HIGH-STRENGTH R-T-B RARE EARTH PERMANENT MAGNET AND PREPARATION METHOD THEREOF

Abstract

The present invention discloses a high-strength R-T-B rare earth permanent magnet and a preparation method thereof. The magnet contains 0.3-1.5 wt. % of an element Zr, and a cast strip prepared through vacuum induction melting and melt spinning is treated at a high temperature to make the element Zr therein precipitate in a form of fibrous Zr compounds from R-rich phases, and the fibrous Zr compounds can be uniformly mixed with magnetic powder after hydrogen decrepitation and powder jet milling and mixing, and gradually grow into rod-like Zr compounds existing in the R-rich intergranular phases during the sintering of a green compact. By adjusting the content of the element Zr, sintering temperature and time and other process parameters, the morphology, size and distribution of Zr compounds can be effectively controlled, and the mechanical properties of the magnet can be improved by strengthening the R-rich intergranular phases without deteriorating the magnetic properties of the magnet.

Claims

1. A high-strength R-T-B rare earth permanent magnet, comprising following components: 29.0-33.0 wt. % of R, wherein R is composed of R.sub.1, R.sub.2 and R.sub.H, R.sub.1 is a rare earth element Nd, R.sub.2 is a rare earth element Pr, and R.sub.H is at least one of rare earth elements Dy, Tb, Ho and Gd; the content of R.sub.2 is between 0.3 wt. % and 10 wt. %; the content of R.sub.H is between 0.02 wt. % and 8.5 wt. %; and the balance of R is R.sub.1, 0.90-1.1 wt. % of B, 5.0 wt. % or less of M, wherein M is composed of M.sub.1, M.sub.2, M.sub.3 and/or M.sub.4, M.sub.1 is an element Al, M.sub.2 is an element Cu, M.sub.3 is an element Ga, and M.sub.4 is at least one of elements Si, Zn, Sn, Ge, Ag and Au; the content of M.sub.1 is between 0.1 wt. % and 1.5 wt. %; the content of M.sub.2 is between 0.01 wt. % and 0.55 wt. %; the content of M.sub.3 is between 0.01 wt. % and 0.6 wt. %; when M contains M.sub.4, the content of M.sub.4 is less than or equal to 3.0 wt. %, 0.3-1.5 wt. % of Zr, T and other unavoidable impurities as the balance, wherein T is at least one of Fe and Co, and more than 85 wt. % of T is Fe, the following formulas are met:
[Zr]/91.22≥0.0025[T]/56  (1),
[B]/10.81≥4[Zr]/91.22  (2), wherein [Zr] is the content of Zr expressed by mass percent, [B] is the content of B expressed by mass percent, and [T] is the content of T expressed by mass percent; and the magnet is of a micro-structure consisting of main phase R2T14B, R-rich intergranular phases and rod-like Zr compounds, wherein the rod-like Zr compounds are mainly distributed in the R-rich intergranular phases of the magnet, and the number of the rod-like Zr compounds in the main phase grains accounts for less than 2.0% of the total number of the Zr compounds.

2. The high-strength R-T-B rare earth permanent magnet according to claim 1, wherein the high-strength R-T-B rare earth permanent magnet is prepared by following steps of: preparing an cast strip from proportioned raw materials of the high-strength R-T-B rare earth permanent magnet through vacuum induction melting and melt spinning, treating the cast strip at a high temperature to obtain a powder through hydrogen decrepitation and jet milling, molding the powder in an oriented magnetic field after particle size distribution optimization, and preparing the molded magnet into the high-strength R-T-B rare earth permanent magnet through vacuum sintering and aging treatment; and treating the cast strip in argon gas at a high temperature of 900-1030° C. and a pressure of 30-50 kPa for 30 minutes to 4 hours.

3. The high-strength R-T-B rare earth permanent magnet according to claim 2, wherein the melting temperature for vacuum induction melting and melt spinning is 1480-1510° C., and the cast strip is prepared through melt spinning at 1440-1460° C.

4. The high-strength R-T-B rare earth permanent magnet according to claim 2, wherein, during particle size distribution optimization, the powder after jet milling is classified by a powder classification device to obtain powder containing more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm.

5. The high-strength R-T-B rare earth permanent magnet according to claim 2, wherein the vacuum sintering temperature is held at 1080-1120° C. for 4 hours to 20 hours.

6. The high-strength R-T-B rare earth permanent magnet according to claim 2, wherein the aging process comprises a first aging stage in which the temperature is held at 700-900° C. for 2 hours to 8 hours and a second aging stage in which the temperature is held at 400-600° C. for 2 hours to 8 hours.

7. The high-strength R-T-B rare earth permanent magnet according to claim 1, wherein the rod-like Zr compounds have a length of 0.5-2.6 μm and an aspect ratio of 2 to 10.

8. The high-strength R-T-B rare earth permanent magnet according to claim 1, wherein the Zr compounds distributed in a micro-structure of the magnet have an area density of 1 to 6 Zr compounds per 100 μm.sup.2.

9. The high-strength R-T-B rare earth permanent magnet according to claim 1, wherein the total content of elements Ti, Nb, Hf and W in the magnet is lower than 0.01 wt. %.

10. A preparation method of the high-strength R-T-B rare earth permanent magnet according to claim 1, comprising following steps of: preparing a cast strip from proportioned raw materials of the high-strength R-T-B rare earth permanent magnet through vacuum induction melting and melt spinning, treating the cast strip at a high temperature to obtain a powder through hydrogen decrepitation and jet milling, compressing the powder in an oriented magnetic field after particle size distribution optimization, and preparing the molded magnet into the high-strength R-T-B rare earth permanent magnet through vacuum sintering and aging treatment; and during vacuum induction melting and melt spinning, the raw materials are melted into alloy melt at 1480-1510° C., and the alloy melt is poured on a rotating copper roller at 1440-1460° C. by a tundish for solidification to obtain the cast strip; treating the cast strip in argon gas at a high temperature of 900-1030° C. and a pressure of 30-50 kPa for 30 minutes to 4 hours; during particle size distribution optimization, the powder after jet milling is classified by a powder classification device to obtain powder containing more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm; and the vacuum sintering temperature is held at 1080-1120° C. for 4 hours to 20 hours.

