CN107077940A - Sintered magnet based on R-FE-B without heavy rare earth element and preparation method thereof - Google Patents

Abstract

The present invention relates to a kind of sintered magnet based on R Fe B, wherein the sintered magnet is included：Principal phase, it is by the R containing LREE₂Fe₁₄B (R is La, Ce, Nd, Pr, Pm, Sm, Eu or Nb) crystal grain is formed；With the second phase, it is made up of such microstructure, and the rich R phases (R is La, Ce, Nd, Pr, Pm, Sm, Eu or Nb) in the microstructure containing LREE surround the crystal grain, and by the R₂Fe₁₄The crystal boundary of B crystal grain formation optionally includes refractory metal element, two adjacent R₂Fe₁₄Contact rate between B crystal grain is 50% or smaller.

Description

Sintered magnet based on R-FE-B without heavy rare earth elements and its preparation method

technical field

The present invention relates to R-Fe-B sintered magnets free of heavy rare earth elements (HREE), and more particularly, to R-Fe-B sintered magnets manufactured based on: liquid coating technology, wherein even without adding heavy rare earth elements In the case of using a high melting point metal precursor to prepare Nd-Fe-B powder with a controlled microstructure, the improved magnetic properties of the sintered magnet are obtained; the sintering process, the controlled microstructure during the sintering process effectively Inhibition of grain growth to obtain increased coercive force of sintered magnets; and techniques related to controlling the addition of small amounts of high-melting point metals and the selective formation of grain boundaries to minimize the reduction in remanence of sintered magnets. The invention also relates to a method for producing said sintered magnet.

Background technique

The maximum magnetic energy product (BH-max) value of Nd-Fe-B sintered magnet is 29MGOe to 53MGOe, which is higher than other permanent magnets such as Alnico permanent magnet alloy (1MGOe to 7.5MGOe), ferrite (1.1 MGOe to 4.5MGOe) and SmCo ₅ (18MGOe to 33MGOe). BH-max is a parameter representing the magnetic properties of the permanent magnet. Nd-Fe-B sintered magnets are currently considered to be the strongest permanent magnets among existing permanent magnets. Due to such excellent magnetic properties, Nd-Fe-B sintered magnets have been widely used in motors for machine tools and industrial robots, electronic information devices, and small motors for automobiles. Since Nd-Fe-B sintered magnets are recently applied to drive motors of hybrid and electric vehicles, Nd-Fe-B sintered magnets have received much attention worldwide. However, the Curie temperature of Nd-Fe-B sintered magnets is only 350°C, at which temperature Nd-Fe-B sintered magnets begin to lose their magnetic properties, and have the disadvantage of deteriorating their magnetic properties at elevated temperatures . In particular, for driving motors of hybrid and electric vehicles exposed to extreme environmental conditions including temperatures up to 200° C., Nd—Fe—B sintered magnets need to be manufactured considering the deterioration of magnetic properties. In general, deterioration of magnetic properties is closely related to coercive force. According to a report on the magnetic properties of a commercially available Nd-Fe-B sintered magnet, its coercive force is 25 kOe at room temperature, and the coercive force temperature coefficient is -0.5%/°C. Therefore, it is known that commercially available Nd-Fe-B sintered magnets lose about 50% and about 75% of their coercivity when exposed to 100°C and 150°C, respectively. Nd-Fe-B sintered magnets for electric motors of hybrid vehicles need to have high coercivity at room temperature to maintain their coercive force within the operating temperature range of vehicle motors.

The most common method for making Nd-Fe-B sintered magnets with high coercivity involves adding elements with high intrinsic coercivity. The coercive force of Nd-Fe-B sintered magnets can be improved by adding representative heavy rare earth elements such as Dy and Tb. Such heavy rare earth elements form intermetallic compounds such as Dy ₂ Fe ₁₄ B and Tb ₂ Fe ₁₄ B having magnetic anisotropy constants of 150 kOe and 220 kOe (at least twice that of Nd ₂ Fe ₁₄ B (67 kOe)), respectively. That is to say, the addition of heavy rare earth elements is very helpful to improve the coercive force of Nd-Fe-B sintered magnets.

However, the production of heavy rare earth elements is smaller than that of light rare earth elements. In particular, heavy rare earth elements are less abundant in the earth's crust and are at least 10 times more expensive than light rare earth elements. For these reasons, continuous efforts have been made to minimize the addition of heavy rare earth elements.

