balance is 0at% to 20at% of Co, Fe, and unavoidable impurities
  • the impurity includes O, and the sintered magnet has an O content of 0.1 at% to 1.0 at%.
  • the at% described in the present invention is an atomic percentage.
  • the rare earth element referred to in the present invention is at least one selected from the group consisting of Nd, Pr, Dy, Tb, Ho, La, Ce, Pm, Sm, Eu, Gd, Er, Tm, Yb, Lu or lanthanum.
  • ICP-MS inductively coupled plasma mass spectrometer
  • FE-EPMA field emission electron probe microscopic analysis
  • ICP-MS model 7700x, Agilent
  • FE ⁇ EPMA Model 8530F, JEOL
  • an amorphous phase and an isotropic quenching are generated in a quenched alloy obtained from a raw material obtained by a high-melting-point metal such as Zr, Hf, Mo, V, W, or Nb (which is limited to about 0.25 at%).
  • a high-melting-point metal such as Zr, Hf, Mo, V, W, or Nb
  • the present invention contains a trace amount of W, that is, a content of 0.03 at% or less, and since it is a non-magnetic element, the dilution effect is small.
  • the quenched magnet alloy contains almost no amorphous phase and isotropic quenching phase. Therefore, in the present invention, the trace amount of W does not lower Br and (BH)max at all, but also Br, (BH) Max is improved.
  • W has a large solid solution limit in the main raw material Fe, so that a trace amount of W in the melt can be uniformly dissolved. Since W has different ionic radii and electronic structures from the rare earth elements, iron and boron of the main constituent elements, there is almost no W in the main phase of R 2 Fe 14 B, and in the cooling process of the melt, along with R 2 The precipitation of the main phase of Fe 14 B is concentrated to the crystal grain boundaries.
  • the composition of the rare earth is more than the composition of the main phase alloy, so the crystal grain boundary rare earth (R) content is large, that is, the R-rich phase (also known as the Nd-rich phase) contains Most of the W (tested by FE-EPMA test, most of the W contained in traces exist in the crystal grain boundaries), after W dissolves into the grain boundary, the affinity of W element with rare earth elements and Cu is poor, crystal The W in the rare earth-rich phase is precipitated and separated during the cooling process.
  • the solidification temperature of the grain boundary reaches about 500-700 °C, it is not easy to form large particles due to the slow diffusion rate of B, C, and O.
  • the rare earth intermetallic compound R 2 Fe 14 B is also precipitated minutely and uniformly, preventing the occurrence of AGG and improving the squareness (SQ) of the produced magnet. Furthermore, since Cu distributed in the grain boundary increases the low melting point liquid phase, the increase of the low melting point liquid phase promotes the migration of W. As can be seen from the EMPA result of Fig. 3, in the present invention, W is distributed in the grain boundary. It is fairly uniform, and the distribution range exceeds the distribution range of the Nd-rich phase, completely covering the entire Nd-rich phase, which can be considered as evidence that W plays a pinning effect and hinders grain growth.
  • a graphite crucible electrolytic cell a barrel-shaped graphite crucible is used as an anode, a tungsten (W) rod is arranged on the crucible axis as a cathode, and a rare earth is used to collect rare earth at the bottom.
  • a rare earth element such as Nd
  • a small amount of W is inevitably mixed therein.
  • other high-melting-point metals such as molybdenum (Mo) may be used as the cathode, and the rare earth metal may be obtained by using molybdenum rhenium to collect the rare earth metal.
  • W may also be an impurity of a raw material (such as pure iron, rare earth metal, B, etc.), and the raw material used in the present invention is selected according to the content of impurities in the raw material; of course, the W content may also be selected.
  • Raw materials such as pure iron, rare earth metals, B, etc.
  • having the detection limit of the equipment which may be regarded as not containing W
  • W W
  • Table 1 shows the W element content of metal Nd in different workshops in different places.
  • R 12 at% to 15.2 at%
  • B 5 at% to 8 at%
  • the balance of 0 at% to 20 at% of the content range of Co, Fe, etc. are conventional choices in the industry, and therefore In the examples, the range of contents of R, B, Fe, and Co was not tested and verified.
  • the present invention requires that the entire manufacturing process of the magnet be completed in a low-oxygen environment, and the O content is controlled to be 0.1 at% to 1.0 at% to obtain the claimed effect of the present invention.
  • the oxygen content is high (2500 ppm).
  • the above rare earth magnet can reduce the generation of AGG, while the rare earth magnet having a lower oxygen content (below 2500 ppm) has a good magnetic property, but is easy to produce AGG, and the present invention contains only a very small amount of W and a small amount of Cu.
  • the effect of reducing AGG is also achieved in a low oxygen content magnet.
  • the X content is below 2.0 at%.
  • the sintered magnet is obtained by the step of preparing a sintered magnet raw material melt into a sintered magnet alloy at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec; a step of coarsely pulverizing a sintered magnet with an alloy and then finely pulverizing it to obtain a fine powder; obtaining a formed body by a magnetic field forming method, and sintering the formed body at a temperature of 900 ° C to 1100 ° C in a vacuum or an inert gas; obtain.
  • the sintering temperature using a temperature of from 900 ° C to 1100 ° C is a conventional choice in the industry, and therefore, in the examples, the range of the sintering temperature was not tested and verified.
  • the degree of dispersion of W is increased in the grain boundary to make the squareness exceed 95%, and the temperature resistance of the magnet is improved.
  • the degree of dispersion of W is improved mainly by controlling the cooling rate of the melt.
  • the sintered magnet has a B content of preferably 5.0 at% to 6.5 at%. Since an excessive amount of B easily reacts with W, the boride phase formed, the hardness of these boride phases is very high, very hard, and the processability is drastically deteriorated, and at the same time, the boride phase (WB 2 phase) is formed due to the formation of large particles. ), the effect of W on the grain boundary migration in the grain boundary is also affected. Therefore, appropriately reducing the B content can reduce the formation of the boride phase and give full play to the pinning effect of W. effect.
  • the sintered magnet has an Al content of preferably 0.8 at% to 2.0 at%.
  • Al 0.8 to 2.0 at%
  • a W-containing R 6 T 13 is formed.
  • the unavoidable impurities mentioned in the present invention also include a small amount of C, N, S, P and other impurities which are inevitably mixed in the raw material or in the manufacturing process, and therefore, the places mentioned in the present invention
  • the coarse pulverization is a step of pulverizing the alloy for sintering magnets to obtain a coarse powder
  • the fine pulverization is a step of pulverizing the coarse powder, and further includes removing the particle diameter from the finely pulverized powder. At least a part of 1.0 ⁇ m or less, thereby reducing the volume of the powder having a particle diameter of 1.0 ⁇ m or less to the entire powder The process of 10% or less of the final volume.
  • the step of subjecting the sintered magnet to RH (heavy rare earth element) grain boundary diffusion treatment is further included.
  • the grain boundary diffusion is generally carried out at a temperature of from 700 ° C to 1050 ° C. This temperature range is a conventional choice in the industry, and therefore, in the examples, the above temperature range was not tested and verified.
  • the magnet of the present invention can achieve very high performance by the grain boundary diffusion of RH, and a dramatic improvement effect can be obtained.
  • the RH is selected from at least one of Dy or Tb.
  • the method further comprises the step of aging treatment: aging the sintered magnet at a temperature of from 400 ° C to 650 ° C.
  • the method further includes a two-stage aging treatment: after the first stage heat treatment of the sintered magnet at a temperature of 800 ° C to 950 ° C for 1 hour to 2 hours, and then the sintered magnet is 450 A secondary heat treatment is carried out at a temperature of from ° C to 660 ° C for 1 hour to 4 hours.
  • the sintered magnet has an O content of from 0.1 at% to 0.5 at%.
  • the ratio of O, W and Cu reaches the optimum ratio, the heat resistance of the sintered magnet is high, the stability of the magnet under dynamic working conditions is high, and when there is no AGG, the oxygen content is low, Hcj Raise.
  • the sintered magnet has a Ga content of 0.05 at% to 0.8 at%.
  • Another object of the present invention is to provide a quenched alloy for a R-Fe-B-Cu based sintered magnet containing W.
  • a quenched alloy for a R-Fe-B-Cu based sintered magnet containing W characterized in that the crystal grain boundary of the quenched alloy has a W-rich region having a W content of 0.004 at% or more and 0.26 at% or less.
  • the W-rich region has a uniformly dispersed distribution in the crystalline grain boundaries and accounts for at least 50% by volume of the crystalline grain boundaries.
  • the present invention has the following characteristics:
  • the present invention improves the performance of the magnets based on the three mass-produced magnet technologies in the background art, and studies on trace elements in detail, and improves the SQ, Hcj, Br, (BH) of the magnet by suppressing the AGG during sintering. Max, the results show that a small amount of W-pinned crystals in the grain boundary of the uniform pinning effect grain boundary migration, can effectively prevent the occurrence of abnormal grain growth (AGG), and can obtain significant improvement .
  • the content of W contained in the present invention is extremely small and uniformly dispersed, and the high-melting-point large-particle metal boride phase hardly occurs, and even if it appears, it is only present in a small amount, so there is almost no processing deterioration.
  • the present invention contains a trace amount of W (non-magnetic element), that is, a content of 0.03 at% or less, and the dilution effect is small. Further, the quenched magnet alloy does not contain an amorphous phase and an isotropic quenching phase at all. It is detected by FE-EPMA that most of the W contained in a trace amount exists in the crystal grain boundary. Therefore, in the present invention, the trace amount of W does not decrease Br and (BH)max at all, but also increases Br and (BH)max.
  • W non-magnetic element
  • the composition of the present invention contains a trace amount of Cu and W, so that the high melting point in the grain boundary [such as WB 2 phase (melting point 2365 ° C), etc.] intermetallic phase cannot be formed, but produces more RCu (melting point 662). °C), RCu 2 (melting point 840 ° C), Nd-Cu eutectic alloy (melting point 492 ° C) and other low melting point phase, as a result, almost all of the crystal grain boundaries except the W phase at the grain boundary diffusion temperature, crystal The efficiency of the boundary diffusion is excellent, the squareness and the coercive force are increased to an unprecedented extent, and the squareness is more than 99%, thereby obtaining a high-performance magnet with good heat resistance.
  • the WB 2 phase herein includes a WFeB alloy, a WFe alloy, a WB alloy, and the like.
  • Figure 1 is a schematic diagram of the principle of the grain boundary migration of the Pinning effect.
  • Fig. 2 is a result of EPMA detection of the quenched alloy sheet of Example 3 of the first embodiment.
  • Fig. 3 is a result of EPMA detection of the sintered magnet of Example 3 of the first embodiment.
  • the BHH, magnetic performance evaluation process, and AGG measurement mentioned in each embodiment are defined as follows:
  • BHH is the sum of (BH)max and Hcj and is one of the evaluation criteria for the comprehensive performance of the magnet.
  • Magnetic performance evaluation process The sintered magnet was magnetically tested using the NIM ⁇ 10000H BH bulk rare earth permanent magnet non-destructive measurement system of China Metrology Institute.
  • AGG measurement The sintered magnet was polished in a direction perpendicular to the orientation direction, and the AGG mentioned in the present invention was a crystal grain having a particle diameter exceeding 40 ⁇ m per the average AGG amount included in 1 cm 2 .
  • the detection limit of FE-EPMA detection mentioned in each example is about 100 ppm, and the detection conditions are as follows:
  • the maximum resolution of the FE ⁇ EPMA device is 3 nm, and the resolution can reach 50 nm under the above detection conditions.