Description

BRIEF DESCRIPTION OF THE DRAWINGS

[0046] FIGS. 1(a) and (b) are micro-structures of magnets in Experiment No. 5 and Experiment No. 8, respectively;

[0047] FIGS. 2(a) and (b) are micro-structures of the magnets in Experiment No. 21 and Experiment No. 23, respectively;

[0048] FIG. 3 is a micro-structure of magnets in Experiment No. 29;

[0049] FIGS. 4(a) and (b) are micro-structures of cast strips in Experiment No. 30 and Experiment No. 31, respectively; and

[0050] FIG. 5 is a micro-structure of magnets in Experiment No. 31.

DETAILED DESCRIPTION OF THE EMBODIMENTS OF THE INVENTION

[0051] A technical solution of the present invention will be further described with reference to embodiments, but the scope of protection of the present invention is not limited thereto.

[0052] According to the present invention, during vacuum induction melting and melt spinning, proportioned raw materials with a purity over 99.9% were put into a crucible in descending order of melting point, and the vacuum degree in a furnace reached 10.sup.−3 Pa to 10.sup.−4 Pa and the dew point was lower than −50° C. after vacuumizing. Argon gas was filled into a furnace to make the gas pressure reach 30-50 kPa, the furnace was heated to 1480-1510° C., and the temperature was held for 3-5 minutes after the raw materials were completely melted. The temperature of an alloy melt obtained was reduced to 1440-1460° C. and held for pouring. The speed of a copper roller was adjusted to 70-75 rpm, and the crucible rotated at a certain speed, so that the alloy melt was conveyed to a cooling roller through a tundish for solidification, and cooled down after falling on a water-cooling tray.

[0053] A cast strip prepared by melting was treated at high temperature, and was put in a molybdenum boat and then in the vacuum sintering furnace. The furnace was heated to 900-1030° C. after the vacuum degree therein reached 10.sup.−3 Pa to 10.sup.−4 Pa. Argon gas at a pressure of 30-50 kPa was filled in the furnace when the temperature reached a target temperature and was held for 30 minutes to 4 hours. The furnace was cooled down to room temperature after heat treatment. This process had no requirement on cooling speed, so either furnace cooling or air cooling might be adopted.

[0054] The alloy sheet was prepared into a powder through hydrogen decrepitation and jet milling. The hydrogen pressure in a reaction kettle was generally 0.01-0.09 MPa during hydrogen decrepitation, and the hydrogen absorption process ended when the pressure change in the reaction kettle did not exceed 0.5% within 10 minutes in hydrogen abstraction reactions. After the hydrogen absorption reactions, the temperature was raised to 400-600° C. while vacuumizing and held for 2 hours to 6 hours to release the hydrogen gas from the cast strip, and the cast strip was cooled down to obtain coarse powder through hydrogen decrepitation. The obtained coarse powder was put into a jet mill, and the pressure of a nozzle was adjusted to 0.6-0.8 MPa, and the coarse powder particles were crushed by colliding with each other under the action of the high-speed gas, and the inert gas was generally helium and nitrogen. A turbo selector and a cyclone separator of the jet mill were controlled to obtain powder of different particle sizes.

[0055] The powder after jet milling was further classified by a powder classification device for particle size distribution optimization. The powder containing more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm was preferred.

[0056] The powder might be added with lubricant and/or antioxidant before compression in an orientated magnetic field, and the conventional lubricant or antioxidant was available for the purpose of protecting magnetic powder. The dosage of the lubricant might be 0.01-0.1% by mass of the powder, and that of the antioxidant might be 0.01-0.14% by mass of the powder.

[0057] The orientation magnetic field was preferably 3-6 T, and the compressing pressure was 5-7 MPa. Cold isostatic pressing was performed on a green compact after orienting compression at a pressure of 150-180 MPa. The green compact density was 3.6-4.0 g/cm.sup.3 after orientated compressing and 4.6 g/cm.sup.3 after cold isostatic pressing.

[0058] Each magnet was densified by vacuum sintering. The vacuum sintering process was performed under a vacuum degree of 10.sup.−3 Pa to 10.sup.−4 Pa and at a temperature of 1080-1120° C. for 4 hours to 20 hours. In order to avoid the rare earth elements on a surface layer of the magnet from volatilizing during high-temperature sintering, inert gas at a pressure of 30-50 kPa was filled into a sintering furnace after a target sintering temperature was reached, and the inert gas might be argon gas and helium gas.

[0059] The sintered magnet needed to go through two aging stages, i.e., the sintered magnet was aged at 700-900° C. for 2 hours to 8 hours in a first aging stage, cooled down to below 100° C. at a speed of not less than 20° C./min, then aged for 2 hours to 8 hours in a second aging stage at 400-600° C., and finally cooled down to below 80° C. at a speed of not less than 30° C./min.

[0060] The magnets were crushed and sampled at the center to detect the composition thereof by ICP-MS. The micro-structures of the magnets were observed by Scanning Electron Microscopy (SEM), and the micro-area composition of the magnets was analyzed by Electron Probe Micro-Analysis (EPMA). Three-point bending specimens with dimensions of 25(±0.01) mm×6(±0.01) mm×5(±0.01 mm) were fabricated by an inside diameter slicer and a double-sided grinding machine, with the height direction thereof parallel to the orientation direction of the magnets. According to GB/T31967.2-2015, the bending strength of the magnets shall be measured by a three-point bending method. In each experimental group, 10 specimens were measured to take the average value through calculation. The three-point bending indenter was a cylinder with a diameter of 5 mm and a down speed of 0.1 mm/min, and two support columns had a diameter of 5 mm and a spacing of 14.5 mm. Specimens with dimensions of ϕ10 mm×10 mm for magnetic property measurement were prepared by wire-electrode cutting, double-sided grinding and face grinding, and the magnetic properties of the magnets were measured by NIM equipment.

