EP1780736B1 - R-T-B-Legierung, Verfahren zur Herstellung von Flocken der R-T-B-Legierung, feines Pulver für R-T-B- Seltenerd-Dauermagnete und R-T-B Seltenerd-Dauermagnete - Google Patents

R-T-B-Legierung, Verfahren zur Herstellung von Flocken der R-T-B-Legierung, feines Pulver für R-T-B- Seltenerd-Dauermagnete und R-T-B Seltenerd-Dauermagnete Download PDF

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EP1780736B1
EP1780736B1 EP06020850A EP06020850A EP1780736B1 EP 1780736 B1 EP1780736 B1 EP 1780736B1 EP 06020850 A EP06020850 A EP 06020850A EP 06020850 A EP06020850 A EP 06020850A EP 1780736 B1 EP1780736 B1 EP 1780736B1
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phase
alloy
type
average
type alloy
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EP1780736A1 (de
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Shiro Sasaki
Hiroshi Hasegawa
Kenichiro Nakajima
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Resonac Holdings Corp
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Showa Denko KK
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    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/06—Continuous casting of metals, i.e. casting in indefinite lengths into moulds with travelling walls, e.g. with rolls, plates, belts, caterpillars
    • B22D11/0611—Continuous casting of metals, i.e. casting in indefinite lengths into moulds with travelling walls, e.g. with rolls, plates, belts, caterpillars formed by a single casting wheel, e.g. for casting amorphous metal strips or wires
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/10—Ferrous alloys, e.g. steel alloys containing cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium with copper
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01F—MAGNETS; INDUCTANCES; TRANSFORMERS; SELECTION OF MATERIALS FOR THEIR MAGNETIC PROPERTIES
    • H01F1/00—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties
    • H01F1/01—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials
    • H01F1/03—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity
    • H01F1/032—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials
    • H01F1/04—Magnets or magnetic bodies characterised by the magnetic materials therefor; Selection of materials for their magnetic properties of inorganic materials characterised by their coercivity of hard-magnetic materials metals or alloys
    • H01F1/047—Alloys characterised by their composition
    • H01F1/053—Alloys characterised by their composition containing rare earth metals
    • H01F1/055—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5
    • H01F1/057—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B
    • H01F1/0571—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B in the form of particles, e.g. rapid quenched powders or ribbon flakes
    • H01F1/0575—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B in the form of particles, e.g. rapid quenched powders or ribbon flakes pressed, sintered or bonded together
    • H01F1/0577—Alloys characterised by their composition containing rare earth metals and magnetic transition metals, e.g. SmCo5 and IIIa elements, e.g. Nd2Fe14B in the form of particles, e.g. rapid quenched powders or ribbon flakes pressed, sintered or bonded together sintered
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F9/00—Making metallic powder or suspensions thereof
    • B22F9/02—Making metallic powder or suspensions thereof using physical processes
    • B22F9/04—Making metallic powder or suspensions thereof using physical processes starting from solid material, e.g. by crushing, grinding or milling
    • B22F2009/044—Making metallic powder or suspensions thereof using physical processes starting from solid material, e.g. by crushing, grinding or milling by jet milling
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F2998/00—Supplementary information concerning processes or compositions relating to powder metallurgy
    • B22F2998/10—Processes characterised by the sequence of their steps
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F2999/00—Aspects linked to processes or compositions used in powder metallurgy
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01F—MAGNETS; INDUCTANCES; TRANSFORMERS; SELECTION OF MATERIALS FOR THEIR MAGNETIC PROPERTIES
    • H01F41/00—Apparatus or processes specially adapted for manufacturing or assembling magnets, inductances or transformers; Apparatus or processes specially adapted for manufacturing materials characterised by their magnetic properties
    • H01F41/02—Apparatus or processes specially adapted for manufacturing or assembling magnets, inductances or transformers; Apparatus or processes specially adapted for manufacturing materials characterised by their magnetic properties for manufacturing cores, coils, or magnets
    • H01F41/0253—Apparatus or processes specially adapted for manufacturing or assembling magnets, inductances or transformers; Apparatus or processes specially adapted for manufacturing materials characterised by their magnetic properties for manufacturing cores, coils, or magnets for manufacturing permanent magnets
    • H01F41/0273—Imparting anisotropy

Definitions

  • the present invention relates to a production method of a flake of an R-T-B type alloy.