  • Raw material preparation process preparation of 99.5% purity Nd, Dy, industrial Fe-B, industrial pure Fe, purity 99.9% Co and purity 99.5% Cu, Al, purity 99.999% W, atomic percentage at% Formulated.
  • the W content in selected Nd, Dy, Fe, B, Al, Cu, and Co is below the detection limit of the existing device, and the source of W is additionally added. W metal.
  • Each serial number group was prepared according to the elemental composition in Table 2, and 100 kg of raw materials were weighed and prepared.
  • Casting process Ar gas is introduced into the melting furnace after vacuum melting to bring the gas pressure to 50,000 Pa, and then cast by a single roll quenching method to obtain a quenched alloy at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec. The quenched alloy was heat treated at 600 ° C for 60 minutes and then cooled to room temperature.
  • the Cu, Nd, and W components of the quenched alloy prepared in Example 3 were subjected to FE-EPMA (Field Emission Electron Probe Microanalysis) [JEOL, 8530F], and the results are shown in FIG. It can be observed that W is distributed in the R-rich phase with a higher dispersion.
  • FE-EPMA Field Emission Electron Probe Microanalysis
  • the FE-EPMA test was performed on the quenched alloy sheets of Examples 2, 3, 4, 5 and 6, and the W-rich region was uniformly dispersed in the crystal grain boundaries and accounted for at least 50% by volume of the grain boundaries of the alloy, wherein The W-rich region is a region having a W content of 0.004 at% or more and 0.26 at% or less.
  • Hydrogen breaking pulverization process a hydrogen-breaking furnace in which a quenching alloy is placed is evacuated at room temperature, and then a hydrogen gas having a purity of 99.5% is introduced into a hydrogen-breaking furnace to a pressure of 0.1 MPa, and after standing for 2 hours, the temperature is raised while evacuating. The vacuum was evacuated at a temperature of 500 ° C, and then cooled, and the powder after the pulverization of hydrogen was taken out.
  • the powder after the hydrogen pulverization was subjected to jet milling at a pressure of 0.4 MPa in a nitrogen gas atmosphere having an oxidizing gas content of 100 ppm or less to obtain a fine powder, and the average particle size of the fine powder was 4.5 ⁇ m.
  • Oxidizing gas refers to oxygen or moisture.
  • the finely pulverized fine powder (30% by weight based on the total weight of the fine powder) was classified by a classifier to remove particles having a particle diameter of 1.0 ⁇ m or less, and the classified fine powder was mixed with the remaining unfractionated fine powder.
  • the volume of the powder having a particle diameter of 1.0 ⁇ m or less is reduced to 10% or less of the entire volume of the powder.
  • Methyl octanoate was added to the powder after the jet mill pulverization, and the methyl octanoate was added in an amount of 0.2% by weight of the mixed powder, followed by thorough mixing with a V-type mixer.
  • Magnetic field forming process Using a right-angle oriented magnetic field forming machine, the above-mentioned methyl octanoate-added powder was once formed into a cube having a side length of 25 mm in a 1.8 T orientation magnetic field at a molding pressure of 0.4 ton/cm 2 . After one forming, it demagnetizes in a magnetic field of 0.2T.
  • Sintering process Each formed body is moved to a sintering furnace for sintering, and the sintering is maintained at a temperature of 200 ° C and 800 ° C for 2 hours under a vacuum of 10 -3 Pa, and then sintered at a temperature of 1030 ° C for 2 hours, and then passed through. After the Ar gas was introduced to bring the gas pressure to 0.1 MPa, it was cooled to room temperature.
  • Heat treatment process The sintered body was heat-treated at a temperature of 460 ° C for 1 hour in high-purity Ar gas, and then cooled to room temperature and taken out.
  • the heat-treated sintered body is processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • the magnets of the sintered bodies of Examples 1 to 7 were directly subjected to magnetic property detection to evaluate their magnetic properties.
  • the evaluation results of the examples of the magnets are shown in Tables 3 and 4:
  • the amorphous phase and the isotropic phase in Table 3 are the amorphous phases and the isotropy in the quenched alloy. Sexual phase.
  • the W-rich phase in Table 3 is a region of 0.004 at% or more and 0.26 at% or less.
  • the W content in selected Nd, Pr, Tb, Fe, B, Al, and Cu is below the detection limit of the existing equipment, and the source of W is additionally added. W metal.
  • Each serial number group was prepared according to the elemental composition in Table 5, and 100 kg of raw materials were weighed and prepared.
  • Casting process Ar gas is introduced into a melting furnace after vacuum melting to bring the gas pressure to 30,000 Pa, and then casting is performed by a single roll quenching method, and a quenched alloy is obtained at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec. The quenched alloy was heat treated at 600 ° C for 60 minutes and then cooled to room temperature.
  • the FE-EPMA was tested on the quenched alloy sheets of Examples 2 to 7.
  • the W-rich region was uniformly dispersed in the crystal grain boundaries and accounted for at least 50% by volume of the alloy crystal grain boundaries, wherein the W-rich region was W content. It is an area of 0.004 at% or more and 0.26 at% or less.
  • Hydrogen breaking and pulverizing process a hydrogen breaking furnace in which a quenching alloy is placed is evacuated at room temperature, and then a hydrogen gas having a purity of 99.5% is introduced into the hydrogen breaking furnace to a pressure of 0.1 MPa, and after standing for 125 minutes, the temperature is raised while vacuuming. The vacuum was evacuated at a temperature of 500 ° C for 2 hours, and then cooled, and the hydrogen-crushed powder was taken out.
  • the pressure in the pulverizing chamber is The powder after the hydrogen pulverization was subjected to jet mill pulverization under a pressure of 0.41 MPa to obtain a fine powder, and the average particle size of the fine powder was 4.30 ⁇ m.
  • Oxidizing gas refers to oxygen or moisture.
  • Methyl octanoate was added to the powder after the jet mill pulverization, and the methyl octanoate was added in an amount of 0.25% by weight of the mixed powder, and then thoroughly mixed by a V-type mixer.
  • Magnetic field forming process Using a right angle oriented type magnetic field forming machine, the above-mentioned methyl octanoate-added powder was once formed into a cube having a side length of 25 mm in a 1.8 T orientation magnetic field at a molding pressure of 0.3 ton/cm 2 . After one forming, it demagnetizes in a magnetic field of 0.2T.
  • each formed body is moved to a sintering furnace for sintering, and the sintering is maintained at a temperature of 200 ° C and 800 ° C for 3 hours under a vacuum of 10 -3 Pa, and then sintered at a temperature of 1020 ° C for 2 hours. After the Ar gas was introduced to bring the gas pressure to 0.1 MPa, it was cooled to room temperature.
  • Heat treatment process The sintered body was heat-treated at a temperature of 620 ° C for 1 hour in high-purity Ar gas, and then cooled to room temperature and taken out.
  • the heat-treated sintered body is processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • the magnets of the sintered bodies of Examples 1 to 8 were directly subjected to magnetic property measurement to evaluate their magnetic properties.
  • the evaluation results of the examples of the magnets are shown in Tables 6 and 7:
  • the amorphous phase and the isotropic phase in Table 6 are the amorphous phase and the isotropic phase in the quenched alloy.
  • the W-rich phase in Table 6 is a region of 0.004 at% or more and 0.26 at% or less.
  • FE-EPMA detection (JEOL, 8530F) was performed on Examples 2 to 7, and as a result of the detection, it was observed that W uniformly pinned the grain boundary with a high degree of dispersion. Migration to prevent the formation of AGG.
  • the W content in the selected Nd, Fe, B, Cu, and Co raw materials is below the detection limit of the existing equipment, and the source of W is the additionally added W metal.
  • the ingredients were weighed and prepared, and 700 Kg of raw materials were prepared.
  • Smelting process The prepared raw materials are placed in a crucible made of alumina, and vacuum-melted at a temperature of 1500 ° C or lower in a vacuum of 10 ⁇ 2 Pa in a high-frequency vacuum induction melting furnace.
  • Casting process Ar gas is introduced into the melting furnace after vacuum melting to bring the gas pressure to 50,000 Pa, and then cast by a single roll quenching method to obtain a quenched alloy at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec. The quenched alloy was heat treated at 600 ° C for 60 minutes and then cooled to room temperature.
  • the FE-EPMA test is performed on the quenched alloy sheet.
  • the W-rich region has a uniform dispersion distribution in the crystal grain boundary and accounts for at least 50% by volume of the crystal grain boundary of the alloy.
  • the W-rich region has a W content of 0.004 at% or more. Area below 0.26at%.
  • Hydrogen breaking and pulverizing process the hydrogen quenching furnace in which the quenching alloy is placed is evacuated at room temperature, and then hydrogen gas having a purity of 99.5% is introduced into the hydrogen breaking furnace to a pressure of 0.1 MPa, and after standing for 97 minutes, the temperature is raised while vacuuming. The vacuum was evacuated at a temperature of 500 ° C for 2 hours, and then cooled, and the hydrogen-crushed powder was taken out.
  • the powder after the hydrogen pulverization is divided into 7 parts, and each of the powders is pulverized under the pressure of the pulverization chamber at a pressure of 0.42 MPa in an atmosphere having an oxidizing gas content of 10 to 3,000 ppm or less.
  • the powder was subjected to jet mill pulverization to obtain a fine powder having an average particle size of 4.51 ⁇ m.
  • Oxidizing gas refers to oxygen or moisture.
  • Methyl octanoate was added to each of the powders pulverized by the jet mill, and the amount of methyl octanoate added was 0.1% by weight of the powder after mixing, and then thoroughly mixed by a V-type mixer.
  • Magnetic field forming process The above-mentioned powder added with methyl octanoate was formed into a side length of 25 mm at a molding pressure of 0.2 ton/cm 2 in a 1.8 T oriented magnetic field using a magnetic field forming machine of a right angle orientation type. The cube is demagnetized in a magnetic field of 0.2 T after one forming.
  • each formed body is moved to a sintering furnace for sintering, and the sintering is maintained at a temperature of 200 ° C and 700 ° C for 2 hours under a vacuum of 10 -3 Pa, and then sintered at a temperature of 1020 ° C for 2 hours, and then passed through.
  • the Ar gas was introduced to bring the gas pressure to 0.1 MPa, it was cooled to room temperature.
  • Heat treatment process The sintered body is subjected to a primary heat treatment at 900 ° C for 1 hour in a high-purity Ar gas, and then subjected to a secondary heat treatment at a temperature of 500 ° C for 1 hour, and after cooling to room temperature, it is taken out.
  • the heat-treated sintered body is processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • the sintered magnet was placed in a 150 ° C environment for 30 min, then cooled naturally to room temperature, and the magnetic flux was measured. The measured results were compared with the measured data before heating to calculate the flux decay before and after heating. rate.
  • the W-rich phase in Table 9 is a region of 0.004 at% or more and 0.26 at% or less.
  • FE-EPMA detection (JEOL, 8530F) was performed on Examples 2 to 6, and as a result of the detection, it was also observed that W was uniformly pinned with a high degree of dispersion. The migration of the boundary prevents the formation of AGG.
  • Raw material preparation process preparation of 99.5% purity Nd, Dy, industrial Fe-B, industrial pure Fe, purity 99.9% Co and purity 99.5% Cu, Al, purity 99.999% W, atomic percentage at% Formulated.
