Patent Application
20160293307
This non-provisional application claims priority under 35 U.S.C. §119(a) on Patent Application Nos. 2015-072343 and 2016-025548 filed in Japan on Mar. 31, 2015 and Feb. 15, 2016, respectively, the entire contents of which are hereby incorporated by reference.
This invention relates to an R—Fe—B base sintered magnet having a high coercivity and a method for preparing the same.
While Nd—Fe—B sintered magnets, referred to as Nd magnets, hereinafter, are regarded as the functional material necessary for energy saving and performance improvement, their application range and production volume are expanding every year. Since many applications are used in high temperature, the Nd magnets are required to have not only a high remanence but also a high coercivity. On the other hand, since the coercivity of Nd magnets are easy to decrease significantly at a elevated temperature, the coercivity at room temperature must be increased enough to maintain a certain coercivity at a working temperature.
As the means for increasing the coercivity of Nd magnets, it is effective to substitute Dy or Tb for part of Nd in Nd2Fe14B compound as main phase. For these elements, there are short resource reserves in the world, the commercial mining areas in operation are limited, and geopolitical risks are involved. These factors indicate the risk that the price is unstable or largely fluctuates. Under the circumstances, the development for a new process and a new composition of R—Fe—B magnets with a high coercivity, which include minimizing the content of Dy and Tb, is required.
From this standpoint, several methods are already proposed. Patent Document 1 discloses an R—Fe—B base sintered magnet having a composition of 12-17 at % of R (wherein R stands for at least two of yttrium and rare earth elements and essentially contains Nd and Pr), 0.1-3 at % of Si, 5-5.9 at % of B, 0-10 at % of Co, and the balance of Fe (with the proviso that up to 3 at % of Fe may be substituted by at least one element selected from among Al, Ti, V, Cr, Mn, Ni, Cu, Zn, Ga, Ge, Zr, Nb, Mo, In, Sn, Sb, Hf, Ta, W, Pt, Au, Hg, Pb, and Bi), containing a R2(Fe,(Co),Si)14B intermetallic compound as main phase, and exhibiting a coercivity of at least 10 kOe. Further, the magnet is free of a B-rich phase and contains at least 1 vol % based on the entire magnet of an R—Fe(Co)—Si phase consisting essentially of 25-35 at % of R, 2-8 at % of Si, up to 8 at % of Co, and the balance of Fe. During sintering or post-sintering heat treatment, the sintered magnet is cooled at a rate of 0.1 to 5° C./min at least in a temperature range from 700° C. to 500° C., or cooled in multiple stages including holding at a certain temperature for at least 30 minutes on the way of cooling, for thereby generating the R—Fe(Co)—Si phase in grain boundary.
Patent Document 2 discloses a Nd—Fe—B alloy with a low boron content, a sintered magnet prepared by the alloys, and their process. In the sintering process, the magnet is quenched after sintering below 300° C., and an average cooling rate down to 800° C. is ΔT1/Δt1<5K/min.
Patent Document 3 discloses an R-T-B magnet comprising R2Fe14B main phase and some grain boundary phases. One of grain boundary phase is R-rich phase with more R than the main phase and another is Transition Metal-rich phase with a lower rare earth and a higher transition metal concentration than that of main phase. The R-T-B rare earth sintered magnet is prepared by sintering at 800 to 1,200° C. and heat-treating at 400 to 800° C.
Patent Document 4 discloses an R-T-B rare earth sintered magnet comprising a grain boundary phase containing an R-rich phase having a total atomic concentration of rare earth elements of at least 70 at % and a ferromagnetic transition metal-rich phase having a total atomic concentration of rare earth elements of 25 to 35 at %, wherein an area proportion of the transition metal-rich phase is at least 40% of the grain boundary phase. The green body of magnet alloy powders is sintered at 800 to 1,200° C., and then heat-treated with multiple steps. First heat-treatment is in the range of 650 to 900° C., then sintered magnet is cooled down to 200° C. or below, and second heat-treatment is in range of at 450 to 600° C.
Patent Document 5 discloses an R-T-B rare earth sintered magnet comprising a main phase of R2Fe14B and a grain boundary phase containing more R than that of the main phase, wherein easy axis of magnetization of R2Fe14B compound is in parallel to the c-axis, the shape of the crystal grain of R2Fe14B phase is elliptical shape elongated in a perpendicular direction to the c-axis, and the grain boundary phase contains an R-rich phase having a total atomic concentration of rare earth elements of at least 70 at % and a transition metal-rich phase having a total atomic concentration of rare earth elements of 25 to 35 at %. It is also described that magnet are sintered at 800 to 1,200° C. and subsequent heat treatment at 400 to 800° C. in an argon atmosphere.