  • An R-T-B type magnet having a maximum magnetic energy product among permanent magnets is being used for HD (hard disk), MRI (magnetic resonance imaging), various types of motors and the like by virtue of its high-performance characteristics.
  • HD hard disk
  • MRI magnetic resonance imaging
  • various types of motors and the like by virtue of its high-performance characteristics.
  • the R-T-B type magnet comprises Nd, Fe and B as the main components and therefore, the magnets of this type are collectively called an Nd-Fe-B type or R-T-B type magnet.
  • R is primarily Nd with a part being replaced by another rare earth element such as Pr, Dy and Tb and is at least one member selected from rare earth elements including Y; T is Fe with a part being replaced by a transition metal such as Co and Ni; and B is boron and may be partially replaced by C or N.
  • one species or a combination of a plurality of species selected from Cu, Al, Ti, V, Cr, Ga, Mn, Nb, Ta, Mo, W, Ca, Sn, Zr, Hf and the like may be added as the additive element.
  • the R-T-B type alloy which works out to an R-T-B type magnet is an alloy where a ferromagnetic R 2 T 14 B phase contributing to the magnetization activity is the main phase and coexists with a nonmagnetic, rare earth element-enriched and low-melting point R-rich phase.
  • This alloy is an active metal and therefore, generally melted or cast in vacuum or in an inert gas.
  • a sintered magnet is usually produced by a powder metallurgy process as follows.
  • the alloy ingot is ground into an alloy powder of about 3 ⁇ m (as measured by FSSS (Fisher sub-sieve sizer)), press-shaped in a magnetic field, sintered at a high temperature of about 1,000 to 1,100°C in a sintering furnace, then subjected to, if desired, heat treatment and machining, and further plated for enhancing the corrosion resistance, thereby completing a sintered magnet.
  • FSSS Fisher sub-sieve sizer
  • the R-rich phase plays the following important roles:
  • ⁇ -Fe has deformability and remains in the grinder without being ground, and this not only decreases the grinding efficiency at the grinding of alloy but also affects the compositional fluctuation or particle size distribution. If ⁇ -Fe still remains in the magnet after sintering, reduction in the magnetic characteristics of the magnet results. Accordingly, ⁇ -Fe has been dealt with as a material which should be eliminated from the raw material alloy as much as possible. For this purpose, an alloy has been heretofore subjected to a homogenization treatment at a high temperature for a long time to eliminate ⁇ -Fe.
  • ⁇ -Fe When the amount of ⁇ -Fe in the raw material alloy is small, this may be removed by a homogenization heat treatment.
  • ⁇ -Fe is present as a peritectic nucleus and therefore, its elimination requires solid phase diffusion for a long time. In the case of an ingot having a thickness of several cm and a rare earth content of 33% or less, elimination of ⁇ -Fe is practically impossible.
  • SC method strip casting method
  • the SC method is a method of solidifying an alloy through rapid cooling, where a molten alloy is cast on a copper roll of which inside is water-cooled, and a flake of 0.1 to 1 mm is produced.
  • the molten alloy is supercooled to the temperature where the main R 2 T 14 B phase is produced, so that an R 2 T 14 B phase can be produced directly from a molten alloy and the precipitation of ⁇ -Fe can be suppressed.
  • the alloy comes to have a fine crystal texture, so that an alloy having a texture allowing for fine dispersion of an R-rich phase can be produced.
  • the R-rich phase expands by reacting with hydrogen in a hydrogen atmosphere and becomes a brittle hydride.