Embodiment One

[0061] When taking low melting point metals as raw materials, metals with purity of more than 99.9 wt. % should be adopted, and when taking elements with a melting point higher than pure iron as raw materials, an alloy of these elements and iron should be adopted. The raw materials were put in a crucible in descending order of melting point, and the vacuum degree in a furnace reached 10.sup.−4 Pa and the dew point was lower than −50° C. after vacuumizing. Argon gas was filled into the furnace to make the gas pressure reach 30 kPa, the furnace was heated to 1490° C., and the temperature was held for 3 minutes after the raw materials were completely melted. An alloy melt obtained was cooled down to 1450° C. for pouring. The speed of a copper roller was adjusted to 70 rpm, and the crucible rotated at a certain speed, so that the alloy melt was conveyed to a cooling roller through a tundish for solidification, and cooled down after falling on a water-cooling tray, to obtain a cast strip with a thickness of 0.25±0.05 mm.

[0062] The cast strip prepared by melting was treated at a high temperature, and was put in a molybdenum boat and then in the heating furnace. The furnace was heated to 1000° C. after the vacuum degree therein reached 10.sup.−4 Pa. Argon gas at a pressure of 30 kPa was filled in the furnace when the temperature reached a target temperature and was held for 2 hours. The furnace was cooled down to room temperature after heat treatment.

[0063] The alloy sheet was subjected to hydrogen absorption reactions at a hydrogen pressure of 0.09 MPa. After the hydrogen absorption reactions, the temperature was raised to 550° C. while vacuumizing and held for 4 hours to release the hydrogen gas from the cast strip, and the cast strip was cooled down to obtain coarse powder through hydrogen decrepitation. After cooling, 0.05 wt. % of zinc stearate was added to the coarse powder and mixed for 3 hours. The mixed coarse powder was further milled by jet milling with nitrogen gas to obtain fines, and the nitrogen pressure was 0.6 MPa. The fines after jet milling were further classified by a powder classification device, so that the fines contained more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm.

[0064] Then, 0.03 wt. % of organic lubricant (magnetic powder lubricant 3# produced by Tianjin Yuesheng New Materials Research Institute) was added to the fines and mixed for 3 hours. The uniformly mixed fines were compressed in an oriented magnetic field which was a 3.5 T static magnetic field at a pressing pressure of 5 MPa, and the density of the pressed magnet was 3.9-4.0 g/cm.sup.3. Cold isostatic pressing was performed at a pressure of 160 MPa, and the density of the pressed magnet was greater than 4.6 g/cm.sup.3.

[0065] Each magnet was densified by vacuum sintering. The vacuum sintering process was that the magnet was sintered at a vacuum degree of 10.sup.−4 Pa and a temperature of 1080-1120° C. to ensure that the density of the sintered magnet was at least 7.53 g/cm.sup.3, and the temperature was held for 4 hours to 20 hours. In order to avoid the rare earth elements on a surface layer of the magnet from volatilizing during high-temperature sintering, argon gas at 30 kPa was filled into a sintering furnace after a target sintering temperature was reached.

[0066] The sintered magnet needed to go through two aging stages, i.e., the sintered magnet was aged at 860° C. for 3 hours in a first aging stage, cooled down to below 100° C. at a speed of not less than 20° C./min, aged for 3 hours in a second aging stage at 520° C., and finally cooled down to below 80° C. at a speed of not less than 30° C./min.

[0067] The magnets were crushed and sampled at the center to detect the composition thereof by ICP-MS. The micro-structures of the magnets were observed by Scanning Electron Microscopy (SEM), and the micro-area composition of the magnets was analyzed by Electron Probe Micro-Analysis (EPMA). According to GB/T31967.2-2015, the bending strength of the magnets shall be measured by a three-point bending method. In each experimental group, 10 specimens were measured to take the average value through calculation.

[0068] The composition of the magnets in each experimental group was expressed by mass percent, as shown in Table 1.

TABLE-US-00001 TABLE 1 No. Nd Pr Fe Al Dy Co Cu Ga B Zr Ti Nb Formula (1) Formula (2) 1 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0 / / / / 2 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.05 / / Non- Conforming conforming 3 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.1 / / Non- Conforming conforming 4 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.28 / / Conforming Conforming 5 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.30 / / Conforming Conforming 6 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5 / / Conforming Conforming 7 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.6 / / Conforming Conforming 8 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 1 / / Conforming Conforming 9 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 1.5 / / Conforming Conforming 10 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 2 / / Conforming Conforming 11 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 2.5 / / Conforming Non- conforming 12 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 3 / / Conforming Non- conforming 13 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5  0.002 / Conforming Conforming 14 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5 0.2 / Conforming Conforming 15 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5 0.5 / Conforming Conforming 16 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5 /  0.002 Conforming Conforming 17 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5 / 0.2 Conforming Conforming 18 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5 / 0.5 Conforming Conforming 19 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.5 0.2 0.2 Conforming Conforming

[0069] The bending strength of the magnets should be measured by a three-point bending method. In each experimental group, 10 specimens were measured to take the average value through calculation. In the back-scattering pattern of SEM, the size, the morphology, and the calculated proportion of precipitates in the R-rich intergranular phases and the calculated distribution density of precipitates (the quantity of precipitates in an area of 100 μm.sup.2) in an area of 1000 μm×1000 μm were shown in Table 2 below.

TABLE-US-00002 TABLE 2 Average Proportion of Dis- length of precipitates in tribution Bending pre- R-rich density Morphology strength/ cipitates intergranular (Number/ of Zr No. MPa (μm) phases (%) 100 μm.sup.2) compounds 1 360 / / / / 2 365 / / / / 3 380 / / / / 4 410 0.21 100 0.3 Rod-like 5 525 0.51 100 2 Rod-like 6 580 0.85 100 4 Rod-like 7 620 0.93 99.8 4 Rod-like 8 550 1.8 99.6 5 Rod-like 9 465 2.53 99.2 6 Rod-like 10 355 2.64 94.6 8 Rod-like 11 320 1.56 92.1 9 Blocky 12 280 1.83 88.4 9 Blocky 13 560 0.79 100 4 Rod-like 14 500 0.63 98.6 3 Rod-like 15 480 0.59 96.7 3 Blocky 16 565 0.82 100 4 Rod-like 17 515 0.61 98.8 4 Rod-like 18 465 0.6 97.5 2 Blocky 19 420 0.74 96.9 2 Blocky

[0070] Microstructures of the magnets were observed by SEM, as shown in FIGS. 1(a) and 1 (b). The figures showed that Zr compounds could be found in the intergranular phases of the magnet when the content of the element Zr in the magnet was 0.3 wt. %. However, due to the low dosage of the element Zr, there was less quantity of Zr compounds in the magnet. On the contrary, when the content of the element Zr in the magnet increased to 1.0 wt. %, the number of Zr compounds in the R-rich intergranular phases obviously increased.