Patent Document 6 discloses a rare earth magnet comprising R2T14B main phase and an intergranular grain boundary phase, wherein the intergranular grain boundary phase has a thickness of 5 nm to 500 nm and the magnetism of the phase is not ferromagnetism. It is described that the intergranular grain boundary phase is formed from a non-ferromagnetic compound due to add element M such as Al, Ge, Si, Sn or Ga, though this phase contains the transition metal elements. Furthermore by adding Cu to the magnet, a crystalline phase with a La6Co11Ga3-type crystal structure can be uniformly and widely formed as the intergranular grain boundary phase, and a thin R—Cu layer may be formed at the interface between the La6Co11Ga3-type grain boundary phase and the R2T14B main phase crystal grains. As a result, the interface of the main phase is passivated, a lattice distortion of main phase can be suppressed, and nucleation of the magnetic reversal domain can be inhibited. The method of preparing the magnet involves post-sintering heat treatment at a temperature in the range of 500 to 900° C., and cooling at the rate of least 100° C./min, especially at least 300° C./min.
Patent Document 7 and 8 disclose an R-T-B sintered magnet comprising a main phase of Nd2Fe14B compound, an intergranular grain boundary which is enclosed between two main phase grains and which has a thickness of 5 nm to 30 nm, and a grain boundary triple junction which is the phase surrounded by three or more main phase grains.
Patent Document 1: JP 3997413 (U.S. Pat. No. 7,090,730, EP 1420418)
Patent Document 2: JP-A 2003-510467 (EP 1214720)
Patent Document 3: JP 5572673 (US 20140132377)
Patent Document 4: JP-A 2014-132628
Patent Document 5: JP-A 2014-146788 (US 20140191831)
Patent Document 6: JP-A 2014-209546 (US 20140290803)
Patent Document 7: WO 2014/157448
Patent Document 8: WO 2014/157451
However, there exists a need for an R—Fe—B sintered magnet which exhibits a high coercivity despite a minimal content of Dy, Tb and Ho.
An object of the invention is to provide an R—Fe—B sintered magnet exhibiting a high coercivity, and a method for preparing the same.
The inventors have found that a desired R—Fe—B base sintered magnet can be prepared by a method comprising the steps of shaping an alloy powder (consisting essentially of 12 to 17 at % of R, 0.1 to 3 at % of M1, 0.05 to 0.5 at % of M2, 4.8+2×m to 5.9+2×m at % of B, up to 10 at % of Co, and the balance of Fe) into a green compact, sintering the green compact, cooling the sintered compact to room temperature, machining the sintered compact into the shape near the desired end product shape, placing a powder of HR-containing compounds or intermetallic compounds (HR stands for at least one element selected from Tb, Dy and Ho) on the surface of the sintered magnet, heating the powder-coated magnet in vacuum at 700 to 1,100° C. for HR to permeate through the grain boundaries and to diffuse among the sintered magnet, cooling the sintered magnet to a temperature of 400° C. or below at a rate of 5 to 100° C./min, and aging treatment including exposing at a temperature in the range of 400 to 600° C. which temperature is lower than the peritectic temperature of (R,HR)—Fe(Co)-M1 phase so as to form the R—Fe(Co)-M1 phase at a grain boundary, and cooling to a temperature of 200° C. or below.
The magnet contains a R2(Fe,(Co))14B intermetallic compound as main phase and a M2 boride phase at a grain boundary triple junction, but not including R1.1Fe4B4 compound phase, has a core/shell structure that the main phase is covered with HR-rich layer composed of (R,HR)2(Fe,(Co))14B, wherein HR is at least one element selected from Tb, Dy and Ho, the thickness of HR-rich layer is in range of 0.01 to 1.0 μm and moreover the outside of HR-rich layer is covered with (R,HR)—Fe(Co)-M1 phase, wherein at least 50% of the main phase with HR-rich layer is covered with the (R,HR)—Fe(Co)-M1 phase, and a width of the intergranular grain boundary phase is at least 10 nm and at least 50 nm on the average.
The sintered magnet exhibits a coercivity of at least kOe. Continuing experiments to establish appropriate processing conditions and an optimum magnet composition, the inventors have completed the invention.