  • fine cracking commensurate with the dispersion degree of R-rich phase can be introduced.
  • the internal R-rich phase in the alloy produced by the SC method is thus finely dispersed, and this leads to good dispersibility of the R-rich phase also in the magnet after grinding and sintering, thereby succeeding in enhancing the magnetic characteristics of the magnet (see, for example, Patent Document 1).
  • the alloy flake produced by the SC method is excellent also in the texture homogeneity.
  • the texture homogeneity can be compared by the crystal grain diameter or the dispersed state of R-rich phase.
  • a chill crystal is sometimes generated on the casting roll side of the alloy flake (hereinafter referred to as a "mold face side"), but an appropriately fine homogeneous texture yielded by the solidification through rapid cooling can be obtained as a whole.
  • the R-rich phase is finely dispersed and the precipitation of ⁇ -Fe is also suppressed, so that in the production of a sintered magnet, the homogeneity of the R-rich phase in the final magnet can be increased and the adverse effect of ⁇ -Fe on the grinding and magnetism can be prevented.
  • the R-T-B type alloy ingot produced by the SC method has an excellent texture for the production of a sintered magnet.
  • demands for high-level control of the raw material alloy texture, particularly, the presence state of the R-rich phase are increasing.
  • the present inventors have previously made studies on the relationship between the texture of the cast-produced R-T-B type alloy and the behavior at the hydrogen cracking or pulverization and found that in order to control the particle size of the alloy powder for a sintered magnet, the control of the dispersed state of R-rich phase is important (see, for example, Patent Document 2). Also, it has been found that fine division readily occurs in the region where the R-rich phase produced on the mold face side in the alloy (fine R-rich phase region) is extremely finely dispersed, as a result, the grinding stability of the alloy is deteriorated and at the same time, the particle size distribution of the powder is broadened. This finding leads to understanding that reduction of the fine R-rich phase region is necessary for the enhancement of characteristics of the magnet.
  • EP-A1-1 033 415 discloses a main-phase alloy consisting of about 20 % by weight of Nd, 7 % by weight of Dy, 1 % by weight of B and 72 % by weight of Fe. The alloy has been strip cast into flakes of a thickness of 0.30 mm.
  • US 6328825 B1 discloses alloys for rare-earth based permanent magnets which are of the R-T-B type and which contain an R 2 T 17 phase having an average grain size between 3 and 20 micrometers.
  • an object of the present invention is to provide flakes of an R-T-B type alloy as a raw material of a rare earth-based permanent magnet having excellent magnetic characteristics.
  • the present inventors have particularly observed the cross-sectional texture of alloy flakes which are cast and solidified under various conditions, and found that there is a relationship between the precipitated state of 2-17 phase and the magnetic characteristics and when a fine 2-17 phase (R 2 T 17 phase) is precipitated in the alloy, the magnetic characteristics can be enhanced. Also, the present inventors have confirmed the fact that when a sintered magnet is produced from an alloy allowing for the presence of a fine R 2 T 17 phase or an alloy prepared by controlling the cooling rate on the casting roll or the temperature on separating from the casting roll in the SC method, the coercive force thereof is stably increased and excellent magnetic characteristics are obtained. The present invention has been accomplished based on these findings.
  • the present invention provides the following invention.
  • the volume percentage of a region containing an R 2 T 17 phase having an average grain diameter of 3 ⁇ m or less in the short axis direction is 0.5 to 10%, so that a rare earth permanent magnet having a high coercive force and excellent magnetic characteristics can be realized.
  • the alloy flake is produced by the SC method and not only the average thickness is set to from 0.1 to 1 mm but also the average molten alloy supply rate to the casting roll is set to 10 g/sec or more per 1-cm width, so that an R-T-B type alloy having a high coercive force can be obtained.
  • Fig. 1 is a photograph showing one example of the R-T-B type alloy flakes obtained by the method of the present invention, and this photograph is taken when the cross-section of the R-T-B type alloy flake is observed by a scanning electron microscope (SEM).