[0071] The micro-area composition of the magnets was analyzed by EPMA spot scanning. The composition of precipitates from the magnet in Experiment No. 5 and experiment No. 8 was shown in Table 3.

TABLE-US-00003 TABLE 3 Experiment Scanning No. spot Nd Pr Fe Zr B Co Al 5 1 27.57 9.53 36.97 20.5 0.35 4.72 0.36 2 21.78 10.86 39.5 26.77 0.38 / 0.71 8 3 22 7.8 25.45 43.51 0.64 / 0.6 4 4.79 4.03 4.79 84.21 1.6 / 0.58 5 22.14 9.72 24.36 42.5 0.58 / 0.70 6 23.17 10.04 23.84 41.96 0.67 / 0.32

[0072] The composition of precipitates showed that the Zr compounds had a high content of the element Zr ranging from 20 wt. % to 85 wt. %. As the element Zr could replace the element R in the main phase, when the magnet contained a low concentration of the element Zr, the element Zr was not enough to involve in the precipitation, resulting in the failure to precipitate Zr compounds from the cast strip during high-temperature treatment. Therefore, in the present invention, only when the element Zr in the magnet exceeded a certain concentration could the precipitation of Zr compounds be obviously promoted.

[0073] Based on the data from each experimental group, it could be concluded that Zr compounds could be found in the magnet only when the content of the element Zr in the magnet conformed to Formula (1). When the content of the element Zr in the magnet conformed to the Formula (1) but was less than 0.3 wt. %, only a small amount of fine fibrous Zr compounds were precipitated. Due to the low distribution density and too small size of precipitates, these fine precipitates were easy to fracture along with the R-rich intergranular phases when the magnet was stressed, so the mechanical properties of the magnet were improved to a limited extent. The distribution density and average length of precipitates increased gradually when the content of the element Zr in the alloy increased to a value within a recommended range (0.3-1.5 wt. %), and the fibrous Zr compounds precipitated during the high-temperature treatment of the sintered cast strip grew into rod-like Zr compounds. Because the rod-like precipitates could withstand higher load, the intergranular phases of the magnets could be strengthened by changing the propagation direction of cracks. At the same time, the pulling effect of rod-like precipitates under stress might also consume more fracture energy, so the bending strength of the magnets could be effectively improved by the two effects.

[0074] The content of the element B in the magnets also affected the precipitation of Zr compounds. As shown in Experiment No. 11 and Experiment No. 12, when the content of the element B did not conform to the Formula (2), the Zr compounds precipitated out of the cast strip after high-temperature treatment were blocky compounds that would grow into blocky coarse precipitates after sintering. These blocky coarse precipitates contained more than 85 wt. % of the element Zr generally distributed in the R-rich intergranular phases and in direct contact with the main phase grains. Due to the poor wettability between the coarse precipitates and the main phase grains, the poor interface bonding strength between the coarse precipitates and the main phase grains easily led to stress concentration and deteriorated the mechanical properties of the magnets.

[0075] Ti, Nb and other elements that could react with the element with B to produce precipitates which would affect the precipitation of Zr compounds when existing in the magnets. Experiments No. 13 to No. 19 showed that when the content of Ti, Nb and other elements in the magnets increased, the average length of Zr compounds decreased and the bending strength deteriorated accordingly. The reaction of Ti, Nb and other elements with the element B affected the precipitation of Zr compounds during high temperature treatment and the further growth of Zr compounds during sintering, thus changing the size and precipitation of Zr compounds. The experimental results showed that the average length of Zr compounds in the magnets decreased with the addition of Ti and Nb, and the Zr compounds gradually changed from rod-like compounds to bulky compounds with the increase of Ti and Nb content. As a result, the strengthening effect of precipitates on the R-rich intergranular phases was weakened, and the bending strength of the magnets deteriorated. Therefore, when elements such as Ti, Nb, Hf and W existed as inevitable impurities, the total content of these elements should be less than 0.01 wt. %.

Embodiment Two

[0076] When taking low melting point metals as raw materials, metals with purity of more than 99.9 wt. % should be adopted, and when taking elements with a melting point higher than pure iron as raw materials, an alloy of these elements and iron should be adopted. The raw materials were put in a crucible in descending order of melting point, and the vacuum degree in a furnace reached 10.sup.−4 Pa and the dew point was lower than −50° C. after vacuumizing. Argon gas was filled into the furnace to make the gas pressure reach 30 kPa, the furnace was heated to 1490° C., and the temperature was held for 3 minutes after the raw materials were completely melted. An alloy melt obtained was cooled down to 1450° C. for pouring. The speed of a copper roller was adjusted to 70 rpm, and the crucible rotated at a certain speed, so that the alloy melt was conveyed to a cooling roller through a tundish for solidification, and cooled down after falling on a water-cooling tray, to obtain a cast strip with a thickness of 0.25±0.05 mm.

[0077] The cast strip prepared by melting was treated at a high temperature, and was put in a molybdenum boat and then in the heating furnace. The furnace was heated to 1000° C. after the vacuum degree therein reached 10.sup.−4 Pa. Argon gas at a pressure of 30 kPa was filled in the furnace when the temperature reached a target temperature and was held for 2 hours. The furnace was cooled down to room temperature after heat treatment.

[0078] The alloy sheet was subjected to hydrogen absorption reactions at a hydrogen pressure of 0.09 MPa. After the hydrogen absorption reactions, the temperature was raised to 550° C. while vacuumizing and held for 4 hours to release the hydrogen gas from the cast strip, and the cast strip was cooled down to obtain coarse powder through hydrogen decrepitation. After cooling, 0.05 wt. % of zinc stearate was added to the coarse powder and mixed for 3 hours. The mixed coarse powder was further milled by jet milling with nitrogen gas to obtain fines, and the nitrogen pressure was 0.6 MPa. The fines after jet milling were further classified by a powder classification device, so that the fines contained more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm.