It is noted that Patent Document 1 recites a low cooling rate after sintering. Even if R—Fe(Co)—Si grain boundary phase forms a grain boundary triple junction, in fact, the R—Fe(Co)—Si grain boundary phase does not enough cover the main phase or form a intergranular grain boundary phase un-continuously. Because of same reason, Patent Document 2 fails to establish the core/shell structure that the main phase is covered with the R—Fe(Co)-M1 grain boundary phase. Patent Document 3 does not refer to the cooling rate after sintering and post-sintering heat treatment, and it does not descript that an intergranular grain boundary phase is formed. The magnet of Patent Document 4 has a grain boundary phase containing R-rich phase and a ferromagnetic transition metal-rich phase with 25 to 35 at % of R, whereas the R—Fe(Co)-M1 phase of the inventive magnet is not a ferromagnetic phase but an anti-ferromagnetic phase. The post-sintering heat treatment in Patent Document 4 is carried out at the temperature below the peritectic temperature of R—Fe(Co)-M1 phase, whereas the post-sintering heat treatment in the invention is carried out at the temperature above the peritectic temperature of R—Fe(Co)-M1 phase.
Patent Document 5 describes that post-sintering heat treatment is carried out at 400 to 800° C. in an argon atmosphere, but it does not refer to the cooling rate. The description of the structure suggests the lack of the core/shell structure that the main phase is covered with the R—Fe(Co)-M1 phase. In Patent Document 6, it is described that the cooling rate of post-sintering heat treatment is preferably at least 100° C./min, especially at least 300° C./min. The sintered magnet above obtained contains crystalline R6T13M1 phase and amorphous or nano-crystalline R-Cu phase. In this invention, R—Fe(Co)-M1 phase in the sintered magnet shows amorphous or nano-crystalline.
The Patent Document 7 provides the magnet contain the Nd2Fe14B main phase, an intergralunar grain boundary and a grain boundary triple junction. In addition, the thickness of the intergranular grain boundary is in range of 5 nm to nm. However the thickness of the intergranular grain boundary phase is too small to achieve a sufficient improvement in the coercivity. Patent Document 8 describes in Example section substantially the same method for preparing sintered magnet as Patent Document 7, suggesting that the thickness (phase width) of the intergranular grain boundary phase is small.
In one aspect, the invention provides an R—Fe—B base sintered magnet of a composition consisting essentially of 12 to 17 at % of R which is at least two of yttrium and rare earth elements and essentially contains Nd and Pr, 0.1 to 3 at % of M1 which is at least one element selected from the group consisting of Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi, 0.05 to 0.5 at % of M2 which is at least one element selected from the group consisting of Ti, V, Cr, Zr, Nb, Mo, Hf, Ta, and W, 4.8+2×m to 5.9+2×m at % of B wherein m stands for atomic concentration of M2, up to 10 at % of Co, up to 0.5 at % of carbon, up to 1.5 at % of oxygen, up to 0.5 at % of nitrogen, and the balance of Fe, containing R2(Fe,(Co))14B intermetallic compound as a main phase, and having a coercivity of at least 10 kOe at room temperature. The magnet contains a M2 boride phase at a grain boundary triple junction, but not including R1.1Fe4B4 compound phase, has a core/shell structure that the main phase is covered with HR-rich layer composed of (R,HR)2(Fe,(Co))14B, wherein HR is at least one element selected from Tb, Dy and Ho, the thickness of HR-rich layer is in range of 0.01 to 1.0 μm, and moreover the outside of HR-rich layer is covered with grain boundary phases comprising an amorphous and/or sub-10 nm nanocrystalline (R,HR)—Fe(Co)-M1 phase consisting essentially of 25 to 35 at % of (R,HR), with the proviso that R and HR are as defined above and HR is up to 30 at % of R+HR, 2 to 8 at % of M1, up to 8 at % of Co, and the balance of Fe, or the (R,HR)—Fe(Co)-M1 phase and a crystalline phase or a sub-10 nm nanocrystalline and amorphous (R,HR)-M1 phase having at least 50 at % of R, wherein a surface area coverage of the (R,HR)—Fe(Co)-M1 phase on the main phase with HR-rich layer is at least 50%, and the width of the intergranular grain boundary phase is at least nm and at least 50 nm on the average.
Preferably in the (R,HR)—Fe(Co)-M1 phase, M1 consists of 0.5 to 50 at % of Si and the balance of at least one element selected from the group consisting of Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi; M1 consists of 1.0 to 80 at % of Ga and the balance of at least one element selected from the group consisting of Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi; or M1 consists of 0.5 to 50 at % of Al and the balance of at least one element selected from the group consisting of Si, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi.
In a preferred embodiment, the total content of Dy, Tb and Ho is up to 5.5 at %, more preferably up to 2.5 at %.