  • the W content in the selected Nd, Dy, B, Al, Cu, Co, and Fe is below the detection limit of the existing device, and the source of W is additionally added.
  • W metal. Its content is shown in Table 11:
  • Each serial number group was prepared according to the elemental composition in Table 11, and 100 kg of raw materials were weighed and prepared.
  • Casting process Ar gas is introduced into the melting furnace after vacuum melting to bring the gas pressure to 20,000 Pa, and then cast by a single roll quenching method to obtain a quenched alloy at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec. The quenched alloy was subjected to heat treatment at 800 ° C for 10 minutes, and then cooled to room temperature.
  • the FE-EPMA was tested on the quenched alloy sheets of Examples 1 to 7, and the W-rich region was uniformly dispersed in the crystal grain boundaries and accounted for at least 50% by volume of the alloy crystal grain boundaries, wherein the W-rich region was W content. It is an area of 0.004 at% or more and 0.26 at% or less.
  • Hydrogen breaking and pulverizing process the hydrogen quenching furnace in which the quenching alloy is placed is evacuated at room temperature, and then hydrogen gas having a purity of 99.5% is introduced into the hydrogen breaking furnace to a pressure of 0.1 MPa, and after being left for 120 minutes, the temperature is raised while vacuuming. The vacuum was evacuated at a temperature of 500 ° C for 2 hours, and then cooled, and the hydrogen-crushed powder was taken out.
  • the powder after the hydrogen pulverization is subjected to jet milling at a pressure of a pulverization chamber pressure of 0.6 MPa in an atmosphere having an oxidizing gas content of 100 ppm or less to obtain a fine powder, and the average particle size of the fine powder is 4.5 ⁇ m.
  • Oxidizing gas refers to oxygen or moisture.
  • the finely pulverized fine powder (30% by weight based on the total weight of the fine powder) was classified by a classifier to remove the particles having a particle diameter of 1.0 ⁇ m or less, and the classified fine powder was mixed with the remaining unfractionated fine powder. After mixing In the fine powder, the volume of the powder having a particle diameter of 1.0 ⁇ m or less is reduced to 2% or less of the entire volume of the powder.
  • Methyl octanoate was added to the powder after the jet mill pulverization, and the methyl octanoate was added in an amount of 0.2% by weight of the mixed powder, followed by thorough mixing with a V-type mixer.
  • Magnetic field forming process Using a right-angle oriented magnetic field forming machine, the above-mentioned methyl octanoate-added powder was once formed into a cube having a side length of 25 mm in a 1.8 T orientation magnetic field at a molding pressure of 0.2 ton/cm 2 . After one forming, it demagnetizes in a magnetic field of 0.2T.
  • each formed body is moved to a sintering furnace for sintering, and the sintering is maintained at a temperature of 200 ° C and 800 ° C for 2 hours under a vacuum of 10 -3 Pa, and then sintered at a temperature of 1040 ° C for 2 hours, followed by After the Ar gas was introduced to bring the gas pressure to 0.1 MPa, it was cooled to room temperature.
  • Heat treatment process The sintered body was heat-treated at a temperature of 400 ° C for 1 hour in high-purity Ar gas, and then cooled to room temperature and taken out.
  • the heat-treated sintered body is processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • the magnets of the sintered bodies of Examples 1 to 7 were directly subjected to magnetic property detection to evaluate their magnetic properties.
  • the evaluation results of the examples of the magnets are shown in Table 12 and Table 13:
  • the amorphous phase and the isotropic phase in Table 12 are the amorphous phase and the isotropic phase in the quenched alloy.
  • the W-rich phase in Table 12 is a region of 0.004 at% or more and 0.26 at% or less.
  • Raw material preparation process preparation of 99.5% purity Nd, Dy, industrial Fe-B, industrial pure Fe, purity 99.9% Co and purity 99.5% Cu, Al, purity 99.999% W, atomic percentage at% Formulated.
  • the W content in the selected Nd, Dy, B, Al, Cu, Co, and Fe is below the detection limit of the existing device, and the source of W is additionally added.
  • W metal. The content of each element is shown in Table 14:
  • Each serial number group was prepared according to the elemental composition in Table 14, and 100 kg of raw materials were weighed and prepared.
  • Casting process Ar gas is introduced into the melting furnace after vacuum melting to bring the gas pressure to 50,000 Pa, and then cast by a single roll quenching method to obtain a quenched alloy at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec. The quenched alloy was heat treated at 700 ° C for 5 minutes and then cooled to room temperature.
  • Hydrogen breaking and pulverizing process the hydrogen quenching furnace in which the quenching alloy is placed is evacuated at room temperature, and then hydrogen gas having a purity of 99.5% is introduced into the hydrogen breaking furnace to a pressure of 0.1 MPa, and after being left for 120 minutes, the temperature is raised while vacuuming. The vacuum was evacuated at a temperature of 600 ° C for 2 hours, and then cooled, and the powder after the pulverization of hydrogen was taken out.
  • the powder after the hydrogen pulverization is subjected to jet milling at a pressure of a pulverization chamber pressure of 0.5 MPa in an atmosphere having an oxidizing gas content of 100 ppm or less to obtain a fine powder, and the average particle size of the fine powder is 5.0 ⁇ m.
  • Oxidizing gas refers to oxygen or moisture.
  • the finely pulverized fine powder (30% by weight based on the total weight of the fine powder) was sieved to remove the particles having a particle diameter of 1.0 ⁇ m or less, and the fine powder after the sieving was mixed with the remaining unsifted fine powder.
  • the volume of the powder having a particle diameter of 1.0 ⁇ m or less is reduced to 10% or less of the entire volume of the powder.
  • Methyl octanoate was added to the powder after the jet mill pulverization, and the methyl octanoate was added in an amount of 0.2% by weight of the mixed powder, followed by thorough mixing with a V-type mixer.
  • Magnetic field forming process Using a right-angle oriented magnetic field forming machine, the above-mentioned methyl octanoate-added powder was once formed into a cube having a side length of 25 mm in a 1.8 T orientation magnetic field at a molding pressure of 0.2 ton/cm 2 . After one forming, it demagnetizes in a magnetic field of 0.2T.
  • each formed body is moved to a sintering furnace for sintering, and the sintering is maintained at a temperature of 200 ° C and 800 ° C for 2 hours under a vacuum of 10 -3 Pa, and then sintered at a temperature of 1060 ° C for 2 hours, and then passed through. After the Ar gas was introduced to bring the gas pressure to 0.1 MPa, it was cooled to room temperature.
  • Heat treatment process The sintered body was heat-treated at a temperature of 420 ° C for 1 hour in high-purity Ar gas, and then cooled to room temperature and taken out.
  • the heat-treated sintered body is processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • the magnets of the sintered bodies of Examples 1 to 7 were directly subjected to magnetic property detection to evaluate their magnetic properties.
  • the evaluation results of the examples of the magnets are shown in Table 15:
  • the amorphous phase and the isotropic phase in Table 15 are the amorphous phase and the isotropic phase in the quenched alloy.
  • the W-rich phase in Table 15 is a region of 0.004 at% or more and 0.26 at% or less.
  • Each group of sintered bodies obtained in Example 1 was separately processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • Grain boundary diffusion treatment process the magnets processed by each group of sintered bodies are washed, and after the surface is cleaned, the raw materials prepared by using Dy oxide and Tb fluoride in a ratio of 3:1 are sprayed on the magnets in a comprehensive manner.
  • the coated magnet was dried and heat-dissipated at a temperature of 850 ° C for 24 hours in a high-purity Ar gas atmosphere.
  • the trace W in the present invention produces very minute W crystals in the crystal grain boundaries and does not inhibit the diffusion of Dy and Tb, so the diffusion speed is very fast. Further, since an appropriate amount of Cu is contained, an Nd-rich phase having a low melting point is formed, and an effect of further promoting diffusion can be exhibited. Therefore, the magnet of the present invention can achieve very high performance by diffusion of grain boundaries of Dy and Tb.
  • Nd, Dy, Tb having a purity of 99.9%, B having a purity of 99.9%, Fe having a purity of 99.9%, and Cu, Co, Nb, Al, and Ga having a purity of 99.5% are prepared at an atomic percentage at%.
  • the W content in selected Dy, Tb, Fe, B, Cu, Co, Nb, Al, and Ga is below the detection limit of the existing equipment, and the selected one is used.
  • Nd contains W, and the content of W element is 0.01 at%.
  • Each serial number group was prepared according to the elemental composition in Table 18, and 100 kg of raw materials were weighed and prepared.
  • Casting process Ar gas is introduced into the melting furnace after vacuum melting to bring the gas pressure to 35,000 Pa, and then casting is performed by a single roll quenching method, and a quenched alloy is obtained at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec. The quenched alloy was heat treated at 550 ° C for 10 minutes and then cooled to room temperature.
  • Hydrogen breaking and pulverizing process the hydrogen-dissolving furnace in which the quenching alloy is placed is evacuated at room temperature, and then hydrogen gas having a purity of 99.5% is introduced into the hydrogen breaking furnace to a pressure of 0.085 MPa, and after standing for 160 minutes, the temperature is raised while vacuuming. The vacuum was applied at a temperature of 520 ° C, and then cooled, and hydrogen was taken out to break the pulverized powder.
  • the powder after the hydrogen pulverization was subjected to jet milling at a pressure of 0.42 MPa in an atmosphere having an oxidizing gas content of 10 ppm or less to obtain a fine powder, and the average particle size of the fine powder was 4.28 ⁇ m.
  • Oxidizing gas refers to oxygen or moisture.
  • Methyl octanoate was added to the powder after the jet mill pulverization, and the methyl octanoate was added in an amount of 0.25% by weight of the mixed powder, and then thoroughly mixed by a V-type mixer.
  • Magnetic field forming process Using a right angle oriented type magnetic field forming machine, the above-mentioned methyl octanoate-added powder was once formed into a cube having a side length of 25 mm in a 1.8 T orientation magnetic field at a molding pressure of 0.3 ton/cm 2 . After one forming, it demagnetizes in a magnetic field of 0.2T.
  • each formed body is moved to a sintering furnace for sintering, and the sintering is maintained at a temperature of 300 ° C and 800 ° C for 3 hours under a vacuum of 10 -3 Pa, and then sintered at a temperature of 1030 ° C for 2 hours, and then passed through. After the Ar gas was introduced to bring the gas pressure to 0.1 MPa, it was cooled to room temperature.
  • Heat treatment process The sintered body was heat-treated at 600 ° C for 2 hours in high-purity Ar gas, and then cooled to room temperature and taken out.
  • the heat-treated sintered body is processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • the magnets of the sintered bodies of Examples 1 to 8 were directly subjected to magnetic property measurement to evaluate their magnetic properties.
  • the evaluation results of the examples of the magnets are shown in Tables 19 and 20:
  • the amorphous phase and the isotropic phase in Table 19 are the amorphous phase and the isotropic phase in the quenched alloy.
  • the W-rich phase in Table 19 is a region of 0.004 at% or more and 0.26 at% or less.
  • O, C and N which will be the above O, C and N.
  • the content of each element in the magnet is controlled to be 0.1 to 0.5 at%, 0.4 at% or less, and 0.2 at% or less.
  • FE-EPMA detection (JEOL, 8530F) was performed on Examples 1 to 8, and as a result of the detection, it was observed that W uniformly pinned the grain boundary with a high degree of dispersion. Migration to prevent the formation of AGG.
  • the W content in selected Dy, Gd, Tb, Fe, B, Cu, Co, Nb, Al, and Ga is below the detection limit of the existing equipment.