  • the R-T-B type alloy flake shown in Fig. 1 is produced by the SC method.
  • This R-T-B type alloy flake has a composition comprising, in terms of the weight ratio, 22% of Nd, 9% of Dy, 0.95% of B, 1% of Co, 0.3% of Al and 0. 1% of Cu, with the balance being Fe.
  • an R 2 T 17 phase is not precipitated and even in an equilibrium state at an ordinary temperature, an R 2 T 17 phase is not stably present at a temperature of 1,170°C or less, which is the melting point of the R 2 T 14 B phase.
  • the R-rich phase is indicated by a white color and the R 2 T 17 phase is indicated by a slightly darker color than the main R 2 T 14 B phase.
  • the R-T-B type alloy flake is entirely composed of a columnar crystal which is an R 2 T 14 B phase, and an R-rich phase extending in the long axis direction of the columnar crystal.
  • the R 2 T 14 B phase mainly comprises a columnar crystal and partially comprises an equi-axed crystal, and the average crystal grain diameter thereof in the short axis direction is from 10 to 50 ⁇ m.
  • a linear R-rich phase extending along the long axis direction of the columnar crystal or a particulated or partially broken R-rich phase is present at the grain boundary and within the grain.
  • the average distance between R-rich phases present at the grain boundary and within the grain of the R 2 T 14 B phase is from 3 to 10 ⁇ m. Also, as shown in Fig. 1 , a region allowing for coexistence of very fine R 2 T 17 phase and R-rich phase is present in the R-T-B type alloy flake, with each phase occupying an area percentage (volume percentage) of about 3%.
  • the R 2 T 17 phase is an intermetallic compound not having a composition width stably present from the ordinary temperature to high temperature region in the binary phase diagram of a rare earth-iron system.
  • This phase is a soft magnetic phase with in-plane anisotropy at an ordinary temperature and when present in an R-T-B type sintered magnet, functions as a nucleation site of the reversed magnetic domain to cause reduction in the coercive force.
  • this phase disappears in the sintering process and becomes harmless in many cases.
  • the R 2 T 17 phase is an intermetallic compound having no ductility and therefore, scarcely affects the grinding behavior in the magnet production step.
  • the R 2 T 17 phase precipitates as a primary crystal instead of ⁇ -Fe.
  • This is magnetically soft but unlike ⁇ -Fe, its effect on the grinding behavior is small as described above and in the SC method, the production thereof can be prevented similarly to ⁇ -Fe by the large supercooling.
  • Fig. 2 is an enlarged photograph of the photograph shown in Fig. 1 , and this photograph shows the region encircled with a white line in Fig. 1 and the peripheral region thereof.
  • the region encircled with a white line indicates the region where the R 2 T 17 phase is precipitated.
  • the average crystal grain diameter in the short axis direction of the R 2 T 17 phase is preferably smaller.
  • the average crystal grain diameter is approximately from 1 to 2 ⁇ m in the R-T-B type alloy flake shown in Fig. 1 .
  • the phase when the crystal grain diameter of the R 2 T 17 phase becomes large, the phase can hardly disappear at the sintering and if remains in the sintered body, the remaining phase incurs deterioration of the magnetic characteristics.
  • This phase may be caused to disappear by increasing the sintering temperature or sintering time, but the main phase crystal grain is also coarsened to give rise to decrease in the coercive force.
  • the average crystal grain diameter in the short axis direction of the R 2 T 17 phase By controlling the average crystal grain diameter in the short axis direction of the R 2 T 17 phase to 3 ⁇ m or less, the effects of the present invention can be brought out.
  • the adverse effect of the coarse R 2 T 17 phase appears as the decrease in the orientation degree, in addition to the possibility of remaining in the sintered body or the reduction in the coercive force or squareness resulting from increase of the sintering temperature or time.
  • Two causes are considered for the decrease in the orientation degree.