In another aspect, the invention provides a method for preparing the R—Fe—B base sintered magnet defined above, comprising the steps of:
shaping an alloy powder into a green compact, the alloy powder being obtained by finely pulverizing an alloy consisting essentially of 12 to 17 at % of R which is at least two of yttrium and rare earth elements and essentially contains Nd and Pr, 0.1 to 3 at % of M1 which is at least one element selected from the group consisting of Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi, 0.05 to 0.5 at % of M2 which is at least one element selected from the group consisting of Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W, 4.8+2×m to 5.9+2×m at % of B wherein m stands for atomic concentration of M2, up to 10 at % of Co, and the balance of Fe, sintering the green compact at a temperature of 1,000 to 1,150° C.,
cooling the sintered compact to room temperature,
machining the sintered compact into the shape near the desired end product shape,
placing a powder of HR-containing compounds or intermetallic compounds (HR stands for at least one element selected from Tb, Dy and Ho) on the surface of the sintered magnet,
heating the powder-coated magnet in vacuum at 700 to 1,100° C. for HR to permeate through the grain boundaries and to diffuse among the sintered magnet,
cooling the magnet body to a temperature of 400° C. or below at a rate of 5 to 100° C./min, and
aging treatment including exposing at a temperature in the range of 400 to 600° C. which temperature is lower than the peritectic temperature of (R,HR)—Fe(Co)-M1 phase so as to form the (R,HR)—Fe(Co)-M1 phase at a grain boundary, and cooling to a temperature of 200° C. or below.
In a preferred embodiment, the alloy contains Dy, Tb and Ho in a total amount of up to 5.0 at %. In a preferred embodiment, the magnet contains up to 0.5 at % of HR which has been diffused into the magnet as a result of the grain boundary diffusion process. Accordingly, the magnet preferably contains Dy, Tb and Ho in a total amount of up to 5.5 at %.
The R—Fe—B base sintered magnet of the invention exhibits a coercivity of at least 10 kOe despite a low content of Dy, Tb and Ho.
First, the composition of the R—Fe—B sintered magnet is described. The magnet has a composition (expressed in atomic percent) consisting essentially of 12 to 17 at %, preferably 13 to 16 at % of R, 0.1 to 3 at %, preferably 0.5 to 2.5 at % of M1, 0.05 to 0.5 at % of M2, 4.8+2×m to 5.9+2×m at % of B wherein m stands for atomic concentration of M2, up to 10 at % of Co, up to 0.5 at % of carbon, up to 1.5 at % of oxygen, up to 0.5 at % of nitrogen, and the balance of Fe.
Herein, R is at least two of yttrium and rare earth elements and essentially contains neodymium (Nd) and praseodymium (Pr). Preferably Nd and Pr in total account for 80 to 100 at % of R. When the content of R in the sintered magnet is less than 12 at %, the coercivity of the magnet extremely decreases. When the content of R is more than 17 at %, the remanence (residual magnetic flux density, Br) of the magnet extremely decreases. Notably the total content of Dy, Tb and Ho is preferably up to 5.5 at %, more preferably up to 4.5 at %, and even more preferably up to 2.5 at %, based on the magnet composition. When Dy, Tb or Ho is incorporated (or diffused) into the magnet via grain boundary diffusion, the amount of the diffused element is preferably up to 0.5 at %, more preferably 0.05 to 0.3 at %.
M1 is at least one element selected from the group consisting of Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi. When the content of M1 is less than 0.1 at %, the R—Fe(Co)-M1 grain boundary phase is present in an insufficient proportion to improve coercivity. When the content of M1 is more than 3 at %, the squareness of the magnet get worse and the remanence of the magnet decreases significantly. The content of M1 is preferably 0.1 to 3 at %.
An element M2 capable of forming a stable boride is added for the purpose of inhibiting abnormal grain growth during sintering. M2 is at least one element selected from the group consisting of Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W. M2 is desirably added in an amount of 0.05 to 0.5 at %, which enables sintering at a relatively high temperature, leading to improvements in squareness and magnetic properties.
In particular, the upper limit of B is crucial. If the boron (B) content exceeds (5.9+2×m) at % wherein m stands for atomic concentration of M2, the R—Fe(Co)-M1 phase is not formed in grain boundary, but an R1.1Fe4B4 compound phase, which is so-called B-rich phase, is formed. As long as the inventors' investigation is concerned, when the B-rich phase is present in the magnet, the coercivity of the magnet cannot be enhanced enough. If the B content is less than (4.8+2×m) at %, the percent volume of the main phase is reduced so that magnetic properties of the magnet become worse. For this reason, the B content is better to be (4.8+2×m) to (5.9+2×m) at %, preferably (4.9+2×m) to (5.7+2×m) at %.