  • the selected Nd contains W and the content of W element is 0.01 at%.
  • Each serial number group was prepared according to the elemental composition in Table 21, and 100 kg of raw materials were weighed and prepared.
  • Casting process Ar gas is introduced into the melting furnace after vacuum melting to bring the gas pressure to 45,000 Pa, and then casting is performed by a single roll quenching method, and a quenched alloy is obtained at a cooling rate of 10 2 ° C / sec to 10 4 ° C / sec. The quenched alloy was heat treated at 800 ° C for 5 minutes and then cooled to room temperature.
  • Hydrogen breaking and pulverizing process a hydrogen breaking furnace in which a quenching alloy is placed is evacuated at room temperature, and then a hydrogen gas having a purity of 99.5% is introduced into a hydrogen breaking furnace to a pressure of 0.09 MPa, and after standing for 150 minutes, the temperature is raised while vacuuming. The vacuum was evacuated at a temperature of 600 ° C, and then cooled, and the pulverized powder was taken out by hydrogen.
  • the powder after the hydrogen pulverization is subjected to jet milling at a pressure of 0.5 MPa in an atmosphere having an oxidizing gas content of 30 ppm or less to obtain a fine powder, and the average particle size of the fine powder is 4.1 ⁇ m.
  • Oxidizing gas refers to oxygen or moisture.
  • the aluminum stearate was added to the powder after the jet mill pulverization, and the amount of the aluminum stearate added was 0.05% by weight of the mixed powder, followed by thorough mixing with a V-type mixer.
  • Magnetic field forming process The above-mentioned aluminum stearate-added powder was once formed into a side length of 25 mm in a 1.8 T orientation magnetic field at a molding pressure of 0.3 ton/cm 2 using a right angle orientation type magnetic field molding machine. The cube is demagnetized in a magnetic field of 0.2 T after one forming.
  • each formed body is moved to a sintering furnace for sintering, and the sintering is maintained at a temperature of 200 ° C and 800 ° C for 3 hours under a vacuum of 10 -3 Pa, and then sintered at a temperature of 1050 ° C for 2 hours, and then passed through. After the Ar gas was introduced to bring the gas pressure to 0.1 MPa, it was cooled to room temperature.
  • Heat treatment process The sintered body was heat-treated at a temperature of 480 ° C for 2 hours in high-purity Ar gas, and then cooled to room temperature and taken out.
  • the heat-treated sintered body is processed into A magnet having a thickness of 5 mm has a direction of magnetic field orientation of 5 mm.
  • the magnets of the sintered bodies of Examples 1 to 5 were directly subjected to magnetic property measurement to evaluate their magnetic properties.
  • the evaluation results of the examples of the magnets are shown in Table 22 and Table 23:
  • the amorphous phase and the isotropic phase in Table 23 are the amorphous phase and the isotropic phase in the quenched alloy.
  • the W-rich phase in Table 23 is a region of 0.004 at% or more and 0.26 at% or less.
  • the applicant has learned through experiments that the content of Zr should also be controlled to be 0.2 at% or less.
  • FE-EPMA detection (JEOL, 8530F) was performed on Examples 1 to 5, and as a result of the detection, it was observed that W uniformly pinned the grain boundary with a high degree of dispersion. Migration to prevent the formation of AGG.

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Abstract

提供一种含W的R-Fe-B-Cu系烧结磁铁及急冷合金。烧结磁铁含有R 2Fe 14B型主相，所述的R为包括Nd或Pr的至少一种稀土元素，所述稀土磁铁的结晶晶界中含有0.004at%以上、0.26at%以下的富W区域，所述富W区域占所述烧结磁铁的5.0体积%-11.0体积%。该烧结磁铁通过微量的W钉扎结晶物在结晶晶界中偏析钉扎晶界的迁移，可以有效防止晶粒异常长大的产生，获得显著的改善效果。急冷合金的结晶晶界中含有0.004at%以上、0.26at%以下的富W区域，所述富W区域占所述结晶晶界的至少50体积%。

Description

一种含W的R‐Fe‐B‐Cu系烧结磁铁及急冷合金 技术领域
本发明涉及磁铁的制造技术领域，特别是涉及一种在结晶晶界相中含微量W的低氧含量稀土烧结磁铁及急冷合金。
背景技术
近年来，作为稀土烧结磁体(含R2Fe14B型主相)制法的3大新技术被快速运用于量产技术工序中，具体如下：
1、低氧含量磁体制造工序：尽可能降低磁铁中使烧结性能变差、矫顽力劣化的氧含量；
2、原料制造工序：以甩带法为代表制出的原料合金，其至少一部分使用急冷法制造；
3、通过添加微量Cu，可以在更宽的温度范围内进行热处理获得高矫顽力，并缓和矫顽力和冷却速度的依存性(来源于JP2720040等的公开报道)。
将上述3种量产新技术进行组合，通过结晶晶界中的富Nd相量增加和分散性提高的加成作用，可以较易达到非常高的性能。
然而，由于在低氧含量磁铁中添加了Cu，导致烧结过程中低熔点液相增加，在烧结性能明显提高的同时，容易发生晶粒异常长大(AGG)、并使方形度(SQ)显著降低的缺点。
发明内容
本发明的目的在于克服现有技术之不足，提供一种含W的R‐Fe‐B‐Cu系烧结磁铁，该烧结磁铁通过微量的W钉扎结晶物在结晶晶界中均一偏析钉扎(Pinning effect)晶界的迁移，可以有效防止晶粒异常长大(AGG)的产生，获得显著的改善效果。
本发明提供的技术方案如下：
一种含W的R‐Fe‐B‐Cu系烧结磁铁，所述烧结磁铁含有R2Fe14B型主相，所述的R为包括Nd或Pr的至少一种稀土元素，其特征在于：所述稀土磁铁的结晶晶界中具有W含量为0.004at％以上、0.26at％以下的富W区域，所述富W区域在所述结晶晶界相中呈均一分散的分布，并占所述烧结磁铁的5.0体积％～11.0体积％。
本发明中，结晶晶界为烧结磁铁中除主相(R2Fe14B)以外的部位。
在推荐的实施方式中，所述烧结磁铁由包括如下成分的原料制成：
R：12at％～15.2at％，
B：5at％～8at％，
W：0.0005at％～0.03at％，
Cu：0.05at％～1.2at％，
X：5.0at％以下、X为选自Al、Si、Ga、Sn、Ge、Ag、Au、Bi、Mn、Nb、Zr或Cr中的至少1种元素，在X包括Nb和/或Zr之时，Nb和Zr的总含量在0.20at％以下，
余量为0at％～20at％的Co、Fe、以及不可避免的杂质，
所述杂质包括O，且所述烧结磁铁的O含量为0.1at％～1.0at％。
本发明中所述的at％为原子百分比。
本发明所提及的稀土元素选自Nd、Pr、Dy、Tb、Ho、La、Ce、Pm、Sm、Eu、Gd、Er、Tm、Yb、Lu或钇元素中的至少一种。
由于受检测设备的限制，在以往的研究中，对于痕量元素的测试结果准确性难以保证。近年来，随着检测技术的提升，出现了更精准的检测设备，如电感耦合等离子质谱仪ICP‐MS、场发射电子探针显微分析FE‐EPMA等设备。其中，ICP‐MS(型号7700x，Agilent)可检测10ppb含量的元素。FE‐EPMA(型号8530F，JEOL)通过场发射电子枪，在大电流工作仍可保证极细的电子束，最高分辨率达到3nm，对于微区元素含量的检测限达到100ppm左右。