It is generally believed that reducing the grain size by inhibiting grain growth is the most effective way to increase the coercive force of Nd-Fe-B sintered magnets without adding heavy rare earth elements. Therefore, many studies have been attempted to reduce the size of crystal grains. One example is grain refinement by grain boundary pinning effect, which is induced by the addition of high melting point metals such as Mo, Nb and W to form a second phase at grain boundaries or triple crystal intersections. Due to the low solubility of the added refractory metal in the main phase Nd ₂ Fe ₁₄ B, it forms precipitates such as (Mo,Fe) ₃ B ₂ , Nb—Fe—B and W—Fe—B. It is reported that due to the presence of such precipitates in the second phase (eg, at the grain boundaries), they exhibit a grain boundary pinning effect during sintering to inhibit grain growth. However, in this case, the precipitates present in the Nd ₂ Fe ₁₄ B grain boundaries cause the formation of reverse magnetic domains, and the size of the precipitates on the Nd ₂ Fe ₁₄ B grain boundaries increases as the amount of the additive increases. large, leading to a decrease in the coercive force. The presence of precipitates in the main phase (Nd ₂ Fe ₁₄ B) reduces the relative proportion of the main phase, resulting in a decrease in remanence [Non-Patent Documents 1 and 2].

(Non-Patent Document 1) A. Yan, X. Song, M. Song, X. Wanget, J. Alloy. Compd, 257, 273 (1997).

(Non-Patent Document 2) S. Hirosawa, H. Tomizawa, S. Mino, A. Hamamura, IEEE. Trans. Magn, 26, 1960 (1990).

Contents of the invention

Problem to be solved by the present invention

The present invention has been made in consideration of the above problems, and aims to propose a technique related to the manufacture of sintered magnets by: controlling the amount of the second phase to selectively form the second phase at grain boundaries; controlling the amount of the second phase Size enables fine and uniform distribution of the second phase to maximize grain boundary pinning; and suppresses grain growth through maximized grain boundary pinning for increased coercive force and minimized reduction in remanence change.

Another object of the present invention is to provide a method for manufacturing HREE-free sintered magnets through a simplified process.

way of solving the problem

Therefore, the present inventors attempted to induce the formation of molybdenum on the surface of heavy rare earth element (HREE)-free Nd-Fe-B powder by the following steps: molybdenum pentaethoxide (Mo(OC ₂ H ₅ ) ₅ ) dissolved in absolute alcohol, immersing HREE-free Nd-Fe-B powder in the solution to coat the high melting point metal precursor on the surface of the powder, and make the coated powder Thermal decomposition to remove impurities other than molybdenum contained in the precursor. The resulting powders consisted of HREE-free Nd—Fe—B powder as the core and molybdenum as the shell. The fine second phase can be uniformly distributed across the grain boundaries of the entire sample during sintering. Furthermore, the present inventors tried to suppress the intergrain diffusion of the second phase into the main phase Nd ₂ Fe ₁₄ B to minimize the change in remanence by controlling the amount of added molybdenum.

One aspect of the present invention provides an R-Fe-B (R=La, Ce, Nd, Pr, Pm, Sm, Eu, or Nb) sintered magnet comprising: a main phase composed of R ₂ Fe ₁₄ B grain formation; and a second phase having a microstructure in which an R-rich phase comprising a light rare earth element surrounds the grain and is selectively formed in the R ₂ Fe ₁₄ B grain The formed grain boundaries or triple crystal intersections contain high melting point metal elements, wherein the contact ratio between two adjacent R ₂ Fe ₁₄ B crystal grains is 50% or less.

The second phase is selected from Mo ₂ FeB ₂ and MoFe ₂ .

The average diameter of the R ₂ Fe ₁₄ B crystal grains is 5 nm to 6.5 nm.

The coercive force of the sintered magnet is 10kOe to 20kOe, and the remanence is 1T to 1.7T.

Another aspect of the present invention provides a method for manufacturing an R-Fe-B sintered magnet, comprising:

1) mixing the R-Fe-B powder with a solution of the high melting point metal precursor in absolute alcohol to coat the high melting point metal precursor on the surface of the R-Fe-B powder;

II) drying the R-Fe-B powder coated with the high melting point metal precursor, followed by thermal decomposition to prepare the core-shell raw material powder; and

III) Sintering the raw material powder.

The core-shell raw material powder contains 0.03% by weight to 0.20% by weight of Mo on the surface of the R-Fe-B powder.