The addition of Cobalt (Co) to the magnet is optional. For the purpose of improving Curie temperature and corrosion resistance, Co may substitute for up to 10 at %, preferably up to 5 at % of Fe. Co substitution in excess of 10 at % is undesirable because of a substantial loss of the coercivity of the magnet.
For the inventive magnet, the contents of oxygen, carbon and nitrogen are desirably as low as possible. In the production process of the magnet, contaminations of such elements cannot be avoided completely. An oxygen content of up to 1.5 at %, especially up to 1.2 at %, a carbon content of up to 0.5 at %, especially up to 0.4 at %, and a nitrogen content of up to 0.5 at %, especially up to 0.3 at % are permissible. The inclusion of up to 0.1 at % of other elements such as H, F, Mg, P, S, Cl and Ca as the impurity is permissible, and the content thereof is desirably as low as possible.
The balance is iron (Fe). The Fe content is preferably 70 to 80 at %, more preferably 75 to 80 at %.
An average grain size of the magnet is up to 6 μm, preferably 1.5 to 5.5 μm, and more preferably 2.0 to 5.0 μm, and an orientation of the c-axis of R2Fe14B grains, which is an easy axis of magnetization, preferably is at least 98%. The average grain size is measured as follows. First, a cross-section of sintered magnet is polished, immersed into an etchant such as vilella solution (mixture of glycerol:nitric acid:hydrochloric acid=3:1:2) for selectively etching the grain boundary phase, and observed under a laser microscope. On analysis of the image, the cross-sectional area of individual grains is determined, from which the diameter of an equivalent circle is computed. Based on the data of area fraction of each grain size, the average grain size is determined. The average grain size is the average of about 2,000 grain sizes at the different 20 images. The average grain size of the sintered body is controlled by reducing the average particle size of the fine powder during pulverizing.
The microstructure of the magnet contains R2(Fe,(Co))14B phase as a main phase, and (R,HR)—Fe(Co)-M1 phase and (R,HR)-M1 phase as a grain boundary phase. The main phase comprises a HR-rich layer forming at outside of main phase. A thickness of HR-rich layer is up to 1 μm, preferably 0.01 to 1 μm, and more preferably 0.01 to 0.5 μm, and a composition of HR-rich layer is (R,HR)2(Fe,(Co))14B wherein HR is at least one element selected from Tb, Dy and Ho. At the grain boundary phase, (R,HR)—Fe(Co)-M1 phase is formed on the outside of the HR-rich layer to cover the main phase, and which accounts for preferably at least 1% by volume. If the (R,HR)—Fe(Co)-M1 grain boundary phase is less than 1 vol %, a enough high coercivity cannot be obtained. The (R,HR)—Fe(Co)-M1 grain boundary phase is desirably present in a proportion of 1 to 20% by volume, more desirably 1 to 10% by volume. If the (R,HR)—Fe(Co)-M1 grain boundary phase is more than 20 vol %, there may be accompanied a substantial loss of remanence. Herein, the main phase is preferably free of a solid solution of an element other than the above-identified elements. Also R-M1 phase may coexist. Notably precipitation of (R,HR)2(Fe(Co))17 phase is not confirmed. Also the magnet contains M2 boride phase at the grain boundary triple junction, but not R1.1Fe4B4 compound phase. R-rich phase, and phases formed from inevitable elements included in the production process of the magnet such as R oxide, R nitride, R halide and R acid halide may be contained.
The (R,HR)—Fe(Co)-M1 grain boundary phase is a compound containing Fe or Fe and Co, and considered as an intermetallic compound phase having a crystal structure of space group I4/mcm, for example, R6Fe13Ga1. On quantitative analysis by an analytic technique such as electron probe microanalyzer (EPMA), this phase consists of 25 to 35 at % of R, 2 to 8 at % of M1, 0 to 8 at % of Co, and the balance of Fe, the range being inclusive of measurement errors. A Co-free magnet composition may be contemplated, and in this case, as a matter of course, neither the main phase nor the (R,HR)—Fe(Co)-M1 grain boundary phase contains Co. The (R,HR)—Fe(Co)-M1 grain boundary phase is distributed around main phases such that neighboring main phases are magnetically divided, leading to an enhancement in the coercivity.
In the (R,HR)—Fe(Co)-M1 phase, HR substitutes at R site. The content of HR is preferably up to 30 at % of the total content of rare earth elements (R+HR). In general, R—Fe(Co)-M1 phase forms a stable compound phase with a light rare earth element such as La, Pr or Nd. When heavy rare earth elements such as Dy, Tb or Ho substitute for parts of the rare earth elements, a stable phase is yet formed as long as the substitution is up to 30 at %. If the substitution exceeds 30 at %, undesirably a ferromagnetic phase such as (R,HR)1Fe3 phase forms during aging treatment so as to degrade the coercivity and the squareness.