  • a first cause is the in-plane anisotropy of the R 2 T 17 phase. This phase differs also in the magnetization from the R 2 T 14 B phase and therefore, may affect the orientation behavior of the R 2 T 14 B phase during shaping in a magnetic field.
  • the second cause it is considered that a small R 2 T 17 phase coalesces with the adjacent R 2 T 14 B phase or converts into a liquid phase, however, when the R 2 T 17 phase becomes large to an extent of up to the grain size of the main R 2 T 14 B phase, the disappearance takes time and until reaching the disappearance, the phase reacts with a B-rich phase or the like in the neighborhood to produce and grow an R 2 T 14 B phase nucleus.
  • the newly nucleated and grown R 2 T 14 B phase has a random crystal orientation and therefore, the orientation degree as a whole decreases.
  • a region where, as shown in Fig. 2 , an R 2 T 17 phase is precipitated is defined as an "R 2 T 17 phase-containing region".
  • This region can be easily distinguished from the peripheral alloy texture portion primarily comprising a main phase of columnar crystal and an R-rich phase extending in the long axis direction of the columnar crystal.
  • the volume percentage of the phase is preferably from 0.5 to 10%.
  • volume percentage of the R 2 T 17 phase having an average grain diameter of 3 ⁇ m or less in the short axis direction is less than 0.5%, the effect of improving the sinterability and enhancing the magnetic characteristics decreases, whereas if the volume percentage of the R 2 T 17 phase having an average grain diameter of 3 ⁇ m or less in the short axis direction exceeds 10%, the composition or particle size at the grinding greatly fluctuates to cause large fluctuation of magnetic characteristics and also the magnetization decreases due to reduction in the orientation degree.
  • the volume percentage of the R 2 T 17 phase having an average grain diameter of 3 ⁇ m or less in the short axis direction is more preferably from 1 to 5%.
  • the average crystal grain diameter in the short axis direction of the R 2 T 17 phase exceeds 5 ⁇ m, the effect of precipitating an R 2 T 17 phase becomes poor and if the volume percentage of such an R 2 T 17 phase-containing region exceeds 10%, the magnetic characteristics greatly fluctuate. Also, if the average grain diameter in the short axis direction of the R 2 T 17 phase is 10 ⁇ m or more and the volume percentage of the phase is 10% or more, the magnetic characteristics apparently deteriorate.
  • the volume percentage of the region containing an R 2 T 17 phase having an average grain diameter of 10 ⁇ m or more in the short axis direction is more preferably 5% or less.
  • the R 2 T 17 phase present in the R-T-B type alloy flakes exists as a non-equilibrium phase (metastable phase).
  • the precipitate present as a metastable phase is in an energetically high state and therefore, disappears in a high-temperature region where diffusion satisfactorily functions, for example, at about 1/2 of the decomposition temperature shown by the absolute temperature of the compound.
  • the time required for the R 2 T 17 phase present as a non-equilibrium phase to disappear varies depending on the temperature or the size of R 2 T 17 phase, but the disappearance is easily attained as compared with the R 2 T 17 phase present in an equilibrium state and in a magnet production process, the phase disappears within a general sintering time of several hours or less.
  • an R-rich phase having almost the same size is present together in the R 2 T 17 phase precipitation site of the R-T-B type alloy flakes.
  • the R-rich phase expands by absorbing hydrogen to become brittle in the hydrogen cracking step before pulverization and works out to a starting point for fine cracking.
  • the R 2 T 17 phase-containing region is ground more finely than the R 2 T 14 B phase and the effect of fine R 2 T 17 phase is more enhanced.
  • good dispersibility of the R-rich phase is obtained and the sinterability is more improved.
  • the average grain diameter in the short axis direction of the R-rich phase is increased to about 10 ⁇ m, the proportion of fine powder comprising only an R-rich phase increases and the homogeneity in the powder compact decreases, giving rise to worsened sinterability.
  • the homogeneity of the R-rich phase in the sintered body also decreases and therefore, the coercive force is decreased.