[0079] Then, 0.03 wt. % of organic lubricant (magnetic powder lubricant 3# produced by Tianjin Yuesheng New Materials Research Institute) was added to the fines and mixed for 3 hours. The uniformly mixed fines were compressed in an oriented magnetic field which was a 3.5 T static magnetic field at a pressing pressure of 5 MPa, and the density of the pressed magnet was 3.9-4.0 g/cm.sup.3. Cold isostatic pressing was performed at a pressure of 160 MPa, and the density of the pressed magnet was greater than 4.6 g/cm.sup.3.

[0080] Each magnet was densified by vacuum sintering. The vacuum sintering process was that the magnet was sintered at a vacuum degree of 10.sup.−4 Pa and a temperature of 1080-1120° C. to ensure that the density of the sintered magnet was at least 7.53 g/cm.sup.3. The sintering process in different experimental groups was shown in Table 5. In order to avoid the rare earth elements on a surface layer of the magnets from volatilizing during high-temperature sintering, argon gas at 30 kPa was filled into a sintering furnace after a target sintering temperature was reached.

[0081] The sintered magnet needed to go through two aging stages, i.e., the sintered magnet was aged at 860° C. for 3 hours in a first aging stage, cooled down to below 100° C. at a speed of not less than 20° C./min, then aged for 3 hours in a second aging stage at 520° C., and finally cooled down to below 80° C. at a speed of not less than 30° C./min.

[0082] Micro-structures of the magnets were observed by Scanning Electron Microscopy (SEM), and the size and quantity of precipitates and the size of main phase grains were recorded in the back-scattering pattern of SEM. The magnets were crushed and sampled at the center to detect the composition thereof by ICP-MS. The composition of the magnets in each experimental group was expressed by mass fraction, as shown in Table 4.

TABLE-US-00004 TABLE 4 No. Nd Pr Fe Al Dy Co Cu Ga B Zr Ti Nb Formula (1) Formula (2) 20~23 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.6 / / Conforming Conforming 24 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 / 0.6 / / / 25 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 / / 0.6 / / 26 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 / 0.3 0.3 / / 27 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 /  0.25 / / /

[0083] The sintering process in different experimental groups was shown in Table 5.

TABLE-US-00005 TABLE 5 No. Sintering temperature/° C. Holding time/h 20 1060 12 21 1080 7 22 1100 5 23 1130 4 24 1100 5 25 1100 5 26 1100 5 27 1100 5

[0084] The bending strength of the magnets should be measured by a three-point bending method. In each experimental group, 10 specimens were measured to take the average value and the standard deviation through calculation. In the back-scattering pattern of SEM, the size and quantity of precipitates, and the calculated average length and aspect ratio of precipitates in an area of 1000 μm×1000 μm were specifically shown in Table 6 below.

TABLE-US-00006 TABLE 6 Bending Average Maximum Minimum strength Standard length length length Aspect No. (MPa) deviation Type Morphology (μm) (μm) (μm) ratio 20 360 25.8 Zr Rod-like 0.42 1.26 0.1 2-4  compounds 21 585 15.6 Zr Rod-like 0.49 1.93 0.1 2-10 compounds 22 620 16.3 Zr Rod-like 0.89 2.33 0.32 2-10 compounds 23 490 32.8 Zr Rod-like 1.68 6.14 0.44 3-15 compounds 24 470 65.2 Ti Fibrous/rod- / / / / compounds like/blocky 25 455 68.7 Nb Fibrous/rod- / / / / compounds like/blocky 26 480 55.4 Ti/Nb Fibrous/rod- / / / / compounds like/blocky 27 425 63.5 Ti Granular/fibrous / / / / compounds

[0085] Because the magnets in Experiment No. 20 to No. 23 contained a high concentration of the element Zr, it was difficult to densify the magnets by sintering at low temperature. The density of the specimens in Experiment No. 20 was still not up to standard after the temperature was held at 1060° C. for 12 hours, so the magnets in the present invention needed to be sintered at higher temperature.

[0086] It was found through analysis that the length of Zr compounds precipitated out of the cast strip during high-temperature treatment could be adjusted by changing the sintering temperature. In Experiments No. 20 to No. 23, the difference between average grain sizes of the magnets in different experimental groups did not exceed 0.5 μm at different sintering temperatures, provided that the density of the magnets was ensured to reach the standard. By analyzing the size of Zr compounds and mechanical properties of the magnets at different sintering temperatures, it was found that the length of Zr compounds in the magnets increased with the increase of sintering temperature. The micro-structure (shown in FIG. 2 (b)) in Experiment No. 23 (FIG. 2 (b)) showed that there were Zr compounds with a length of 6.14 μm when the sintering temperature was 1130° C. In this experimental group, the bending strength of the magnets first increased and then decreased with the increase of the length of Zr compounds, and reached its maximum when the average length of Zr compounds was 0.89 μm. The rod-like Zr compounds of appropriate sizes distributed in the intergranular phases could strengthen the R-rich intergranular phases and changed the propagation direction of cracks. At the same time, the pulling effect of rod-like precipitates under stress might absorb more fracture energy, so the mechanical properties of the magnets could be effectively improved. The Zr compounds of larger size in the magnets would be interspersed in and in direct contact with several main phase grains. Due to the poor wettability between Zr compounds and the main phase grains, the interface between the two phases would initiate cracks, so too large size of the precipitated phase would deteriorate the mechanical properties of the magnets.

[0087] In this embodiment, the sintering time at different temperatures was controlled to avoid significant change in the grain sizes in each experimental group as well as the influence of the size of the main phase grains on the mechanical properties of the magnets, provided that the density of the magnets was ensured to reach the standard. The experiments showed that only the Zr compounds with a length of 0.5-2.6 μm and an aspect ratio of 2-10 could strengthen the mechanical properties of the magnets.