In the (R,HR)—Fe(Co)-M1 phase, it is preferred that M1 consist of 0.5 to 50 at % (based on M1) of Si and the balance of at least one element selected from the group consisting of Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi; 1.0 to 80 at % (based on M1) of Ga and the balance of at least one element selected from the group consisting of Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi; or 0.5 to 50 at % (based on M1) of Al and the balance of at least one element selected from the group consisting of Si, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi. These elements can form stable intermetallic compounds such as R6Fe13Ga1 and R6Fe13Si1 as mentioned above, and are capable of relative substitution at M1 site. Multiple additions of such elements at M1 site does not bring a significant difference in magnetic properties, but in practice, achieves stabilization of magnet quality by reducing the variation of magnetic properties and a cost reduction by reducing the amount of expensive elements.
The width of the (R,HR)—Fe(Co)-M1 phase in intergranular grain boundary is preferably at least 10 nm, more preferably to 500 nm, even more preferably 20 to 300 nm. If the width of the (R,HR)—Fe(Co)-M1 is less than 10 nm, a coercivity enhancement effect due to magnetic decoupling is not obtainable. Also preferably the width of the (R,HR)—Fe(Co)-M1 grain boundary phase is at least 50 nm on an average, more preferably 50 to 300 nm, and even more preferably 50 to 200 nm.
The (R,HR)—Fe(Co)-M1 phase intervenes between neighboring R2Fe14B main phases with the HR-rich layer on the outside as intergranular grain boundary phase, and is distributed around main phase so as to cover the main phase, that is, forms a core/shell structure with the main phase. A ratio of surface area coverage of the (R,HR)—Fe(Co)-M1 phase relative to the main phase is at least 50%, preferably at least 60%, and more preferably at least 70%, and the (R,HR)—Fe(Co)-M1 phase may even cover overall the main phase. The balance of the intergranular grain boundary phase around the main phase is (R,HR)-M1 phase containing at least 50% of the sum of R and HR.
The crystal structure of the (R,HR)—Fe(Co)-M1 phase is amorphous, nano-crystalline or nano-crystalline including amorphous while the crystal structure of the (R,HR)-M1 phase is crystalline or nano-crystalline including amorphous. Preferably nano-crystalline grains have a size of up to 10 nm. As crystallization of the (R,HR)—Fe(Co)-M1 phase proceeds, the (R,HR)—Fe(Co)-M1 phase agglomerates at the grain boundary triple junction, and the width of the intergranular grain boundary phase becomes thinner and discontinuous, as a result the coercivity of the magnet decrease significantly. Also as crystallization of the (R,HR)—Fe(Co)-M1 phase proceeds, R-rich phase may form at the interface between the HR-rich layer covered on the main phase and the grain boundary phase as the by-product of peritectic reaction, but the formation of the R-rich phase itself does not contribute to a substantial improvement in the coercivity.
Now the method for preparing an R—Fe—B base sintered magnet having the above-defined structure is described. The method generally involves grinding and milling of a mother alloy, pulverizing a coarse powder, compaction into a green body applying an external magnetic field, and sintering.
The mother alloy is prepared by melting raw metals or alloys in vacuum or an inert gas atmosphere, preferably argon atmosphere, and casting the melt into a flat mold or book mold or strip casting. If primary crystal of α-Fe is left in the cast alloy, the alloy may be heat-treated at 700 to 1,200° C. for at least one hour in vacuum or in an Ar atmosphere to homogenize the microstructure and to erase α-Fe phases.
The cast alloy is crushed or coarsely grinded to a size of typically 0.05 to 3 mm, especially 0.05 to 1.5 mm. The crushing step generally uses a Brown mill or hydrogen decrepitation. For the alloy prepared by strip casting, hydrogen decrepitation is preferred. The coarse powder is then pulverized on a jet mill by a high-pressure nitrogen gas, for example, into a fine particle powder with a particle size of typically 0.2 to 30 μm, especially 0.5 to 20 μm on an average. If desired, a lubricant or other additives may be added in any of crushing, milling and pulverizing processes.
Binary alloy method is also applicable to the preparation of the magnet alloy power. In this method, a mother alloy with a composition of approximate to the R2-T14-B1 and a sintering aid alloy with R-rich composition are prepared respectively. The alloy is milled into the coarse powder independently, and then mixture of alloy powder of mother alloy and sintering aid is pulverized as well as above mentioned. To prepare the sintering aid alloy, not only the casting technique mentioned above, but also the melt span technique may be applied.