The refractory metal precursor is molybdenum pentaethoxide (Mo(OC ₂ H ₅ ) ₅ ).

In step II), thermal decomposition is carried out at ambient pressure and at 750°C to 1000°C.

In step II), thermal decomposition is carried out at a reduced pressure of 10 ⁻³ Torr and a temperature of 250°C to 400°C.

In step III), sintering is carried out at 900°C to 1100°C.

In step III), sintering is performed at a heating rate of 5°C/min to 15°C/min.

The present invention also provides molybdenum pentaethoxide as the refractory metal precursor used in the method.

Invention effect

In the R-Fe-B sintered magnet of the present invention, the refractory metal is formed on the surface of the R-Fe-B raw material powder so that the fine second phase is uniformly distributed at the grain boundaries and triple crystal intersections throughout the sample. Due to this uniform distribution, the microstructure of the sintered magnet can be effectively controlled. Therefore, the sintered magnet of the present invention can overcome the limited physical and magnetic properties of the existing R-Fe-B sintered magnet. Furthermore, the sintered magnet of the present invention has no problems associated with the supply and demand of heavy rare earth elements, and thus can be obtained at a reasonable price.

Description of drawings

FIG. 1 is a schematic diagram showing a method for manufacturing a sintered magnet according to the present invention.

Figure _{2 shows the TGA and DSC results of Mo(OC2H5)5 as a} refractory metal precursor.

Figure 3 shows the XRD patterns of the pressed samples containing Mo and the pressed samples not containing Mo.

Figure 4 shows the surface and cross-sectional scanning electron microscope images of Mo-coated Nd-Fe-B powders, which were measured to study the effect of thermal decomposition on the formation of intermetallic compounds in the powders; points A and B are respectively Shell and core of core-shell raw material powder.

Figure 5 shows the microstructural changes observed by SEM (BSE) and EPMA after sintering of Mo-coated Nd-Fe-B powders.

Figure 6 shows the XRD patterns of Mo-containing sintered magnets and Mo-free sintered magnets, which were measured for accurate analysis of the second phase observed in the SEM and EPMA images.

Figure 7 shows scanning electron microscope (BSE) images and optical microscope (OM) images of the following samples: (a) Dy-free powder (without HREE), (b) Mo-free sintered magnet, (c) Sintered Mo-containing magnets (with the addition of 0.03% by weight of Mo), (d) sintered Mo-containing magnets (with the addition of 0.05% by weight of Mo), and (e) sintered Mo-containing magnets (with the addition of 0.2% by weight of Mo), and using Images of the obtained average grain size and grain size distribution.

FIG. 8 shows changes in coercivity of Mo-free sintered magnets and sintered Mo-containing magnets (0.03 wt%, 0.05 wt%, and 0.2 wt% Mo added).

best practice

The present invention will now be described in more detail.

The term "at the intersection of three crystals" as used herein refers to a region in the sintered magnet where three crystal grains contact each other to form an R-rich phase.

The R-Fe-B sintered magnet has a structure in which a main phase formed of R ₂ Fe ₁₄ B grains is surrounded by an R-rich phase. The magnetic and other properties of a sintered magnet are determined by various parameters such as the size and segregation of grains and the thickness of the R-rich phase. In particular, heavy rare earth metals (Dy or Tb) with high intrinsic magnetic anisotropy fields are mainly used to improve the magnetic properties of sintered magnets. However, the low reserves and localization of heavy rare earth elements lead to supply-demand imbalance and price instability, limiting their use.

Therefore, the present inventors conducted research work to develop a high-performance magnetic material having magnetic properties superior to existing R-Fe-B sintered magnets. As a result, the present inventors paid attention to the fact that when high-melting point metals are simply mixed, the size of crystal grains can be limited but precipitates are formed during production, resulting in reduced coercivity and remanence. The present inventors tried to find a solution to the above-mentioned problems. In particular, the inventors found that when the high-melting-point metal precursor was dissolved in anhydrous alcohol, HREE-free Nd-Fe-B powder was immersed in the solution to coat the high-melting-point metal on the HREE-free Nd-Fe-B powder. -B powder, and when the coated powder is sintered, it is possible to manufacture a sintered magnet in which the second phase containing the high-melting point metal is selectively present at the grain boundary or at the intersection of three crystals. The present invention has been accomplished based on this finding.

The present invention aims to provide a sintered magnet different from the existing R-Fe-B sintered magnets in that the size and microstructure of the main phase formed by _{R2Fe14B crystal} grains are effectively is controlled to obtain improved magnetic properties such as high coercivity and remanence.