[0088] When added to magnets alone, elements such as Nb and Ti also reacted with the element with B to produce precipitates. However, the experiments showed that the morphology of compounds produced by reactions between elements such as Nb and Ti and the element Zr was more diverse, including fibrous, rod-like and blocky compounds. Even granular precipitates could be found in case of a low dosage of these elements. In addition, it was difficult to control the size of Nb and Ti precipitates and the proportion of precipitates difference in morphology by controlling the sintering process. Although these precipitates in the magnets could also strengthen the R-rich intergranular phases and improve the bending strength of the magnets to a certain extent. However, due to the deviations in morphology and size of different precipitates, the mechanical properties of different parts of the magnets differed sharply, indicating that the standard deviation of the bending strength of this type of magnets was significantly higher than that of magnets containing the element Zr alone. Therefore, in order to ensure the consistency of mechanical properties of the magnets, the proportion of these elements in the magnets was minimized, and the content of these elements was less than 0.01 wt. % when such elements existed as inevitable impurities.

Embodiment Three

[0089] When taking low melting point metals as raw materials, metals with purity of more than 99.9 wt. % should be adopted, and when taking elements with a melting point higher than pure iron as raw materials, an alloy of these elements and iron should be adopted. The raw materials were put in a crucible in descending order of melting point, and the vacuum degree in a furnace reached 10.sup.−4 Pa and the dew point was lower than −50° C. after vacuumizing. Argon gas was filled into the furnace to make the gas pressure reach 30 kPa, the furnace was heated to 1490° C., and the temperature was held for 3 minutes after the raw materials were completely melted. An alloy melt obtained was cooled down to 1450° C. for pouring. The speed of a copper roller was adjusted to 70 rpm, and the crucible rotated at a certain speed, so that the alloy melt was conveyed to a cooling roller through a tundish for solidification, and cooled down after falling on a water-cooling tray, to obtain a cast strip with a thickness of 0.25±0.05 mm.

[0090] The cast strip prepared by melting was treated at a high temperature, and was put in a molybdenum boat and then in the heating furnace. The furnace was heated to 1000° C. after the vacuum degree therein reached 10.sup.−4 Pa. Argon gas at a pressure of 30 kPa was filled in the furnace when the temperature reached a target temperature and was held for 2 hours. The furnace was cooled down to room temperature after heat treatment.

[0091] The alloy sheet was subjected to hydrogen absorption reactions at a hydrogen pressure of 0.09 MPa. After the hydrogen absorption reactions, the temperature was raised to 550° C. while vacuumizing and held for 4 hours to release the hydrogen gas from the cast strip, and the cast strip was cooled down to obtain coarse powder through hydrogen decrepitation. After cooling, 0.05 wt. % of zinc stearate was added to the coarse powder and mixed for 3 h. The mixed coarse powder was further milled by jet milling with nitrogen gas to obtain fines, and the nitrogen pressure was 0.6 Mpa. Some fines after jet milling were further classified by a powder classification device, so that the fines contained more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm. The fines without being classified by the powder classification device contained 82% of powder particles with SMD of 3.0-6.0 μm.

[0092] Then, 0.03 wt. % of organic lubricant (magnetic powder lubricant 3# produced by Tianjin Yuesheng New Materials Research Institute) was added to the fines and mixed for 3 hours. The uniformly mixed fines were compressed in an oriented magnetic field which was a 3.5 T static magnetic field at a pressing pressure of 5 MPa, and the density of the pressed magnet was 3.9-4.0 g/cm.sup.3. Cold isostatic pressing was performed at a pressure of 160 MPa, and the density of the pressed magnet was greater than 4.6 g/cm.sup.3.

[0093] In this embodiment, the composition in all experimental groups was the same as that in Experiment No. 7, and the green compact was sintered in a vacuum sintering furnace for densification. The vacuum sintering process was performed at a vacuum degree of 10.sup.−4 Pa and a temperature of 1100° C. In order to avoid the rare earth elements on a surface layer of the magnet from volatilizing during high-temperature sintering, argon gas at 30 kPa was filled into a sintering furnace after a target sintering temperature was reached and held for 6 hours.

[0094] The sintered magnet needed to go through two aging stages, i.e., the sintered magnet was aged at 860° C. for 3 hours in a first aging stage, cooled down to below 100° C. at a speed of not less than 20° C./min, then aged for 3 hours in a second aging stage at 520° C., and finally cooled down to below 80° C. at a speed of not less than 30° C./min.

[0095] The bending strength of the sintered magnet was measured with a universal material tester by three-point bending tests. Microstructures of the magnets was observed by SEM, and the size and quantity of precipitates and the size of main phase grains in an area of 1000 μm×1000 μm were recorded in the back-scattering pattern of SEM.

[0096] The process characteristics, the size and distribution of precipitates, the size of main phase grains and the mechanical properties of the magnets in different experimental groups were shown in Table 7.

TABLE-US-00007 TABLE 7 Proportion Number of powder Average Maximum of grains Proportion of particles grain grain greater Bending precipitates in Powder with SMD of size size than 15 strength intergranular No. classification 3.0-6.0 μm (μm) (μm) μm/Nr. (MPa) phases 28 Yes 95% 8.2 17.9 23 615 99.8 29 No 82% 12.3 25.6 325 530 96.7

[0097] Comparison of the above experimental data showed that the grain sizes in Experiment No. 28 were more uniform, indicating that the jet-milled powder after powder classification could significantly reduce the proportion of ultrafine and coarse powder particles, and improve the size consistency of grains in the sintered magnets, and obviously improve the mechanical properties of the magnets. In addition, the uniform particle size distribution could affect the distribution of precipitates. When ultrafine powder particles dominated in the powder particles were melted or merge with surrounding large grains during sintering, Zr compounds were easily encapsulated in the main phase grains during grain growth, as shown in the micro-structure of the magnets in FIG. 3. The distribution of precipitates showed that the proportion of precipitates in the intergranular phases without powder classification in the magnets significantly decreased, i.e., the number of precipitates distributed in the main phase grains increased. These precipitates existing in the main phase grains would initiate cracks in the main phase grains, thus deteriorating the mechanical properties of the magnets. The powder classification could ensure the consistency of the main phase grains in the magnets and also could prevent the ultrafine power particles from encapsulating precipitates during the growth of surrounding grains, so as to give full play to the advantages of precipitates in improving the mechanical properties of the magnets.