  • the hydrogenated R-rich phase is more brittle than the main phase and finely divided in a short time at the initial stage of grinding to increase the fluctuation of composition or particle size at the grinding, and this gives rise to fluctuation of characteristics.
  • the average grain diameter in the short axis direction of the R-rich phase is preferably 3 ⁇ m or less.
  • the R-T-B type alloy obtained by the method of the present invention shown in Fig. 1 is a flake produced by the strip casting method.
  • the R-T-B type alloy can be cast-produced by the following SC method.
  • Fig. 3 is a schematic view showing the apparatus for casting by the SC method.
  • an R-T-B type alloy is melted by using a refractory crucible 1 in a vacuum or inert gas atmosphere because of its active property.
  • the molten alloy after melting the R-T-B type alloy is kept at 1,300 to 1,500°C for a predetermined time and then supplied to a rotating roll 3 for casting (casting roll) with the inside being water-cooled, through a tundish 2 in which, if desired, a rectification mechanism or a slug removal mechanism is provided.
  • the supply rate of the molten alloy and the rotation velocity of the casting roll are controlled according to the desired alloy thickness. In general, the rotation number of the casting roll is approximately from 0.5 to 3 m/s in terms of the peripheral velocity.
  • the material of the casting roll is suitably copper or a copper alloy because of good heat conductivity and easy availability.
  • a cleaning device is provided, whereby the quality of the cast-produced R-T-B type alloy flakes is stabilized.
  • the alloy 4 solidified on the casting roll is separated from the roll on the side opposite the tundish and recovered by a collection container 5. It is disclosed in JP-A-10-36949 that the texture state of the R-rich phase can be controlled by providing a heating and cooling mechanism in the collection container. In the present invention, in order to control the dispersed state of the R-rich phase, the cooling and thermal insulation after separation from the roll may be divided into several steps and thereby controlled.
  • a heating and cooling mechanism is provided before finally collecting the alloy by the collection container and the alloy is heated, thermally insulated and cooled, whereby the size and homogeneity of the alloy texture, the particle size distribution of the fine particle after grinding, the supply to the metal mold, the bulk density, the adjustment of percentage shrinkage at sintering, and the magnetic characteristics can be improved.
  • the R-T-B type alloy obtained by the method of the present invention is a flake having an average thickness of 0.1 to 1 mm. If the average thickness of the flake is less than 0.1 mm, the solidification rate is excessively increased and the R-rich phase is too finely dispersed, whereas if the average thickness of the flake exceeds 1 mm, the solidification rate decreases and this incurs reduction in the dispersibility of the R-rich phase, precipitation of ⁇ -Fe, coarsening of the R 2 T 17 phase, or the like.
  • the average molten alloy supply rate to the casting roll is 10 g/sec or more and 100 g/sec or less per 1-cm width. If the molten alloy supply rate is less than 10 g/sec, the molten alloy is not thinly wetted and spread on the roll but shrinks because of the viscosity of the molten alloy itself or wettability to the casting roll surface and fluctuation of the alloy quality is brought about, whereas if the average molten alloy supply rate to the casting roll exceeds 100 g/sec per 1-cm width, cooling on the casting roll is insufficient and this causes coarsening of the texture, precipitation of ⁇ -Fe, or the like.
  • the supply rate can be controlled to a certain extent by providing a rectification mechanism in the tundish.
  • the average cooling rate of the R-T-B type alloy on the casting roll is preferably from 500 to 3,000°C/sec. If the average cooling rate is less than 500°C/sec, precipitation of ⁇ -Fe or texture coarsening of the R-rich phase, R 2 T 17 phase or the like occurs due to insufficient cooling rate, whereas if the average cooling rate exceeds 3,000°C/sec, the supercooling becomes too large and the production of the R 2 T 17 phase-containing region as a characteristic feature of the present invention decreases.
  • the average temperature of the R-T-B type alloy on separating from the casting roll subtly varies due to fine difference in the degree of contact with the casting roll, fluctuation of the thickness, or the like.