An aspect of the present invention provides a sintered magnet comprising: a main phase formed of R ₂ Fe ₁₄ B grains containing a light rare earth element; and a second phase having a light rare earth element contained therein. The R-rich phase surrounds the microstructure of the grains, and contains refractory metal elements at the grain boundaries formed by the R ₂ Fe ₁₄ B grains or at the intersection points of the three crystals.

The light rare earth elements present in the main phase and the light rare earth elements present in the R-rich phase are independent of each other. That is, the two light rare earth elements may be the same as or different from each other.

R is La, Ce, Nd, Pr, Pm, Sm, Eu or Nb. Nd is used as R in the Examples section below.

The contact ratio between two adjacent R ₂ Fe ₁₄ B grains is 50% or less, preferably 23% to 40%. The contact ratio is a parameter that numerically shows that the grains are almost completely segregated by the R-rich phase. A lower contact ratio indicates that the grains are less in contact with each other.

The contact ratio is defined as the ratio of the grain boundary contact area between the same two phases to the total area of the microstructure grain boundaries. That is, the contact rate is defined as the ratio of the grain boundary area in contact with each other to the total grain boundary area, or the ratio of the grain boundaries adjacent to each other to the total grain boundary area [METALLURGICAL TRANSACTIONS A, R.M.GERMAN, Volume 16A, 1985 July, 1247; METALLOGRAPHY, V. Srikanth, G.S. Upadhyaya, Volume 19, November 4, 1986, 437-445; International Journal of Refractory Metals & Hard Materials, V.T.Golovchan, N.V.Litoshenko, 21, 2003, 241-244] . A higher contact ratio indicates that the grains are more in contact with each other. A lower contact ratio indicates that the grains are isolated from each other.

The second phase is Mo ₂ FeB ₂ or MoFe ₂ . The average grain size of the second phase is lower than the submicron level, and it is uniformly distributed at the grain boundaries formed by R ₂ Fe ₁₄ B grains or at the intersection points of the three crystals. The second phase effectively controls the grain size while preventing Mo from dissolving into the grains, resulting in enhanced coercive force and remanence of the sintered magnet.

The presence of a large amount of Nd and a small amount of Nd _1.x Fe ₄ B ₄ in the initial Nd ₂ Fe ₁₄ B powder can be observed by the microstructure and XRD analysis of this initial powder. _Nd and _{Nd1.xFe4B4 react} with Mo to form new intermetallic compounds, where Mo is formed from high melting point metal precursors during the sintering process. Based on the standard Gibbs free energy of formation, Nd ₂ Fe ₁₄ B is more stable than Nd _1.x Fe ₄ B ₄ , which improves the formation of Mo and Nd _1.x Fe ₄ B ₄ on the surface of Nd ₂ Fe ₁₄ B powder. Possibility of reaction to form intermetallic compounds without _reacting with _Nd2Fe14B . Furthermore, considering the phase diagram of binary alloys, Mo and Nd cannot form intermetallic compounds through the reaction at the corresponding temperature, but Mo and Fe can. As a result, the core-shell structured R-Fe-B powders formed by coating with high-melting point metals can form a second phase during the sintering process through the following chemical reactions:

(1)4Mo+Nd _1.x Fe ₄ B ₄ →2Mo ₂ FeB ₂ +2Fe+1.xNd

(2) _x Fe+Mo→MoFe _x

During sintering, _Nd1.xFe4B4 present _on the surface of Nd _- Fe _- B powder can react with Mo to form Mo2FeB2 phase, as described in reaction ( ₁ ). However, as determined by XRD phase analysis of the samples after sintering was complete, the _MoFex phase was formed by the reaction of Mo with Fe present in the Nd _- rich phase or Fe remaining after the formation of the _Mo2FeB2 phase.

The use of Mo-coated core-shell raw material powders in the present invention enables the formation of a second phase during the sintering process. The second phase effectively suppresses grain growth, limiting grain size deviations to 1.5 μm or less.

The formation of the second phase improves the wettability between the R-rich phase and the grains, and allows the R-rich phase to penetrate better into the grain boundaries.

The average diameter of R ₂ Fe ₁₄ B crystal grains is 5 nm to 6.5 nm, which is a level suitable for sintered magnets. If the diameter of the R ₂ Fe ₁₄ B grains exceeds 6.5 nm, the grains are not easily isolated, and thus magnetic exchange coupling occurs between the grains, resulting in low coercive force.