Embodiment Four

[0098] When taking low melting point metals as raw materials, metals with purity of more than 99.9 wt. % should be adopted, and when taking elements with a melting point higher than pure iron as raw materials, an alloy of these elements and iron should be adopted. The raw materials were put in a crucible in descending order of melting point, and the vacuum degree in a furnace reached 10.sup.−4 Pa and the dew point was lower than −50° C. after vacuumizing. Argon gas was filled into the furnace to make the gas pressure reach 30 kPa, the furnace was heated to 1490° C., and the temperature was held for 3 minutes after the raw materials were completely melted. An alloy melt obtained was cooled down to 1450° C. for pouring. The speed of a copper roller was adjusted to 70 rpm, and the crucible rotated at a certain speed, so that the alloy melt was conveyed to a cooling roller through a tundish for solidification, and cooled down after falling on a water-cooling tray, to obtain a cast strip with a thickness of 0.25±0.05 mm.

[0099] For comparison, some cast strips prepared by melting were treated at high temperature while some were not. The cast strips prepared by melting were put in a molybdenum boat and then in the heating furnace. The furnace was heated to 1000° C. after the vacuum degree therein reached 10.sup.−4 Pa. Argon gas at a pressure of 30 kPa was filled in the furnace when the temperature reached a target temperature and was held for 2 hours. The furnace was cooled down to room temperature after heat treatment.

[0100] Each alloy sheet was subjected to hydrogen absorption reactions at a hydrogen pressure of 0.09 MPa. After the hydrogen absorption reactions, the temperature was raised to 550° C. while vacuumizing and held for 4 hours to release the hydrogen gas from the cast strip, and the cast strip was cooled down to obtain coarse powder through hydrogen decrepitation. After cooling, 0.05 wt. % of zinc stearate was added to the coarse powder and mixed for 3 hours. The mixed coarse powder was further milled by jet milling with nitrogen gas to obtain fines, and the nitrogen pressure was 0.6 MPa. Some fines after jet milling were further classified by a powder classification device, so that the fines contained more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm. The fines without being classified by the powder classification device contained 82% of powder particles with SMD of 3.0-6.0 μm.

[0101] Then, 0.03 wt. % of organic lubricant (magnetic powder lubricant 3# produced by Tianjin Yuesheng New Materials Research Institute) was added to the fines and mixed for 3 hours. The uniformly mixed fines were compressed in an oriented magnetic field which was a 3.5 T static magnetic field at a pressing pressure of 5 MPa, and the density of the pressed magnet was 3.9-4.0 g/cm.sup.3. Cold isostatic pressing was performed at a pressure of 160 MPa, and the density of the pressed magnet was greater than 4.6 g/cm.sup.3.

[0102] In this embodiment, the composition in all experimental groups was the same as that in Experiment No. 7, and the green compact was sintered in a vacuum sintering furnace for densification. The vacuum sintering process was performed at a vacuum degree of 10.sup.−4 Pa and a temperature of 1100° C. In order to avoid the rare earth elements on a surface layer of the magnet from volatilizing during high-temperature sintering, argon gas at 30 kPa was filled into a sintering furnace after a target sintering temperature was reached and held for 6 hours.

[0103] The sintered magnet needed to go through two aging stages, i.e., the sintered magnet was aged at 860° C. for 3 hours in a first aging stage, cooled down to below 100° C. at a speed of not less than 20° C./min, and aged for 3 hours in a second aging stage at 520° C., and finally cooled down to below 80° C. at a speed of not less than 30° C./min.

[0104] Micro-structures of the magnets were observed by SEM. FIGS. 4(a) and (b) showed the micro-structures of the cast strip after high-temperature treatment and without high-temperature treatment, respectively. The figures showed that no Zr compound could be found in the structure of the cast strips without high-temperature treatment, but clusters of fine fibrous Zr compounds could be found in the R-rich phase when the cast strip was treated for 2 hours at a high temperature of 1000° C. These Zr compounds were crushed and uniformly mixed with magnetic powder in the subsequent milling process, acting as a binary powder mixture.

[0105] The bending strength of the magnets should be measured by a three-point bending method. In each experimental group, 10 specimens were measured to take the average value through calculation. In the back-scattering pattern of SEM, the quantity and distribution of precipitates in an area of 1000 μm×1000 μm were specifically shown in Table 8 below.

TABLE-US-00008 TABLE 8 High-temperature Bending Proportion of precipitates Experiment treatment of the strength distributed in main phase No. cast strip (MPa) grains (%) 30 Yes 615 0.25 31 No 480 9.5

[0106] FIG. 5 showed a micro-structure of the magnets in Experiment No. 31. Based on the data in Table 8, it could be found that the main phase Zr compounds in the magnets were distributed in intergranular phases after the high-temperature treatment of cast strips, while a considerable proportion of Zr compounds in the magnets were distributed in the main phase grains without the high-temperature treatment of cast strips. At the same time, the bending strength of the magnets in Experiment No. 30 was significantly higher than that of the magnets in Experiment No. 31.

[0107] The high-temperature treatment was performed on the cast strips of the magnets in Experiment No. 30 before hydrogen decrepitation, which promoted the precipitation of Zr compounds out of the cast strips. In the subsequent preparation of powders, these Zr compounds would be crushed to a certain extent and uniformly mixed with magnetic powder, acting as a binary powder mixture. However, the powders obtained by this method had the advantages of high morphological consistency, uniform particle size distribution and no impurities. In the subsequent sintering process, these Zr compounds grew from fibrous to rod-like compounds existing in the intergranular phases of the magnets. The mechanical properties of the magnets could be improved by strengthening the R-rich intergranular phases of the magnets.

[0108] Since no high-temperature treatment was performed on the cast strips of the magnets in Experiment No. 31, the Zr compounds could only be precipitated out of the magnets in the sintering process during which the precipitation of Zr compounds occurred simultaneously with the magnet densification and the main phase grain growth, and the Zr compounds precipitated in this process was easily encapsulated in the main phase grains. Therefore, the proportion of precipitates in the main phase grains in the magnets in Experiment No. 31 was much higher than that in Experiment No. 30. These precipitates existing in the main phase grains initiated cracks in the main phase grains, thus deteriorating the mechanical properties of the magnets.