  • the average temperature of the alloy on separating from the casting roll can be obtained, for example, by scanning the alloy surface in the width direction by a radiation thermometer from start to finish of the casting, thereby measuring the temperature, and averaging the measured values.
  • the average temperature of the alloy on separating from the casting roll is preferably 100 to 400°C lower, more preferably 100 to 300°C lower, than the solidification temperature of the R 2 T 14 B phase in an equilibrium state of the molten R-T-B type alloy.
  • the melting temperature of the R 2 T 14 B phase is acknowledged to be 1, 150°C in the Nd-Fe-B ternary system but varies according to the substitution of Nd by other rare earth elements, the substitution of Fe by other transition elements, and the kind and amount added of the other additive element. If the difference between the average temperature of the R-T-B type alloy on separating from the casting roll and the solidification temperature of the R 2 T 14 B phase in an equilibrium state of the R-T-B type alloy is less than 100°C, this corresponds to an insufficient cooling rate, whereas if this difference exceeds 400°C, the supercooling of molten alloy becomes excessively large due to the too high cooling rate.
  • the degree of supercooling of the molten alloy is not uniform in the alloy but varies according to the degree of contact with the casting roll or the distance from the contact part with the casting roll.
  • the alloy temperature on separating from the casting roll varies also within the same casting step (tap) and if the variation width is large, this brings about fluctuation of the texture or quality. Therefore, the variation width of temperature within the tap is suitably smaller than 200°C, preferably 100°C or less, more preferably 50°C, still more preferably 20°C.
  • the average temperature of the R-T-B type alloy on separating from the casting roll is 300°C or more lower than the solidification of the R 2 T 14 B phase in an equilibrium state of the molten alloy composition, the amount of the fine R 2 T 17 phase precipitated decreases and the effect of improving magnetic characteristics becomes poor. This infers that precipitation of the R 2 T 17 phase is generated in a portion where the supercooling degree is relatively small. Also, if the proportion of the heavy rare earth occupying in the rare earth is decreased, the amount of the R 2 T 17 phase precipitated also decreases and the presence of the phase cannot be confirmed, but the effect of enhancing the magnetic characteristics continues.
  • the cooling on the roll is called primary cooling and this reference indicates that cooling is preferably performed to a cast strip temperature of 700 to 1,000°C at a cooling rate of 2 ⁇ 10 3 to 7 ⁇ 10 3 °C/sec.
  • a fine powder for R-T-B type rare earth permanent magnets is first produced from the R-T-B type alloy.
  • the fine powder for R-T-B type rare earth permanent magnets is obtained, for example, by a method of performing hydrogen cracking of a flake comprising the R-T-B type alloy and then pulverizing the flake by using a grinder such as jet mill.
  • a grinder such as jet mill.
  • a hydrogen absorption step of keeping the flake in a hydrogen atmosphere under a predetermined pressure is preferably performed in advance.
  • the obtained fine powder for R-T-B type rare earth permanent magnets is, for example, press-shaped by a shaping machine or the like in a transverse magnetic field and sintered, whereby an R-T-B type rare earth permanent magnet is obtained.
  • the fine R 2 T 17 phase or the fine R-rich phase present together with the R 2 T 17 phase swiftly converts into a liquid phase at the sintering, contributing to the enhancement of sinterability or dispersibility of the R-rich phase, so that a rare earth magnet having a high coercive force and excellent magnetic characteristics can be realized.
  • the R 2 T 17 phase-containing alloy includes, for example, an alloy where an R 2 T 17 phase-containing alloy powder by the SC method is mixed with an alloy powder having an R 2 T 14 B phase as the main phase, which is obtained by the SC method, to increase the volume percentage of the R 2 T 14 B phase (see, for example, JP-A-7-45413 ).
  • the R 2 T 17 phase-containing alloy described in JP-A-7-45413 is formulated such that the R 2 T 17 phase precipitates in an equilibrium state resulting from decrease in the B amount. In this case, the volume percentage of the R 2 T 17 phase in the alloy increases and the crystal gain diameter of the R 2 T 17 phase in the alloy also increases.