The coercive force of the sintered magnet having the above structure is 10 kOe to 20 kOe, and the remanence is 1 T to 1.7 T, which are at a higher level than the existing sintered magnets.

Another aspect of the present invention provides a method for manufacturing the microstructured sintered magnet, as shown in FIG. 1 .

The production of core-shell structured raw powders is usually achieved by dry coating processes such as physical vapor deposition, chemical vapor deposition or spray coating. On the contrary, the method of the present invention adopts a liquid coating process to prepare a core-shell structure raw material powder with a uniform shell thickness in a faster and simpler manner.

Core-shell raw material powders were prepared using a liquid coating process, in which refractory metals were coated on R-Fe-B powders prepared by strip casting. First, the R-Fe-B powder was immersed in a solution of the high-melting-point metal precursor in absolute alcohol to coat the R-Fe-B powder with the high-melting-point metal precursor.

Then, drying and thermal decomposition are performed to decompose organic compounds from the coated R-Fe-B powder.

Thermal decomposition is preferably carried out at ambient pressure and at 750°C to 1000°C, which is determined to be the optimum condition by the TGA and DSC analysis results in Figure 2 .

Thermal decomposition can be performed at a reduced pressure of 10 ⁻³ Torr and a temperature of 250°C to 400°C.

Most preferably, the refractory metal precursor is molybdenum pentaethoxide (Mo(OC ₂ H ₅ ) ₅ ).

The raw material powder of core-shell structure has been surrounded before sintering, which is effective for the isolation of grains.

As described above, the raw material powder of the core-shell structure contains 0.03% by weight to 0.20% by weight of the Mo shell on the surface of the R-Fe-B powder. The Mo shell content is 0.03% by weight or more, preferably 0.03% to 0.2% by weight. When the Mo shell content is 0.2% by weight, the maximum increase in coercive force can be ensured.

If the Mo shell content is less than the lower limit defined above, it is difficult to effectively limit the size of crystal grains. Furthermore, if the Mo shell content exceeds the upper limit defined above, the refractory metal (Mo) diffuses excessively into the crystal grains, deteriorating the coercive force of the sintered magnet. It is therefore preferred to adjust the Mo shell content to the range defined above.

As described above, the steps of the process for preparing the raw material powder of the core-shell structure are reduced, enabling rapid preparation of the raw material powder of the core-shell structure, and are advantageous in coating efficiency compared with conventional dry processes. In particular, this process avoids the need for additional equipment such as a sputtering system, and is therefore advantageous in terms of cost-effectiveness compared to conventional dry processes.

Finally, the core-shell structure raw material powder is sintered at 900°C to 1100°C to manufacture the desired R-Fe-B sintered magnet.

In particular, when the sintering temperature reaches about 635 °C, the R-rich phase begins to appear in the liquid phase.

When the temperature is further increased, the shell containing the high-melting metal containing the core diffuses along the grain boundaries into the R-rich liquid phase and forms a second phase around the grains. The second phase further isolates the grains.

The shell comprising the refractory metal reacts with the Nd _1.x Fe ₄ B ₄ present on the surface of the Nd ₂ Fe ₁₄ B core powder to form a second phase. The second phase effectively suppresses grain growth during the sintering process so that the grain size can be kept at a low level.

The formation of a second phase at the grain boundaries induces a change in intergranular capillary forces to improve wettability. The improved wettability enables better penetration of the R-rich phase between the grain boundaries, promoting the segregation of grains.

The sintering process suppresses the size of the sintered particles and homogenizes the microstructure. The relative density of the sintered magnet is 99% or higher, the coercive force is 10kOe to 20kOe, and the remanence is 1T to 1.7T, which are higher than the existing sintered magnets. Furthermore, the sintering process ensures high performance of sintered magnets even without using any heavy rare earth elements. High-performance sintered magnets can be used as a substitute for HREE sintered magnets in magnetic materials for engines, generators, and green energy and its application components.

Another aspect of the invention relates to the use of molybdenum pentaethoxide as a refractory metal precursor used in a method for manufacturing R-Fe-B sintered magnets. More specifically, it relates to the use of molybdenum pentaethoxide (Mo(OC ₂ H ₅ ) ₅ ) for improving the coatability of Mo as a refractory metal on the surface of particles such as R-Fe-B powder particles.