Embodiment Five

[0109] When taking low melting point metals as raw materials, metals with purity of more than 99.9 wt. % should be adopted, and when taking elements with a melting point higher than pure iron as raw materials, an alloy of these elements and iron should be adopted. The raw materials were put in a crucible in descending order of melting point, and the vacuum degree in a furnace reached 10.sup.−4 Pa and the dew point was lower than −50° C. after vacuumizing. Argon gas was filled into the furnace to make the gas pressure reach 30 kPa, the furnace was heated to 1490° C., and the temperature was held for 3 minutes after the raw materials were completely melted. An alloy melt obtained was cooled down to 1450° C. for pouring. The speed of a copper roller was adjusted to 70 rpm, and the crucible rotated at a certain speed, so that the alloy melt was conveyed to a cooling roller through a tundish for solidification, and cooled down after falling on a water-cooling tray, to obtain a cast strip with a thickness of 0.25±0.05 mm.

[0110] The cast strip prepared by melting was treated at a high temperature, and was put in a molybdenum boat and then in the heating furnace. The furnace was heated to 1000° C. after the vacuum degree therein reached 10.sup.−4 Pa. Argon gas at a pressure of 30 kPa was filled in the furnace when the temperature reached a target temperature and was held for 2 hours. The furnace was cooled down to room temperature after heat treatment.

[0111] The alloy sheet was subjected to hydrogen absorption reactions at a hydrogen pressure of 0.09 MPa. After the hydrogen absorption reactions, the temperature was raised to 550° C. while vacuumizing and held for 4 hours to release the hydrogen gas from the cast strip, and the cast strip was cooled down to obtain coarse powder through hydrogen decrepitation. After cooling, 0.05 wt. % of zinc stearate was added to the coarse powder and mixed for 3 hours. The mixed coarse powder was further milled by jet milling with nitrogen gas to obtain fine particles, and the nitrogen pressure was 0.6 MPa. The fines after jet milling were further classified by a powder classification device, so that the fines contained more than 95% of powder particles with a surface mean diameter (SMD) of 3.0-6.0 μm.

[0112] Then, 0.03 wt. % of organic lubricant (magnetic powder lubricant 3# produced by Tianjin Yuesheng New Materials Research Institute) was added to the fines and mixed for 3 hours. The uniformly mixed fines were compressed in an oriented magnetic field which was a 3.5 T static magnetic field at a pressing pressure of 5 MPa, and the density of the pressed magnet was 3.9-4.0 g/cm.sup.3. Cold isostatic pressing was performed at a pressure of 160 MPa, and the density of the pressed magnet was greater than 4.6 g/cm.sup.3.

[0113] Each magnet was densified by vacuum sintering. The vacuum sintering process was performed at a vacuum degree of 10.sup.−4 Pa and a temperature of 1100° C. to ensure that the density of the sintered magnet was at least 7.53 g/cm.sup.3. In order to avoid the rare earth elements on a surface layer of the magnet from volatilizing during high-temperature sintering, argon gas at 30 kPa was filled into a sintering furnace after a target sintering temperature was reached and held for 6 h.

[0114] The sintered magnet needed to go through two aging stages, i.e., the sintered magnet was aged at 860° C. for 3 hours in a first aging stage, cooled down to below 100° C. at a speed of not less than 20° C./min, then aged for 3 hours in a second aging stage at 520° C., and finally cooled down to below 80° C. at a speed of not less than 30° C./min.

[0115] The magnets were crushed and sampled at the center to detect the composition thereof by ICP-MS. Three-point bending specimens were fabricated by an inside diameter slicer and a double-sided grinding machine. According to GB/T31967.2-2015, the bending strength of the magnets should be measured by a three-point bending method. In each experimental group, 10 specimens were measured to take the average value through calculation. Specimens with dimensions of ϕ10 mm×10 mm for magnetic property measurement were prepared by wire-electrode cutting, double-sided grinding and face grinding, and the magnetic properties of the magnets were measured by NIM equipment.

[0116] The composition of the magnets in each experimental group was expressed by mass fraction, as shown in Table 9.

TABLE-US-00009 TABLE 9 No. Nd Pr Fe Al Dy Co Cu Ga B Zr 32 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 / 33 30.6 0.5 Bal 1 0.5 0.1 0.15 0.1 0.96 0.6 34 31.1 / Bal 1 0.5 0.1 0.15 0.1 0.96 0.6 35 31.6 0.5 Bal / 0.5 0.1 0.15 0.1 0.96 0.6 36 30.75 0.5 Bal 1 0.5 0.1 / 0.1 0.96 0.6 37 30.7 0.5 Bal 1 0.5 0.1 0.15 / 0.96 0.6

[0117] Mechanical properties and coercivity of the magnet were shown in Table 10.

TABLE-US-00010 TABLE 10 No. Bending strength (MPa) Coercivity (kOe) 32 360 20.13 33 618 20.09 34 602 19.65 35 613 18.24 36 596 19.80 37 620 19.74

[0118] In this embodiment, by adjusting the composition of the magnets, some elements in Pr, Al, Cu and Ga might be selectively removed, and the content of removed elements was replaced by the same weight percentage of an element Nd. The data such as magnet composition, mechanical properties and coercivity showed that the change in the content of elements such as Pr, Al, Cu and Ga would not affect the bending strength of the magnets containing 0.6 wt. % of the element Zr, but would affect the magnetic properties of the magnets to a great extent. If elements Al, Ga, Cu and Pr were added to magnets containing 0.6 wt. % of the element Zr, the coercivity of the magnets was basically the same as that of magnets without the element Zr. However, the absence of one of the four elements would lead to a great decrease in the coercivity of magnets. Therefore, these four elements must be added to improve the mechanical properties of the magnets by Zr compounds in this experiment to ensure that the coercivity of the magnets did not decrease.

[0119] In the present invention, rod-like Zr compounds were distributed in the R-rich intergranular phases, and the mechanical properties of the magnets were improved by strengthening the R-rich intergranular phases. However, the distribution of rod-like Zr compounds in intergranular phases would hinder the flow and distribution of intergranular phases during the second aging stage, thus deteriorating the coercivity of the magnets. References suggested that when the element Nd in R-T-B magnets was replaced with a certain amount of the element Pr, the element Pr was mainly distributed in the intergranular phases. The combined addition of low melting point elements Al, Cu and Ga significantly reduced the melting point of grain boundary phase and enhanced the wettability between the intergranular phases and the main phase grains, so as to improve the distribution of intergranular phases and the coercivity of magnets in the second aging stage.