The composition of the alloy is essentially 12 to 17 at % of R which is at least two of yttrium and rare earth elements and essentially contains Nd and Pr, 0.1 to 3 at % of M1 which is at least one element selected from the group consisting of Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb, and Bi, 0.05 to 0.5 at % of M2 which is at least one element selected from the group consisting of Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W, 4.8+2×m to 5.9+2×m at % of B wherein m stands for atomic concentration of M2, up to at % of Co, and the balance of Fe.
The fine powder above obtained is compacted under an external magnetic field by a compression molding machine. The green compact is then sintered in a furnace in vacuum or in an inert gas atmosphere typically at a temperature of 900 to 1,250° C., preferably 1,000 to 1,150° C. for 0.5 to 5 hours.
In the practice of the invention, the HR-rich layer composed of (R,HR)2(Fe,(Co))14B enclosing the main phase of the magnet may be formed by a grain boundary diffusion process. In this case, the sintered compact is machined into a magnet body of desired shape approximate to the end product shape, and then HR element in the powder enclosure is introduced from the magnet body surface into the bulk along the grain boundary phase.
The grain boundary diffusion process of introducing HR element in the magnet body from the surface into the bulk along the grain boundary phase may be (1) a process of placing powder of HR-containing compounds or intermetallic compounds on the surface of the magnet body and heat treating in vacuum or inert gas atmosphere (e.g., dip coating process), (2) a process of forming a thin film of HR-containing compounds or intermetallic compounds on the surface of the magnet body in high vacuum atmosphere and heat treating in vacuum or inert gas atmosphere (e.g., sputtering process), or (3) a process of heating HR element in a high-vacuum atmosphere to create a HR-containing vapor phase, and supplying and diffusing the HR element into the magnet body via the vapor phase (e.g., vapor diffusion process).
Suitable HR-containing compounds or intermetallic compounds include metals, oxides, halides, oxi-halides, hydroxides, carbides, carbonates, nitrides, hydrides, borides of HR, and their mixture, and intermetallic compounds of HR and transition metals such as Fe, Co and Ni wherein part of the transition metal may be substituted by at least one element selected from among Si, Al, Ti, V, Cr, Mn, Cu, Zn, Ga, Ge, Pd, Ag, Cd, Zr, Nb, Mo, In, Sn, Sb, Hf, Ta, W, Pt, Au, Hg, Pb, and Bi.
Preferably the HR-rich layer composed of (R,HR)2(Fe,(Co))14B has a thickness of 10 nm to 1 μm. If the thickness of HR-rich layer is less than 10 nm, any coercivity enhancement effect undesirably is not be obtained. If the thickness of a HR-rich layer is more than 10 μm, the remanence is decreased significantly. The thickness of the HR-rich layer may be controlled by adjusting the amount of HR element added or the amount of HR element diffused into the magnet bulk, or the temperature and time of sintering step, or the temperature and time of grain boundary diffusion treatment.
In the HR-rich layer, HR substitutes at R site. The content of HR is preferably up to 30 at % of the total content of rare earth elements (R+HR). If the HR content exceeds 30 at %, undesirably a ferromagnetic phase such as (R,HR)1Fe3 phase forms during aging treatment, to degrade the coercivity and the squareness.
In order to form the grain boundary phase composed of (R,HR)—Fe(Co)-M1 phase and (R,HR)-M1 phase, the compact as sintered is cooled to a temperature of 400° C. or below, especially 300° C. or below, typically room temperature. The cooling rate is preferably 5 to 100° C./min, more preferably 5 to 50° C./min, though not limited thereto. After sintering, the sintered compact is heated at a temperature in the range of 700 to 1,100° C. which temperature is exceeding peritectic temperature of R—Fe(Co)-M1 phase. (It is called post-sintering heat treatment.) The heating rate is preferably 1 to 20° C./min, more preferably 2 to 10° C./min, though not limited thereto. The peritectic temperature depends on the additive element M1. For example, the peritectic temperature is 640° C. at M1=Cu, 750 to 820° C. at M1=Al, 850° C. at M1=Ga, 890° C. at M1=Si, and 1,080° C. at M1=Sn. The holding time at the temperature is preferably at least 1 hour, more preferably 1 to 10 hours, and even more preferably 1 to 5 hours. The heat treatment atmosphere is preferably vacuum or an inert gas atmosphere such as Ar gas.