When coating particles such as R-Fe-B powder particles with refractory metal (Mo) by liquid coating process, the presence of pentaethoxy molybdenum (Mo(OC ₂ H ₅ ) ₅ ) makes the R-Fe-B A thin, uniform coating (shell) is formed on the powder surface.

best practice

The present invention will be explained in more detail with reference to the following examples. However, these examples should not be construed to limit or define the scope and disclosure of the invention. It is understood that other embodiments of the invention (experimental results of which are not explicitly shown) can be readily practiced by those skilled in the art based on the teachings of the invention including the following examples. It is also to be understood that such modifications and changes are intended to fall within the scope of the appended claims.

Example 1

(1) A sample having a composition of Nd ₁₄ Fe ₈₀ B ₆ (Nd:14, Fe:80, B:6 (atomic %)) was produced. First, raw materials were liquefied at 1600°C and alloy strips were prepared by strip casting. The alloy ribbon was hydrogenated to form microcracks at grain boundaries, jet-milled, and classified into powders with an average particle size (D ₅₀ ) of 5.0 μm. The particle size distribution was 2 μm to 10 μm with a standard deviation of 0.94.

(2) Pentaethoxy Mo (Mo(OC ₂ H ₅ ) ₅ ) was used as a high melting point metal precursor. The refractory metal precursor was dissolved in absolute (ethanol) alcohol. The powder was immersed in this solution, dried under an argon atmosphere, and thermally decomposed at 750 °C for 30 minutes to remove organic compounds. R ₂ Fe ₁₄ B from the powder forms the core, and Mo from the refractory metal precursor forms the shell.

(3) Next, the obtained core-shell structure raw material powder was compacted using a magnetic field press under a static magnetic field of 20 kOe to produce a pressed sample with a size of 20×12×15 mm ³ . The compaction pressure was 1.2 tons, and the relative density of the pressed sample was 48%.

Subsequently, the pressed samples were sintered at 1070° C. for 4 hours in a vacuum furnace maintained at vacuum (≦2.4×10 ⁻⁶ Torr) to fabricate Nd—Fe—B sintered magnets. Under these temperature and time conditions, sintering sufficiently induces a uniform distribution of the Nd-rich liquid phase in the Nd ₂ Fe ₁₄ B grain boundaries.

Figure _{2 shows the TGA and DSC results of Mo(OC2H5)5 as a} refractory metal precursor. As shown in Figure 2, weight changes of the refractory metal precursors were observed at two temperatures of 290 °C and 750 °C. The DSC curve showed that the reaction was exothermic.

These results show the optimal conditions for the thermal decomposition of the powder of Example 1(1) to the core-shell raw material powder of Example 1(2).

Figure 3 shows the XRD patterns of the pressed samples containing Mo and the pressed samples not containing Mo. In both samples, peaks corresponding to Nd ₂ Fe ₁₄ B and Nd-rich phases were found. Peaks corresponding to the Nd _1.x Fe ₄ B ₄ phase were also found. It is known that Nd _1.x Fe ₄ B ₄ is attributed to the presence of relatively little Fe compared to the amount of B when preparing Nd-Fe-B powder and Nd _1.x Fe ₄ B ₄ exists in Nd-Fe-B powder s surface. However, in the Mo-containing pressed samples, a low-intensity peak corresponding to the Mo phase was observed.

Figure 4 shows the surface and cross-sectional scanning electron microscope images of Mo-coated Nd-Fe-B powders. In Fig. 4, points A and B are the shell and the core of the core-shell raw material powder, respectively.

As shown in Figure 4, EDS analysis indicated that only Nd, Fe and O were present within the Nd-Fe-B powder, and Mo as well as Nd, Fe and O were present on the surface of the Nd-Fe-B powder.

These results indicate that the refractory metal element Mo is only partially coated on the surface of the Nd—Fe—B powder.