本发明与以往添加量较多的Zr、Hf、Mo、V、W、Nb等高熔点金属(多限定为0.25at％左右)原料所获得急冷合金中会产生非晶质相及各向同性急冷相，使结晶取向度变差、导致Br、(BH)max显著降低的趋势不同，本发明含有微量的W，即0.03at％以下的含有量，由于其为非磁性元素，稀释效果较少，且急冷后的磁体合金中几乎不含有非晶质相和各向同性急冷相，因此，本发明中W的微量含有完全不会降低Br、(BH)max，反而还会使Br、(BH)max提高。
从现有文献报道提供的状态来看，W在主要原料Fe中具有较大的固溶限，所以，熔融液中的微量W能均匀溶解。而由于W与主要构成元素的稀土元素、铁、硼的离子半径及电子构造不同，所以，R2Fe14B主相中几乎不存在W，W在熔融液的冷却过程中，随着R2Fe14B主相的析出，向结晶晶界浓缩。在原料配比组成时，按稀土类多于主相合金的成分进行设计，所以结晶晶界稀土(R)含量较多，也就是说，富R相(亦称为富Nd相)中含有绝大部分的W(经FE‐EPMA检测验证，微量含有的W绝大部分存在结晶晶界中)，W溶入晶界之后，由于W元素与稀土类元素、Cu的亲和性较差，晶界中富稀土类相中的W在冷却过程中析出分离，在达到晶界的凝固温度500～700℃左右时，由于处于B、C、O扩散速度较慢的区域，不容易形成大颗粒的W2B、WC、WO化合物，W以微小并且均一分散的方式实现析出。在将原料合金粉碎之后，进入成形烧结工序，主相晶粒会在烧结过程中长大，但是，由于结晶晶界中存在的W钉扎(Pinning effect)晶界的迁移，可以有效防止晶粒异常长大(AGG)的产生，对于SQ、Hcj性能的提升起到非常好效果。钉扎(Pinning effect)晶界迁移的原理用图1举例说明，图1中的黑点代表W钉扎结晶物，2代表合金溶液，3代表晶粒，箭头表示晶粒生长方向，由图1中所见，W钉扎结晶物在晶粒生长过程中，积聚在晶粒生长方向的表面上，隔断了晶粒与外部的物质迁移过程，从而阻碍晶粒长大。
同样地，稀土类金属间化合物R2Fe14B也由于W微小且均一的析出，防止AGG的发生，使所制得磁铁的方形度(SQ)提升。再者，由于分布在晶界中的 Cu增加了低熔点液相，低熔点液相的增加促进了W的迁移，从图3的EMPA结果可以看到，本发明中，W在晶界中分布相当均匀，且分布范围超过富Nd相的分布范围，完全包覆了整个富Nd相，可以认为是W发挥钉扎效果、阻碍晶粒长大的证据。
再者，以往的方式中，由于添加量较多的Zr、Hf、Mo、V、W、Nb等高熔点金属，会出现高熔点金属的硼化物相，这些硼化物相的硬度非常高，非常硬，会使加工性能急剧劣化。而本发明中的W由于含有量非常微量，几乎不会出现高熔点金属硼化物相、就算出现但是也只会是微乎其量的存在，所以几乎没有加工劣化。
需要说明的是，在目前较多采用的稀土制备方法中，有采用石墨坩埚电解槽，圆桶形石墨坩埚作阳极，坩埚轴线上配置钨(W)棒做阴极，且底部用钨坩埚收集稀土金属的方式。在上述制备稀土元素(如Nd)的过程中，不可避免有少量W混入其中。当然，也可以使用钼(Mo)等其他高熔点金属做阴极，同时使用钼坩埚收集稀土金属的方式，获得完全不含W的稀土元素。
在本发明中，W也可以是原料(如纯铁、稀土金属、B等)等的杂质，其根据原料中杂质的含量选定本发明所使用的原料；当然，也可以选择W含量在现有设备的检测限以下(可视为不含有W)的原料(如纯铁、稀土金属、B等)，采用加入本发明所描述含量的W金属原料的方式。简而言之，只要原料中含有必要量的W即可，不管W的来源为何。表1中举例显示了不同产地不同工场的金属Nd的W元素含量。
表1 不同产地不同工场的金属Nd的W元素含量
Figure PCTCN2015075512-appb-000001
Figure PCTCN2015075512-appb-000002
表1中的2N5所代表的含义为99.5％。
需要说明的是，本发明中，R：12at％～15.2at％、B：5at％～8at％、余量为0at％～20at％的Co和Fe等的含量范围为本行业的常规选择，因此，在实施例中，没有对R、B、Fe和Co的含量范围加以试验和验证。
另外，本发明需要在低氧环境中完成磁铁的全部制造工序，使O含量控制在0.1at％～1.0at％，才能获得本发明所声称的效果，一般而言，具有较高氧含量(2500ppm以上)的稀土磁铁可以减少AGG的产生，而较低氧含量(2500ppm以下)的稀土磁铁虽然具有很好的磁性能，却容易产生AGG，而本发明仅含有极微量的W和少量的Cu，在低氧含量磁铁中也同样实现了减少AGG的效果。
需要说明的是，由于磁铁的低氧制造工序已是现有技术，且本发明的所有实施例全部采用低氧制造方式，在此不再予以详细描述。
在推荐的实施方式中，X的含量在2.0at％以下。
在推荐的实施方式中，所述的烧结磁铁由如下的工序制得：将烧结磁铁原料成分熔融液以102℃/秒～104℃/秒的冷却速度制备成烧结磁铁用合金的工序；将烧结磁铁用合金粗粉碎后再通过微粉碎制成细粉的工序；用磁场成形法获得成形体，并在真空或惰性气体中以900℃～1100℃的温度对所述成形体进行烧结后获得。烧结温度采用900℃～1100℃的温度为本行业的常规选择，因此，在实施例中，没有对烧结温度的范围加以试验和验证。
通过上述的方式，在晶界中提高W的分散度，使方形度超过95％，提高磁铁的耐温性能。
我们经过研究发现，提高W分散度的方法有如下几种：
1)调节烧结磁铁成分熔融液制成烧结磁铁用合金的冷却速度，冷却速度越高，W的分散度越好；
2)控制烧结磁铁成分熔融液的粘度，粘度越小，W的分散度越好；
3)调节烧结后的冷却速度，冷却速度越快，造成的晶格缺陷减少，W的分散度越高。
本发明中，主要通过控制熔融液冷却速度来获得提高W的分散度。
在推荐的实施方式中，所述烧结磁铁的B含量优选为5.0at％～6.5at％。由于过多量的B容易与W反应，形成的硼化物相，这些硼化物相的硬度非常高，非常硬，会使加工性急剧劣化，同时，由于形成了大颗粒的硼化物相(WB2相)，W在结晶晶界中均一钉扎(Pinning effect)晶界迁移的效果也受到影响，因此，适当降低B含量，可降低硼化物相的形成，充分发挥W的均一钉扎(Pinning effect)效果。通过FE‐EPMA分析，在B量大于6.5at％之时，会在结晶晶界中产生较多含B的R(T,B)2型相，而在B含量为5.0at％～6.5at％之时，产生了含W的R6T13X(X＝Al、Cu、Ga等)型相，这个相的产生使矫顽力和方形度变优，并具有弱磁性，W有利于R6T13X型相产生，并提高其稳定性。
在推荐的实施方式中，所述烧结磁铁的Al含量优选为0.8at％～2.0at％，根据FE‐EPMA的分析，当Al为0.8～2.0at％时，会生成含W的R6T13X(X＝Al、Cu、Ga等)型相，这个相的产生使矫顽力和方形度变优，并具有弱磁性，Al有利于R6T13X型相产生，并提高其稳定性。
需要说明的是，本发明中提及的不可避免的杂质还包括在原料中或者在制造过程中不可避免混入的少量C、N、S、P及其他杂质，因此，本发明中提及的所述烧结磁铁在制作过程中，最好将C含量控制在1at％以下，更优选在0.4at％以下，N含量则控制在0.5at％以下，S含量则控制在0.1at％以下，P含量则控制在0.1at％以下。
在推荐的实施方式中，所述粗粉碎为烧结磁铁用合金吸氢破碎、得到粗粉的工序，所述微粉碎为粗粉气流粉碎的工序，还包括从微粉碎后的粉末中除去粒径1.0μm以下的至少一部分，由此使粒径1.0μm以下的粉末体积减少至全体粉 末体积的10％以下的工序。
在推荐的实施方式中，还包括将所述烧结磁体进行RH(重稀土元素)晶界扩散处理的工序。晶界扩散一般在700℃～1050℃的温度下进行，这一温度范围为本行业的常规选择，因此，在实施例中，没有对上述温度范围加以试验和验证。
在对上述烧结磁铁实施晶界扩散时，微量的W可以在结晶晶界中产生非常微小的W结晶，不会阻碍RH的扩散，所以，扩散速度非常快。另外，由于含有适量的Cu，所以形成了低熔点的富Nd相、可发挥进一步促进扩散的效果。所以，本发明的磁铁通过RH的晶界扩散，可以获得非常高的性能，得到飞跃性的提高效果。
在推荐的实施方式中，所述的RH选自Dy或Tb中的至少一种。
在推荐的实施方式中，还包括时效处理的步骤：对所述烧结磁铁在400℃～650℃的温度进行时效处理。
在推荐的实施方式中，还包括二段时效处理的步骤：对所述烧结磁铁在800℃～950℃的温度下进行1小时～2小时的一级热处理之后，再对所述烧结磁铁在450℃～660℃的温度下进行1小时～4小时的二级热处理。
在推荐的实施方式中，所述烧结磁铁的O含量为0.1at％～0.5at％。在上述的区间内，O、W、Cu配比达到了最佳配比，烧结磁铁的耐热性能高，磁铁在动态工作条件下的稳定性高，不存在AGG的时候，氧含量低，Hcj升高。
在推荐的实施方式中，所述烧结磁铁的Ga含量为0.05at％～0.8at％。
本发明的另一目的在于提供一种含W的R‐Fe‐B‐Cu系烧结磁铁用急冷合金。
一种含W的R‐Fe‐B‐Cu系烧结磁铁用急冷合金，其特征在于：所述急冷合金的结晶晶界中具有W含量在0.004at％以上、0.26at％以下的富W区域，所述富W区域在所述结晶晶界中呈均一分散的分布，并占所述结晶晶界的至少50体积％。
与现有技术相比，本发明具有如下的特点：
1)本发明基于背景技术中3种量产磁铁技术提高所制磁铁的性能，悉心进行了与微量元素相关的研究，通过抑制烧结时的AGG来提高磁铁的SQ、Hcj、Br、(BH)max，结果表明，微量的W钉扎结晶物在结晶晶界中均一钉扎(Pinning effect)晶界的迁移，可以有效防止晶粒异常长大(AGG)的产生，并可以获得显著的改善效果。
2)本发明中所含有的W的含量非常微量、并均一分散，几乎不会出现高熔点大颗粒金属硼化物相、就算出现但是也只会微乎其量的存在，所以几乎没有加工劣化。
3)本发明含有微量的W(非磁性元素)，即0.03at％以下的含有量，稀释效果较少，另外，急冷后的磁体合金中完全不含有非晶质相和各向同性急冷相，经FE‐EPMA检测，微量含有的W大部分存在结晶晶界中，所以本发明中W的微量含有完全不会降低Br、(BH)max，反而还会使Br、(BH)max提高。
4)在本发明的组份中含有微量Cu、W，使得晶界中的高熔点【如WB2相(熔点2365℃)等】金属间化合物相无法生成，而产生较多如RCu(熔点662℃)、RCu2(熔点840℃)、Nd‐Cu共晶合金(熔点492℃)等的低熔点相，作为结果，在晶界扩散温度下结晶晶界中除W相外几乎全部溶解，晶界扩散的效率极佳，方形度和矫顽力以前所未有的程度增加，方形度更是达到99％以上，从而获得了耐热性能良好的高性能磁铁。这里的WB2相包括WFeB合金、WFe合金、WB合金等。
5)微量的W可以促进R6T13X(X＝Al、Cu、Ga等)型相的形成，这个相的产生使矫顽力和方形度变优，并具有弱磁性。
附图说明
图1为W钉扎(Pinning effect)晶界迁移的原理示意图。
图2为实施例一的实施例3的急冷合金片的EPMA检测结果。
图3为实施例一的实施例3的烧结磁体的EPMA检测结果。
具体实施方式
以下结合实施例对本发明作进一步详细说明。
各实施例中提及的BHH、磁性能评价过程、AGG测定的定义如下：
BHH为(BH)max和Hcj的总和，是磁铁综合性能的评价标准之一。
磁性能评价过程：烧结磁铁使用中国计量院的NIM‐10000H型BH大块稀土永磁无损测量系统进行磁性能检测。
AGG测定：将烧结磁铁沿垂直于取向方向的方向抛光，每1cm2所包括的平均AGG数量，本发明中提及的AGG为粒径超过40μm的晶粒。
各实施例中提及的FE‐EPMA检测的检测限为100ppm左右，检测条件如下：