This post-sintering heat treatment can combine with the grain boundary diffusion treatment. In this case, the sintered compact may be machined nearly into a body of desired end product shape, for example, by cutting and grinding, and then powder of HR-containing compounds or intermetallic compounds are placed on the surface of the sintered compact obtained by the above method. The sintered magnet body which is enclosed in the HR-containing compound powder, is heat treated in vacuum at a temperature of 700 to 1,100° C. for 1 to 50 hours as the grain boundary diffusion treatment. After the heat treatment, the magnet body is cooled to a temperature of 400° C. or below, preferably 300° C. or below. The cooling rate down to 400° C. or below is 5 to 100° C./min, preferably 5 to 50° C./min, and more preferably 5 to 20° C./min. If the cooling rate is less than 5° C./min, then (R,HR)—Fe(Co)-M1 phase segregates at grain boundary triple junction, and magnetic properties are degraded substantially. A cooling rate of more than 100° C./min is effective for inhibiting precipitation of (R,HR)—Fe(Co)-M1 phase during the cooling step, but the dispersion of (R,HR)-M1 phase in the microstructure is insufficient. As a result, squareness of the sintered magnet becomes worse.
The aging treatment is performed after post-sintering heat treatment. The aging treatment is desirably carried out at a temperature of 400 to 600° C., more preferably 400 to 550° C., and even more preferably 450 to 550° C., for 0.5 to 50 hours, more preferably 0.5 to 20 hours, and even more preferably 1 to 20 hours, in vacuum or an inert gas atmosphere such as Ar gas. The temperature is lower than the peritectic temperature of (R,HR)—Fe(Co)-M1 phase so as to form the (R,HR)—Fe(Co)-M1 phase at a grain boundary. If the aging temperature is blow 400° C., a reaction rate of forming (R,HR)—Fe(Co)-M1 phase is too slow. If the aging temperature is above 600° C., the reaction rate to form (R,HR)—Fe(Co)-M1 phase increases significantly so that the (R,HR)—Fe(Co)-M1 grain boundary phase segregates at the grain boundary triple junction, and magnetic properties are degraded substantially. The heating rate to a temperature in the range of 400 to 600° C. is preferably 1 to 20° C./min, more preferably 2 to 10° C./min, though not limited thereto.
Examples are given below for further illustrating the invention although the invention is not limited thereto.
The alloy was prepared specifically by using rare earth metals (Neodymium or Didymium), electrolytic iron, Co, ferro-boron and other metals and alloys, weighing them with a designated composition, melting at high-frequency induction furnace in an Ar atmosphere, and casting the molten alloy on the water-cooling copper roll. The thickness of the obtained alloy was about 0.2 to 0.3 mm. The alloy was powdered by the hydrogen decrepitation process, that is, hydrogen absorption at normal temperature and subsequent heating at 600° C. in vacuum for hydrogen desorption. A stearic acid as lubricant with the amount of 0.07 wt % was added and mixed to the coarse alloy powder. The coarse powder was pulverized into a fine powder with a particle size of about 3 m on an average by using a jet milling machine with a nitrogen jet stream. Fine powder was molded while applying a magnetic field of 15 kOe for orientation. The green compact was sintered in vacuum at 1,050 to 1,100° C. for 3 hours, and cooled below 200° C.
The sintered compact was machined into a piece of 20 mm×20 mm×3 mm. The piece was coated with terbium oxide by immersing it in a slurry obtained by mixing 50 wt % of terbium oxide particles with a particle size of 0.5 μm on an average in deionized water, and then drying. The coated piece was held in vacuum at 900-950° C. for 10-20 hours, cooled to 200° C., and aged for 2 hours. Table 1 tabulates the composition of a magnet, although oxygen, nitrogen and carbon concentrations are shown in Table 2. Table 2 tabulates the temperature and time of diffusion treatment, the cooling rate from diffusion treatment temperature to 200° C., the temperature of aging treatment, and magnetic properties. The composition of R—Fe(Co)-M1 phase is shown in Table 3.
The content of (R,HR) in (R,HR)-M1 phase was 50 to 92 at %.
A cross section of the sintered magnet obtained in Example 1 was observed under an electron probe microanalyzer (EPMA). It is observed from
Japanese Patent Application Nos. 2015-072343 and 2016-025548 are incorporated herein by reference.
Although some preferred embodiments have been described, many modifications and variations may be made thereto in light of the above teachings. It is therefore to be understood that the invention may be practiced otherwise than as specifically described without departing from the scope of the appended claims.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2015-072343 | Mar 2015 | JP | national |
| 2016-025548 | Feb 2016 | JP | national |