Figure 5 shows the microstructural changes observed by SEM (BSE) and EPMA after sintering of Mo-coated Nd-Fe-B powders. SEM (BSE) analysis showed that, except for the corresponding Nd ₂ Fe ₁₄ B( Hard magnetic phase) and a bright area corresponding to the Nd-rich phase, a second phase having a contrast different from that of the Nd-rich phase (non-magnetic phase) exists at the grain boundaries. EPMA imaging was performed to analyze the observed elements of the second phase. The second phase was determined by BSE imaging. The results show that a large number of Mo atoms are present in the second phase. The size of the second phase at the three-crystal intersections and grain boundaries is less than 1 μm (submicron), and it is uniformly distributed throughout the sintered sample. The uniform distribution of the second phase is considered to be due to the uniform distribution of Mo element on the Nd-Fe-B surface by the liquid coating process used for powder preparation. The uniform formation of the second phase at grain boundaries and triple-crystal intersections during the sintering process may inhibit grain boundary movement. It was also analyzed that the second phase containing Mo element does not exist in the Nd ₂ Fe ₁₄ B phase. It is believed that the addition of a very small amount of Mo induces the formation of a _second phase, which effectively suppresses the dissolution of Mo in the _Nd2Fe14B phase. Furthermore, microstructural changes showing segregation of the Nd ₂ Fe ₁₄ B phase can be observed because the Nd-rich phase is very continuous. Sintered magnets with microstructured Nd-rich phases can be expected as microstructures capable of effectively controlling coercive force reduction through exchange between ferromagnets in a nucleating coercive force mechanism.

Figure 6 shows the XRD patterns of Mo-containing sintered magnets and Mo-free sintered magnets, which were measured for accurate analysis of the second phase observed in the SEM and EPMA images. The XRD analysis of the sintered samples showed that in the case of the sintered Mo-containing samples, besides the Nd ₂ Fe ₁₄ B phase and the Nd-rich phase, there was also a large-volume Mo ₂ FeB ₂ phase and a small-volume MoFe ₂ phase. A large amount of Nd and a small amount of Nd _1.x Fe ₄ B ₄ phase may be attached on the surface of the initial Nd ₂ Fe ₁₄ B powder and may react with Mo on the powder surface to form an intermetallic compound phase. Based on the standard Gibbs free energy of formation, Nd ₂ Fe ₁₄ B is more stable than Nd _1.x Fe ₄ B ₄ . Furthermore, considering the binary alloy phase diagram, Mo and Nd cannot form compounds, but Mo and Fe can. As a result, there is a possibility that the powder of the core (Nd ₂ Fe ₁₄ B powder)-shell (Mo element) structure can form an intermetallic compound phase (second phase) by the following chemical reaction during the sintering process:

(1) 4Mo+Nd ₂ Fe ₁₄ B→2Mo ₂ FeB ₂ +2Fe+1.xNd

(2)XFe+Mo→MoFe _x

During the sintering process, the Nd _1.x Fe ₄ B ₄ present on the surface of the Nd—Fe—B powder can react with Mo to form the Mo ₂ FeB ₂ phase, as shown in reaction (1). However, a number of possibilities for _MoFex phase formation are contemplated, as determined by XRD phase analysis of samples after sintering is complete. For example, the _MoFex phase can be formed by the reaction of Mo with a small amount of Fe present _on the Nd _- rich phase or Fe remaining after the formation of the Mo2FeB2 phase. The phase observed here is the _MoFe2 phase (one based on Mo-Fe compounds).

Figure 7 shows scanning electron microscope (BSE) images and optical microscope (OM) images of the following samples: (a) Dy-free powder (without HREE), (b) Mo-free sintered magnet, (c) Sintered Mo-containing magnets (0.03 wt% Mo addition), (d) sintered Mo-containing magnets (0.05 wt% Mo addition), and (e) sintered Mo-containing magnets (0.2 wt% Mo addition), and Use the images to obtain the average grain size and grain size distribution. The average grain size and grain size distribution of about 1,000 to 1,100 Nd ₂ Fe ₁₄ B phases were measured at a magnification of 500×. The imaging results showed that the average grain sizes of the sintered samples containing 0.03 wt%, 0.05 wt% and 0.20 wt% Mo were 6.07 ± 0.13 μm, 5.88 ± 0.11 μm and 5.60 ± 0.11 μm, respectively, which were larger than those without Mo ( 7.4±0.22μm) as small as about 1.33μm to 1.8μm. To analyze the grain size distribution, the standard deviation of the measured grains was calculated. For Mo additions of 0.03 wt%, 0.05 wt%, and 0.20 wt%, the standard deviations of the Mo-containing samples were 1.53 μm, 1.42 μm, and 1.3 μm, respectively, while the standard deviation of the Mo-free samples was 2.5 μm. That is, the standard deviation decreases as the amount of Mo added increases. From these results, it can be concluded that as the addition of Mo increases, the grain size becomes uniform. The reason why the grain size decreases even when a small amount of Mo is added is that Mo is uniformly distributed on the surface of the Nd—Fe—B powder by a liquid coating process using an Mo organic compound.