Figure PCTCN2015075512-appb-000003
FE‐EPMA设备最高分辨率达到3nm，在上述检测条件下，分辨率也可达到50nm。
实施例一
原料配制过程：准备纯度99.5％的Nd、Dy、工业用Fe‐B、工业用纯Fe、纯度99.9％的Co和纯度99.5％的Cu、Al，纯度为99.999％的W，以原子百分比at％配制。
为准确控制W的使用配比，该实施例中，所选用的Nd、Dy、Fe、B、Al、Cu和Co中的W含量在现有设备的检测限以下，W的来源为额外添加的W金属。
各元素的含量如表2所示：
表2 各元素的配比(at％)
Figure PCTCN2015075512-appb-000004
Figure PCTCN2015075512-appb-000005
各序号组按照表2中元素组成进行配制，分别称量、配制了100Kg的原料。
熔炼过程：每次取1份配制好的原料放入氧化铝制的坩埚中，在高频真空感应熔炼炉中在10‐2Pa的真空中以1500℃以下的温度进行真空熔炼。
铸造过程：在真空熔炼后的熔炼炉中通入Ar气体使气压达到5万Pa后，使用单辊急冷法进行铸造，以102℃/秒～104℃/秒的冷却速度获得急冷合金，将急冷合金在600℃进行60分钟的保温热处理，然后冷却到室温。
对实施例3制成的急冷合金的Cu、Nd和W等成分进行FE‐EPMA(场发射电子探针显微分析)【日本电子株式会社(JEOL)，8530F】检测，结果如图2中所示，可以观察到，W以较高的分散度分布在富R相中。
对实施例2、3、4、5和6的急冷合金片进行FE‐EPMA检测，富W区域在结晶晶界中呈均一分散的分布，并占合金结晶晶界的至少50体积％，其中，富W区域是W含量为0.004at％以上、0.26at％以下的区域。
氢破粉碎过程：在室温下将放置急冷合金的氢破用炉抽真空，而后向氢破用炉内通入纯度为99.5％的氢气至压力0.1MPa，放置2小时后，边抽真空边升温，在500℃的温度下抽真空，之后进行冷却，取出氢破粉碎后的粉末。
微粉碎工序：在氧化气体含量100ppm以下的氮气气氛下，在粉碎室压力为0.4MPa的压力下对氢破粉碎后的粉末进行气流磨粉碎，得到细粉，细粉的平均粒度为4.5μm。氧化气体指的是氧或水分。
使用分级器对部分微粉碎后的细粉(占细粉总重量30％)分级，除去粒径1.0μm以下的粉粒，再将分级后的细粉与剩余未分级的细粉混合。混合后的细粉中，粒径1.0μm以下的粉末体积减少至全体粉末体积的10％以下。
在气流磨粉碎后的粉末中添加辛酸甲酯，辛酸甲酯的添加量为混合后粉末重量的0.2％，再用V型混料机充分混合。
磁场成形过程：使用直角取向型的磁场成型机，在1.8T的取向磁场中，在0.4ton/cm2的成型压力下，将上述添加了辛酸甲酯的粉末一次成形成边长为25mm的立方体，一次成形后在0.2T的磁场中退磁。
为使一次成形后的成形体不接触到空气，将其进行密封，再使用二次成形机(等静压成形机)在1.4ton/cm2的压力下进行二次成形。
烧结过程：将各成形体搬至烧结炉进行烧结，烧结在10‐3Pa的真空下，在200℃和800℃的温度下各保持2小时后，以1030℃的温度烧结2小时，之后通入Ar气体使气压达到0.1MPa后，冷却至室温。
热处理过程：烧结体在高纯度Ar气中，以460℃温度进行1小时热处理后，冷却至室温后取出。
加工过程：经过热处理的烧结体加工成
Figure PCTCN2015075512-appb-000006
厚度5mm的磁铁，5mm方向为磁场取向方向。
实施例1～7烧结体制成的磁铁直接进行磁性能检测，评定其磁特性。实施例磁铁的评价结果如表3、表4中所示：
表3 实施例的显微结构评价情况
Figure PCTCN2015075512-appb-000007
表3中的非晶质相和各向同性相考察的是急冷合金中的非晶质相和各向同 性相。
表3中的富W相是0.004at％以上、0.26at％以下的区域。
表4 实施例的磁性能评价情况
Figure PCTCN2015075512-appb-000008
在整个实施过程中，特别注意控制O、C和N的含量，将上述O、C和N三种元素分别磁铁中的含量控制在0.1～0.5at％、0.3at％以下和0.1at％以下。
作为结论我们可以得出：本发明中，在磁铁中W含量小于0.0005at％之时，由于W含量过少，难以发挥钉扎效果，而原料中Cu的存在，容易引起AGG，导致SQ和Hcj降低，相对地，在W含量超过0.03at％之时，会产生一部分的WB2相，使方形度、磁性能降低，另外，其所获得的急冷合金中会产生非晶质相及各向同性急冷相，使磁铁性能急剧降低。
对实施例3制成烧结磁铁的Cu、Nd和W等成分进行FE‐EPMA(场发射电子探针显微分析)【日本电子株式会社(JEOL)，8530F】检测，结果如图3中所示，可以观察到，W以较高的分散度均一钉扎(Pinning effect)晶界的迁移，防止AGG的形成。
同样地，对实施例2、4、5和6进行FE‐EPMA检测，同样可以观察到，W以较高的分散度均一钉扎(Pinning effect)晶界的迁移，防止AGG的形成。
实施例二
在原料配制过程：准备纯度99.9％的Nd、Pr、Tb、纯度99.9％的B、纯度99.9％ 的Fe、纯度99.999％的W和纯度99.5％的Cu、Al，以原子百分比at％配制。
为准确控制W的使用配比，该实施例中，所选用的Nd、Pr、Tb、Fe、B、Al和Cu中的W含量在现有设备的检测限以下，W的来源为额外添加的W金属。
各元素的含量如表5所示：
表5 各元素的配比(at％)
Figure PCTCN2015075512-appb-000009
各序号组按照表5中元素组成进行配制，分别称量、配制了100Kg的原料。
熔炼过程：每次取1份配制好的原料放入氧化铝制的坩埚中，在高频真空感应熔炼炉中在10‐2Pa的真空中以1500℃以下的温度进行真空熔炼。
铸造过程：在真空熔炼后的熔炼炉中通入Ar气体使气压达到3万Pa后，使用单辊急冷法进行铸造，以102℃/秒～104℃/秒的冷却速度获得急冷合金，将急冷合金在600℃进行60分钟的保温热处理，然后冷却到室温。
对实施例2～7的急冷合金片进行FE‐EPMA检测，富W区域在结晶晶界中呈均一分散的分布，并占合金结晶晶界的至少50体积％，其中，富W区域是W含量为0.004at％以上、0.26at％以下的区域。
氢破粉碎过程：在室温下将放置急冷合金的氢破用炉抽真空，而后向氢破用炉内通入纯度为99.5％的氢气至压力0.1MPa，放置125分钟后，边抽真空边升温，在500℃的温度下抽真空2小时，之后进行冷却，取出氢破粉碎后的粉末。
在微粉碎工序：在氧化气体含量100ppm以下的气氛下，在粉碎室压力为 0.41MPa的压力下对氢破粉碎后的粉末进行气流磨粉碎，得到细粉，细粉的平均粒度为4.30μm。氧化气体指的是氧或水分。
在气流磨粉碎后的粉末中添加辛酸甲酯，辛酸甲酯的添加量为混合后粉末重量的0.25％，再用V型混料机充分混合。
磁场成形过程：使用直角取向型的磁场成型机，在1.8T的取向磁场中，在0.3ton/cm2的成型压力下，将上述添加了辛酸甲酯的粉末一次成形成边长为25mm的立方体，一次成形后在0.2T的磁场中退磁。
为使一次成形后的成形体不接触到空气，将其进行密封，再使用二次成形机(等静压成形机)在1.0ton/cm2的压力下进行二次成形。
烧结过程：将各成形体搬至烧结炉进行烧结，烧结在10‐3Pa的真空下，在200℃和800℃的温度下各保持3小时后，以1020℃的温度烧结2小时，之后通入Ar气体使气压达到0.1MPa后，冷却至室温。
热处理过程：烧结体在高纯度Ar气中，以620℃温度进行1小时热处理后，冷却至室温后取出。
加工过程：经过热处理的烧结体加工成
Figure PCTCN2015075512-appb-000010
厚度5mm的磁铁，5mm方向为磁场取向方向。
实施例1～8的烧结体制成的磁铁直接进行磁性能检测，评定其磁特性。实施例磁铁的评价结果如表6和表7中所示：
表6 实施例的显微结构评价情况
Figure PCTCN2015075512-appb-000011
Figure PCTCN2015075512-appb-000012
表6中的非晶质相和各向同性相考察的是急冷合金中的非晶质相和各向同性相。
表6中的富W相是0.004at％以上、0.26at％以下的区域。
表7 实施例的磁性能评价情况
Figure PCTCN2015075512-appb-000013
在整个实施过程中，特别注意控制O、C和N的含量，将上述O、C和N三种元素分别磁铁中的含量控制在0.1～0.5at％、0.4at％以下和0.2at％以下。
作为结论我们可以得出：Cu含量小于0.05at％之时，矫顽力的热处理温度依存性会变大，磁铁性能降低，相对地，在Cu含量大于1.2at％之时，由于Cu的低熔点现象，AGG的产生量也会增加，即使W的钉扎(Pinning effect)效应也难以阻止AGG大量形成，由此可知，在低氧含量磁铁中，存在合适的Cu、W范围。
同样地，对实施例2～7进行FE‐EPMA检测【日本电子株式会社(JEOL)，8530F】，作为检测结果，可以观察到，W以较高的分散度均一钉扎(pinning effect)晶界的迁移，防止AGG的形成。
实施例三
在原料配制过程：准备纯度99.5％的Nd、工业用Fe‐B、工业用纯Fe、纯度 99.9％的Co、纯度99.5％的Cu和纯度99.999％的W，以原子百分比at％配制。
为准确控制W的使用配比，该实施例中，所选用的Nd、Fe、B、Cu和Co原料中的W含量在现有设备的检测限以下，W的来源为额外添加的W金属。
各元素的含量如表8所示：
表8 各元素的配比(at％)
Figure PCTCN2015075512-appb-000014
按照表8中元素组成进行配制，称量、配制了700Kg的原料。
熔炼过程：取配制好的原料放入氧化铝制的坩埚中，在高频真空感应熔炼炉中在10‐2Pa的真空中以1500℃以下的温度进行真空熔炼。
铸造过程：在真空熔炼后的熔炼炉中通入Ar气体使气压达到5万Pa后，使用单辊急冷法进行铸造，以102℃/秒～104℃/秒的冷却速度获得急冷合金，将急冷合金在600℃进行60分钟的保温热处理，然后冷却到室温。
对急冷合金片进行FE‐EPMA检测，富W区域在结晶晶界中呈均一分散的分布，并占合金结晶晶界的至少50体积％，其中，富W区域是W含量为0.004at％以上、0.26at％以下的区域。
氢破粉碎过程：在室温下将放置急冷合金的氢破用炉抽真空，而后向氢破用炉内通入纯度为99.5％的氢气至压力0.1MPa，放置97分钟后，边抽真空边升温，在500℃的温度下抽真空2小时，之后进行冷却，取出氢破粉碎后的粉末。
微粉碎工序：将氢破粉碎后的粉末分成7份，将每份粉末分别在氧化气体含量10～3000ppm以下的气氛下，在粉碎室压力为0.42MPa的压力下对各份氢破粉碎后的粉末进行气流磨粉碎，得到细粉，细粉的平均粒度为4.51μm。氧化气体指的是氧或水分。
在各份气流磨粉碎后的粉末中分别添加辛酸甲酯，辛酸甲酯的添加量为混合后粉末重量的0.1％，再用V型混料机充分混合。
磁场成形过程：分别使用直角取向型的磁场成型机，在1.8T的取向磁场中， 在0.2ton/cm2的成型压力下，将上述添加了辛酸甲酯的粉末一次成形成边长为25mm的立方体，一次成形后在0.2T的磁场中退磁。
为使一次成形后的成形体不接触到空气，将其进行密封，再使用二次成形机(等静压成形机)在1.4ton/cm2的压力下进行二次成形。
烧结过程：将各成形体搬至烧结炉进行烧结，烧结在10‐3Pa的真空下，在200℃和700℃的温度下各保持2小时后，以1020℃的温度烧结2小时，之后通入Ar气体使气压达到0.1MPa后，冷却至室温。