FIG. 8 shows changes in coercive force of sintered magnets without Mo and sintered magnets containing Mo (0.03 wt%, 0.05 wt%, and 0.2 wt% Mo added).

As shown in FIG. 8, the coercive force of the sample containing no Mo was 11.88 kOe (remanence: 1.37 T), while the coercive forces of the samples containing 0.03 wt%, 0.05 wt% and 0.20 wt% Mo were respectively 12.83kOe, 13.1kOe and 13.95kOe. In particular, the coercive force of the sample containing 0.20% by weight of Mo was 2.07 kOe higher than that of the sample containing no Mo. In particular, even when the amount of Mo was increased, the remanence of the samples containing Mo remained unchanged or decreased only slightly (1.35T to 1.37T).

From these results, it can be concluded that the addition of refractory metal (Mo) effectively suppresses grain formation and growth, making the grain size uniform, as shown in FIG. 7 .

Furthermore, the application of an improved liquid coating process using high melting point metal precursors to fabricate the sintered magnets of the present invention enables the fine second phase to be uniformly distributed in the sintered magnets and induces small and uniform grain growth. In particular, adding only a very small amount of high-melting-point metal precursors is also very effective in controlling grain growth.

Taken together, it is believed that the enhanced coercive force is explained by the addition of very small amounts of high melting point metal precursors, which induce the selective formation of the second phase at the grain boundaries to effectively limit the grain size growth during sintering. In particular, the microstructure effectively suppresses the dissolution of Mo in Nd ₂ Fe ₁₄ B, minimizing the reduction in remanence.

Industrial applicability

In the R-Fe-B sintered magnet of the present invention, the refractory metal is formed on the surface of the R-Fe-B raw material powder so that the fine second phase is uniformly distributed at the grain boundaries and the intersections of the three crystals throughout the sample. Due to this uniform distribution, the microstructure of the sintered magnet can be effectively controlled. Therefore, the sintered magnet of the present invention can overcome the limited physical and magnetic properties of the existing R-Fe-B sintered magnet. Furthermore, the sintered magnet of the present invention has no problems associated with the supply and demand of heavy rare earth elements, and thus can be obtained at a reasonable price.

Claims (11)

1. a kind of R-Fe-B (R=La, Ce, Nd, Pr, Pm, Sm, Eu or Nb) sintered magnet, comprising：Principal phase, the principal phase is by wrapping R containing LREE₂Fe₁₄B crystal grain is formed；Mutually there is the rich R phases for wherein including LREE with the second phase, described second The microstructure of the crystal grain is surrounded, and by the R₂Fe₁₄The grain boundaries of B crystal grain formation or three brilliant point of intersection include Gao Rong Point metallic element, two of which adjacent R₂Fe₁₄Contact rate between B crystal grain is 50% or smaller.

2. R-Fe-B sintered magnets according to claim 1, wherein described second is mutually Mo₂FeB₂Or MoFe₂。

3. R-Fe-B sintered magnets according to claim 1, wherein the R₂Fe₁₄The average diameter of B crystal grain be 5nm extremely 6.5nm。

4. R-Fe-B sintered magnets according to claim 1, wherein the coercivity of the sintered magnet be 10kOe extremely 20kOe, remanent magnetism is 1T to 1.7T.

5. a kind of method for manufacturing R-Fe-B sintered magnets, including I) by R-Fe-B powder and refractory metal precursor in nothing Solution in water alcohol is mixed so that the refractory metal precursor is coated on the surface of the R-Fe-B powder, II) dry warp The R-Fe-B powder of the refractory metal precursor coating, then is thermally decomposed to prepare core-shell structure copolymer material powder, and III) The material powder is sintered.

6. method according to claim 5, wherein the refractory metal precursor is five ethyoxyl molybdenum (Mo (OC₂H₅)₅)。

7. method according to claim 5, wherein, in step II) in, the thermal decomposition in environmental pressure and 750 DEG C extremely Carried out at 1000 DEG C.

8. method according to claim 5, wherein, in step II) in, the thermal decomposition is 10^-3The decompression of support and 250 DEG C To carrying out at a temperature of 400 DEG C.

9. method according to claim 5, wherein, in step III) in, it is described to be sintered at 900 DEG C to 1100 DEG C OK.

10. method according to claim 5, wherein, in step III) in, the sintering is with 5 DEG C/min to 15 DEG C/minute The rate of heat addition of clock is carried out.

11. five ethyoxyl molybdenums are used as the refractory metal precursor used in method according to claim 5.