热处理过程：烧结体在高纯度Ar气中，以900℃温度进行1小时的一级热处理后，再在500℃的温度下进行1小时的二级热处理，冷却至室温后取出。
加工过程：经过热处理的烧结体加工成
Figure PCTCN2015075512-appb-000015
厚度5mm的磁铁，5mm方向为磁场取向方向。
热减磁的测定：烧结磁铁置于150℃环境中保温30min，然后再自然冷却降温到室温，测量磁通，测量的结果和加热前的测量数据比较，计算加热前和加热后的磁通衰减率。
实施例1～7的烧结体制成的磁铁直接进行磁性能检测，评定其磁特性。实施例磁铁的评价结果如表9和表10中所示：
表9 实施例的显微结构评价情况
Figure PCTCN2015075512-appb-000016
表9中的富W相是0.004at％以上、0.26at％以下的区域。
表10 实施例的磁性能评价情况
Figure PCTCN2015075512-appb-000017
在整个实施过程中，特别注意控制C和N的含量，将上述C和N三种元素分别磁铁中的含量控制在0.2at％以下和0.25at％以下。
作为结论我们可以得出：即使有适量的W、Cu的存在，在磁铁的O含量小于0.1at％之时，其超过了W钉扎效果的界限，处于非常容易发生AGG的状态，所以仍旧会发生AGG现象，使磁铁性能急剧降低。相对地，即使有适量的W、Cu的存在，在磁铁的O含量超过1.0at％之时，氧含量的分散性开始变差，产生了氧含量较多与氧含量较少的地方，由于不均匀所以增加了AGG的产生，使矫顽力和方形度降低。
同样地，对实施例2～6进行FE‐EPMA检测【日本电子株式会社(JEOL)，8530F】，作为检测结果，同样可以观察到，W以较高的分散度均一钉扎(Pinning effect)晶界的迁移，防止AGG的形成。
实施例四
原料配制过程：准备纯度99.5％的Nd、Dy、工业用Fe‐B、工业用纯Fe、纯度99.9％的Co和纯度99.5％的Cu、Al，纯度为99.999％的W，以原子百分比at％配制。
为准确控制W的使用配比，该实施例中，所选用的Nd、Dy、B、Al、Cu、Co和Fe中的W含量在现有设备的检测限以下，W的来源为额外添加的W金属。 其含量如表11所示：
表11 各元素的配比(at％)
Figure PCTCN2015075512-appb-000018
各序号组按照表11中元素组成进行配制，分别称量、配制了100Kg的原料。
熔炼过程：每次取1份配制好的原料放入氧化铝制的坩埚中，在高频真空感应熔炼炉中在10‐2Pa的真空中以1550℃以下的温度进行真空熔炼。
铸造过程：在真空熔炼后的熔炼炉中通入Ar气体使气压达到2万Pa后，使用单辊急冷法进行铸造，以102℃/秒～104℃/秒的冷却速度获得急冷合金，将急冷合金在800℃进行10分钟的保温热处理，然后冷却到室温。
对实施例1～7的急冷合金片进行FE‐EPMA检测，富W区域在结晶晶界中呈均一分散的分布，并占合金结晶晶界的至少50体积％，其中，富W区域是W含量为0.004at％以上、0.26at％以下的区域。
氢破粉碎过程：在室温下将放置急冷合金的氢破用炉抽真空，而后向氢破用炉内通入纯度为99.5％的氢气至压力0.1MPa，放置120分钟后，边抽真空边升温，在500℃的温度下抽真空2小时，之后进行冷却，取出氢破粉碎后的粉末。
微粉碎工序：在氧化气体含量100ppm以下的气氛下，在粉碎室压力为0.6MPa的压力下对氢破粉碎后的粉末进行气流磨粉碎，得到细粉，细粉的平均粒度为4.5μm。氧化气体指的是氧或水分。
使用分级器对部分微粉碎后的细粉(占细粉总重量30％)进行分级，除去粒径1.0μm以下的粉粒，再将分级后的细粉与剩余未分级的细粉混合。混合后 的细粉中，粒径1.0μm以下的粉末体积减少至全体粉末体积的2％以下。
在气流磨粉碎后的粉末中添加辛酸甲酯，辛酸甲酯的添加量为混合后粉末重量的0.2％，再用V型混料机充分混合。
磁场成形过程：使用直角取向型的磁场成型机，在1.8T的取向磁场中，在0.2ton/cm2的成型压力下，将上述添加了辛酸甲酯的粉末一次成形成边长为25mm的立方体，一次成形后在0.2T的磁场中退磁。
为使一次成形后的成形体不接触到空气，将其进行密封，再使用二次成形机(等静压成形机)在1.0ton/cm2的压力下进行二次成形。
烧结过程：将各成形体搬至烧结炉进行烧结，烧结在10‐3Pa的真空下，在200℃和800℃的温度下各保持2小时后，以1040℃的温度烧结2小时，之后通入Ar气体使气压达到0.1MPa后，冷却至室温。
热处理过程：烧结体在高纯度Ar气中，以400℃温度进行1小时热处理后，冷却至室温后取出。
加工过程：经过热处理的烧结体加工成
Figure PCTCN2015075512-appb-000019
厚度5mm的磁铁，5mm方向为磁场取向方向。
实施例1～7烧结体制成的磁铁直接进行磁性能检测，评定其磁特性。实施例磁铁的评价结果如表12和表13中所示：
表12 实施例的显微结构评价情况
Figure PCTCN2015075512-appb-000020
表12中的非晶质相和各向同性相考察的是急冷合金中的非晶质相和各向同性相。
表12中的富W相是0.004at％以上、0.26at％以下的区域。
表13 实施例的磁性能评价情况
Figure PCTCN2015075512-appb-000021
在整个实施过程中，特别注意控制O、C和N的含量，将上述O、C和N三种元素分别磁铁中的含量控制在0.1～0.5at％、0.3at％以下和0.1at％以下。
对实施例1～7进行FE‐EPMA检测【日本电子株式会社(JEOL)，8530F】，作为检测结果，可以观察到，W以较高的分散度均一钉扎(pinning effect)晶界的迁移，防止AGG的形成。
结论：通过FE‐EPMA分析，在B量大于6.5at％之时，较多含B的R(T,B)2型相会在结晶晶界中产生，而在B含量为5at％～6.5at％之时，产生了含W的R6T13X(X＝Al、Cu等)型相，这个相的产生使矫顽力和方形度变优，并具有弱磁性，W有利于R6T13X型相产生，并提高其稳定性。
实施例五
原料配制过程：准备纯度99.5％的Nd、Dy、工业用Fe‐B、工业用纯Fe、纯度99.9％的Co和纯度99.5％的Cu、Al，纯度为99.999％的W，以原子百分比at％配制。
为准确控制W的使用配比，该实施例中，所选用的Nd、Dy、B、Al、Cu、Co和Fe中的W含量在现有设备的检测限以下，W的来源为额外添加的W金属。 各元素的含量如表14所示：
表14 各元素的配比(at％)
Figure PCTCN2015075512-appb-000022
各序号组按照表14中元素组成进行配制，分别称量、配制了100Kg的原料。
熔炼过程：每次取1份配制好的原料放入氧化铝制的坩埚中，在高频真空感应熔炼炉中在10‐2Pa的真空中以1500℃以下的温度进行真空熔炼。
铸造过程：在真空熔炼后的熔炼炉中通入Ar气体使气压达到5万Pa后，使用单辊急冷法进行铸造，以102℃/秒～104℃/秒的冷却速度获得急冷合金，将急冷合金在700℃进行5分钟的保温热处理，然后冷却到室温。
氢破粉碎过程：在室温下将放置急冷合金的氢破用炉抽真空，而后向氢破用炉内通入纯度为99.5％的氢气至压力0.1MPa，放置120分钟后，边抽真空边升温，在600℃的温度下抽真空2小时，之后进行冷却，取出氢破粉碎后的粉末。
微粉碎工序：在氧化气体含量100ppm以下的气氛下，在粉碎室压力为0.5MPa的压力下对氢破粉碎后的粉末进行气流磨粉碎，得到细粉，细粉的平均粒度为5.0μm。氧化气体指的是氧或水分。
对部分微粉碎后的细粉(占细粉总重量30％)过筛，除去粒径1.0μm以下的粉粒，再将过筛后的细粉与剩余未过筛的细粉混合。混合后的细粉中，粒径1.0μm以下的粉末体积减少至全体粉末体积的10％以下。
在气流磨粉碎后的粉末中添加辛酸甲酯，辛酸甲酯的添加量为混合后粉末重量的0.2％，再用V型混料机充分混合。
磁场成形过程：使用直角取向型的磁场成型机，在1.8T的取向磁场中，在0.2ton/cm2的成型压力下，将上述添加了辛酸甲酯的粉末一次成形成边长为25mm的立方体，一次成形后在0.2T的磁场中退磁。
为使一次成形后的成形体不接触到空气，将其进行密封，再使用二次成形机(等静压成形机)在1.0ton/cm2的压力下进行二次成形。
烧结过程：将各成形体搬至烧结炉进行烧结，烧结在10‐3Pa的真空下，在200℃和800℃的温度下各保持2小时后，以1060℃的温度烧结2小时，之后通入Ar气体使气压达到0.1MPa后，冷却至室温。
热处理过程：烧结体在高纯度Ar气中，以420℃温度进行1小时热处理后，冷却至室温后取出。
加工过程：经过热处理的烧结体加工成
Figure PCTCN2015075512-appb-000023
厚度5mm的磁铁，5mm方向为磁场取向方向。
实施例1～7烧结体制成的磁铁直接进行磁性能检测，评定其磁特性。实施例磁铁的评价结果如表15中所示：
表15 实施例的显微结构评价情况
Figure PCTCN2015075512-appb-000024
表15中的非晶质相和各向同性相考察的是急冷合金中的非晶质相和各向同性相。
表15中的富W相是0.004at％以上、0.26at％以下的区域。
表16 实施例的磁性能评价情况
Figure PCTCN2015075512-appb-000025
在整个实施过程中，特别注意控制O、C和N的含量，将上述O、C和N三种元素分别磁铁中的含量控制在0.1～0.5at％、0.3at％以下和0.1at％以下。
对实施例1～7进行FE‐EPMA检测【日本电子株式会社(JEOL)，8530F】，作为检测结果，可以观察到，W以较高的分散度均一钉扎(pinning effect)晶界的迁移，防止AGG的形成。
结论：根据FE‐EPMA的分析，当Al为0.8～2.0％时，会生成含W的R6T13X(X＝Al、Cu等)型相，这个相的产生使矫顽力变优，并具有弱磁性，W有利于R6T13X型相产生，并提高其稳定性。
实施例六
将实施例一制得的各组烧结体分别加工成
Figure PCTCN2015075512-appb-000026
厚度5mm的磁铁，5mm方向为磁场取向方向。
晶界扩散处理过程：将各组烧结体加工成的磁铁洗净，表面洁净化后，分别使用Dy氧化物和Tb氟化物按3:1比例配制成的原料，全面喷雾涂覆在磁铁上，将涂覆后的磁铁干燥，在高纯度Ar气体气氛中，以850℃的温度扩散热处理24小时。
磁性能评价过程：烧结磁铁使用中国计量院的NIM‐10000H型BH大块稀土永磁无损测量系统进行磁性能检测。结果如表17中所示：
表17 实施例的矫顽力评价情况
Figure PCTCN2015075512-appb-000027
从表17中可以看到，本发明中的微量W在结晶晶界中产生非常微小的W結晶，不会阻碍Dy，Tb的扩散，所以，扩散速度非常快。另外，由于含有适量的Cu，所以形成了低熔点的富Nd相、可发挥进一步促进扩散的效果。所以，本发明的磁铁通过Dy、Tb的晶界扩散，可以获得非常高的性能。
实施例七
在原料配制过程：准备纯度99.9％的Nd、Dy、Tb、纯度99.9％的B、纯度99.9％的Fe、和纯度99.5％的Cu、Co、Nb、Al、Ga，以原子百分比at％配制。
为准确控制W的使用配比，该实施例中，所选用的Dy、Tb、Fe、B、Cu、Co、Nb、Al和Ga中的W含量在现有设备的检测限以下，所选用的Nd中则含有W，W元素的含量为0.01at％。
各元素的含量如表18所示：
表18 各元素的配比(at％)
Figure PCTCN2015075512-appb-000028
Figure PCTCN2015075512-appb-000029
各序号组按照表18中元素组成进行配制，分别称量、配制了100Kg的原料。
熔炼过程：每次取1份配制好的原料放入氧化铝制的坩埚中，在高频真空感应熔炼炉中在10‐2Pa的真空中以1500℃以下的温度进行真空熔炼。
铸造过程：在真空熔炼后的熔炼炉中通入Ar气体使气压达到3.5万Pa后，使用单辊急冷法进行铸造，以102℃/秒～104℃/秒的冷却速度获得急冷合金，将急冷合金在550℃进行10分钟的保温热处理，然后冷却到室温。
氢破粉碎过程：在室温下将放置急冷合金的氢破用炉抽真空，而后向氢破用炉内通入纯度为99.5％的氢气至压力0.085MPa，放置160分钟后，边抽真空边升温，在520℃的温度下抽真空，之后进行冷却，取出氢破粉碎后的粉末。
微粉碎工序：在氧化气体含量10ppm以下的气氛下，在粉碎室压力为0.42MPa的压力下对氢破粉碎后的粉末进行气流磨粉碎，得到细粉，细粉的平均粒度为4.28μm。氧化气体指的是氧或水分。
在气流磨粉碎后的粉末中添加辛酸甲酯，辛酸甲酯的添加量为混合后粉末重量的0.25％，再用V型混料机充分混合。
磁场成形过程：使用直角取向型的磁场成型机，在1.8T的取向磁场中，在0.3ton/cm2的成型压力下，将上述添加了辛酸甲酯的粉末一次成形成边长为25mm的立方体，一次成形后在0.2T的磁场中退磁。
为使一次成形后的成形体不接触到空气，将其进行密封，再使用二次成形机(等静压成形机)在1.0ton/cm2的压力下进行二次成形。
烧结过程：将各成形体搬至烧结炉进行烧结，烧结在10‐3Pa的真空下，在300℃和800℃的温度下各保持3小时后，以1030℃的温度烧结2小时，之后通入Ar气体使气压达到0.1MPa后，冷却至室温。
热处理过程：烧结体在高纯度Ar气中，以600℃温度进行2小时热处理后，冷却至室温后取出。
加工过程：经过热处理的烧结体加工成
Figure PCTCN2015075512-appb-000030
厚度5mm的磁铁，5mm方向为磁场取向方向。
实施例1～8的烧结体制成的磁铁直接进行磁性能检测，评定其磁特性。实施例磁铁的评价结果如表19和表20中所示：
表19 实施例的显微结构评价情况
Figure PCTCN2015075512-appb-000031
表19中的非晶质相和各向同性相考察的是急冷合金中的非晶质相和各向同性相。
表19中的富W相是0.004at％以上、0.26at％以下的区域。
表20 实施例的磁性能评价情况
Figure PCTCN2015075512-appb-000032
在整个实施过程中，特别注意控制O、C和N的含量，将上述O、C和N三 种元素分别磁铁中的含量控制在0.1～0.5at％、0.4at％以下和0.2at％以下。
作为结论我们可以得出：Ga含量小于0.05at％之时，矫顽力的热处理温度依存性会变大，磁铁性能降低，相对地，在Ga含量大于0.8at％之时，由于Ga为非磁性元素，导致Br、(BH)max降低。
同样地，对实施例1～8进行FE‐EPMA检测【日本电子株式会社(JEOL)，8530F】，作为检测结果，可以观察到，W以较高的分散度均一钉扎(pinning effect)晶界的迁移，防止AGG的形成。
实施例八
在原料配制过程：准备纯度99.9％的Nd、Dy、Gd、Tb、纯度99.9％的B、纯度99.9％的Fe、和纯度99.5％的Cu、Co、Nb、Al、Ga，以原子百分比at％配制。各元素的含量如表5所示。
为准确控制W的使用配比，该实施例中，所选用的Dy、Gd、Tb、Fe、B、Cu、Co、Nb、Al和Ga中的W含量在现有设备的检测限以下，所选用的Nd中则含有W，W元素的含量为0.01at％。
各元素的含量如表21所示：
表21 各元素的配比(at％)
Figure PCTCN2015075512-appb-000033
各序号组按照表21中元素组成进行配制，分别称量、配制了100Kg的原料。
熔炼过程：每次取1份配制好的原料放入氧化铝制的坩埚中，在高频真空感应熔炼炉中在10‐2Pa的真空中以1450℃以下的温度进行真空熔炼。
铸造过程：在真空熔炼后的熔炼炉中通入Ar气体使气压达到4.5万Pa后，使用单辊急冷法进行铸造，以102℃/秒～104℃/秒的冷却速度获得急冷合金，将 急冷合金在800℃进行5分钟的保温热处理，然后冷却到室温。
氢破粉碎过程：在室温下将放置急冷合金的氢破用炉抽真空，而后向氢破用炉内通入纯度为99.5％的氢气至压力0.09MPa，放置150分钟后，边抽真空边升温，在600℃的温度下抽真空，之后进行冷却，取出氢破粉碎后的粉末。
微粉碎工序：在氧化气体含量30ppm以下的气氛下，在粉碎室压力为0.5MPa的压力下对氢破粉碎后的粉末进行气流磨粉碎，得到细粉，细粉的平均粒度为4.1μm。氧化气体指的是氧或水分。
在气流磨粉碎后的粉末中添加硬脂酸铝，硬脂酸铝的添加量为混合后粉末重量的0.05％，再用V型混料机充分混合。
磁场成形过程：使用直角取向型的磁场成型机，在1.8T的取向磁场中，在0.3ton/cm2的成型压力下，将上述添加了硬脂酸铝的粉末一次成形成边长为25mm的立方体，一次成形后在0.2T的磁场中退磁。
为使一次成形后的成形体不接触到空气，将其进行密封，再使用二次成形机(等静压成形机)在1.0ton/cm2的压力下进行二次成形。
烧结过程：将各成形体搬至烧结炉进行烧结，烧结在10‐3Pa的真空下，在200℃和800℃的温度下各保持3小时后，以1050℃的温度烧结2小时，之后通入Ar气体使气压达到0.1MPa后，冷却至室温。
热处理过程：烧结体在高纯度Ar气中，以480℃温度进行2小时热处理后，冷却至室温后取出。
加工过程：经过热处理的烧结体加工成
Figure PCTCN2015075512-appb-000034
厚度5mm的磁铁，5mm方向为磁场取向方向。
实施例1～5的烧结体制成的磁铁直接进行磁性能检测，评定其磁特性。实施例磁铁的评价结果如表22和表23中所示：
表22 实施例的显微结构评价情况
Figure PCTCN2015075512-appb-000035
Figure PCTCN2015075512-appb-000036
表23中的非晶质相和各向同性相考察的是急冷合金中的非晶质相和各向同性相。
表23中的富W相是0.004at％以上、0.26at％以下的区域。
表23 实施例的磁性能评价情况
Figure PCTCN2015075512-appb-000037
在整个实施过程中，特别注意控制O、C和N的含量，将上述O、C和N三种元素分别磁铁中的含量控制在0.1～0.5at％、0.4at％以下和0.2at％以下。
作为结论我们可以得出：在Nb含量大于0.2at％之时，由于Nb的含量升高，在急冷合金片中观察到了非晶质相，而由于非晶质相的存在，导致Br、Hcj降低。
与添加Nb的情形相同，申请人通过试验得知，Zr的含量也应当控制在0.2at％以下。
同样地，对实施例1～5进行FE‐EPMA检测【日本电子株式会社(JEOL)，8530F】，作为检测结果，可以观察到，W以较高的分散度均一钉扎(pinning effect)晶界的迁移，防止AGG的形成。
上述实施例仅用来进一步说明本发明的几种具体的实施方式，但本发明并不局限于实施例，凡是依据本发明的技术实质对以上实施例所作的任何简单修改、 等同变化与修饰，均落入本发明技术方案的保护范围内。
工业实用性
本发明中所含有的W的含量非常微量、并均一分散，可以促进R6T13X(X＝Al、Cu、Ga等)型相的形成，有效防止晶粒异常长大(AGG)的产生，会使Br、(BH)max提高，并可以获得显著的改善效果，同时避免出现高熔点大颗粒金属硼化物相的加工劣化，具有良好的工业实用性。

Claims (12)

  1. 一种含W的R‐Fe‐B‐Cu系烧结磁铁，所述烧结磁铁含有R2Fe14B型主相，所述的R为包括Nd或Pr的至少一种稀土元素，其特征在于，所述稀土磁铁的结晶晶界中具有W含量为0.004at％以上、0.26at％以下的富W区域，所述富W区域在所述结晶晶界中呈均一分散的分布，并占所述烧结磁铁的5.0体积％～11.0体积％。
  2. 根据权利要求1中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于，所述磁铁由包括如下成分的原料制成：
    R：12at％～15.2at％，
    B：5at％～8at％，
    W：0.0005at％～0.03at％，
    Cu：0.05at％～1.2at％，
    X：5.0at％以下、X为选自Al、Si、Ga、Sn、Ge、Ag、Au、Bi、Mn、Nb、Zr或Cr中的至少1种元素，在X包括Nb和/或Zr之时，Nb和Zr的总含量在0.20at％以下，
    余量为0at％～20at％的Co、Fe、以及不可避免的杂质，
    所述杂质包括O，且所述烧结磁铁的O含量为0.1at％～1.0at％。
  3. 根据权利要求2中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于：X的含量在2.0at％以下。
  4. 根据权利要求3中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于，由如下的步骤制得：将烧结磁铁原料成分熔融液以102℃/秒～104℃/秒的冷却速度制备成烧结磁铁用合金的工序；将烧结磁铁用合金粗粉碎后再通过微粉碎制成细粉的工序；用磁场成形法获得成形体，并在真空或惰性气体中以900℃～1100℃的温度对所述成形体进行烧结后获得。
  5. 根据权利要求1或2或3或4中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于：所述烧结磁铁的B含量优选为5at％～6.5at％。
  6. 根据权利要求1或2或3或4中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于：所述烧结磁铁的Al含量优选为0.8at％～2.0at％。
  7. 根据权利要求4所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于：所述粗粉碎为烧结磁铁用合金吸氢破碎、得到粗粉的工序，所述微粉碎为粗粉气流粉碎的工序，还包括从微粉碎后的粉末中除去粒径1.0μm以下的至少一部分，由此使粒径1.0μm以下的粉末体积减少至全体粉末体积的10％以下的工序。
  8. 根据权利要求1或2或3或4中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于：还包括将所述烧结磁体进行RH晶界扩散处理的工序，所述的RH选自Dy或Tb中的至少一种。
  9. 根据权利要求8中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于，还包括时效处理的步骤：对所述烧结磁铁在400℃～650℃的温度进行时效处理。
  10. 根据权利要求1或2或3或4中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于：所述烧结磁铁的O含量为0.1at％～0.5at％。
  11. 根据权利要求1或2或3或4中所述的一种含W的R‐Fe‐B‐Cu系烧结磁铁，其特征在于：所述烧结磁铁的Ga含量为0.05at％～0.8at％。
  12. 一种含W的R‐Fe‐B‐Cu系烧结磁铁用急冷合金，其特征在于：所述急冷合金的结晶晶界中具有W含量为0.004at％以上、0.26at％以下的富W区域，所述富W区域在所述结晶晶界中呈均一分散的分布，并占所述结晶晶界的至少50体积％。
PCT/CN2015/075512 2014-03-31 2015-03-31 一种含W的R‐Fe‐B‐Cu系烧结磁铁及急冷合金 Ceased WO2015149685A1 (zh)

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