  • the particle size of the R 2 T 17 phase-containing alloy powder needs to be made small. If the particle size is not made small, elevation of the sintering temperature or prolongation of the sintering time is required for obtaining satisfactory diffusion necessary for the disappearance of R 2 T 17 phase, as a result, the texture of the sintered body is coarsened and reduction in the coercive force is caused. Also, it is easily presumed from the compositional formulation that the R 2 T 17 phase described in JP-A-7-45413 is stably present from an ordinary temperature to the decomposition temperature thereof.
  • JP-A-7-45413 indicates that the addition of the R 2 T 17 phase brings about increase of the liquid phase, but is silent on the discussion from the kinetic aspect until reaching the liquid phase.
  • the R 2 T 17 phase constituting the R-T-B type alloy flakes obtained by the method of the present invention is precipitated as a non-equilibrium phase.
  • the R 2 T 17 phase present as a non-equilibrium phase readily disappears as compared with the R 2 T 17 phase present in an equilibrium state and disappears within a sintering time which is generally several hours in the magnet production process.
  • the method for producing an R-T-B type alloy flakes having a composition allowing for precipitation of an R 2 T 17 phase is described, but the production method of an R-T-B type alloy flake of the present invention is not limited to the method for producing an R-T-B type alloy flakes having a composition allowing for precipitation of an R 2 T 17 phase, and an R-T-B type alloy having a composition not allowing for precipitation of an R 2 T 17 phase may be produced by the production method of an R-T-B type alloy flake of the present invention.
  • Raw materials of metallic neodymium, metallic dysprosium, ferroboron, cobalt, aluminum, copper and iron were weiighted to give an alloy composition comprising, in terms of the weight ratio, 22% of Nd, 9% of Dy, 0.95% of B, 1% of Co, 0.3% of Al and 0.1% of Cu, with the balance being Fe, and melted in an alumina crucible in an argon gas atmosphere at 1 atm by using a high-frequency melting furnace, and the molten alloy was cast by the SC method to produce an alloy flake.
  • the rotating roll for casting had a diameter of 600 mm and was made of an alloy obtained by mixing slight amounts of Cr and Zr with copper, and the inside thereof was water-cooled.
  • the peripheral velocity of the roll at the casting was 1.3 m/sec
  • the average molten alloy supply rate to the casting roll was 28 g/sec per 1-cm width
  • the average temperature of the alloy on separating from the casting roll was measured by a radiation thermometer and found to be 890°C. In the measured values, the difference between the maximum temperature and the minimum temperature was 35°C. Since the melting point of the R 2 T 14 B phase of this alloy is about 1,170°C, the difference from the average separation temperature is 280°C.
  • the average cooling rate of the R-T-B type alloy on the casting roll was 980°C/sec and the average thickness was 0.29 mm.
  • the recovery container for housing alloy flakes separated from the roll had a partition plate through which a cooling Ar gas was flowed.
  • the production conditions of the alloy flake are shown in Table 1.
  • “Supply Rate” indicates the average molten alloy supply rate to the casting roll, and this is the amount supplied per 1-cm width per second; “Cooling Rate” indicates the average cooling rate of the R-T-B type alloy on the casting roll; “Solidification Temperature” is a solidification temperature (melting point) of the R 2 T 14 B phase in an equilibrium state of the R-T-B type alloy; “Average Temperature Difference” indicates the temperature difference between the “Solidification Temperature” and the average temperature of the R-T-B type alloy on separating from the casting roll; and “Average Thickness” indicates an average thickness of flakes produced by the strip casting method.
  • the volume percentage was calculated for each of the R 2 T 17 phases having an average grain diameter of 3 ⁇ m or less and an average grain diameter of 5 ⁇ m or more in the region.
  • the average grain diameter and volume percentage of each texture of the alloy flake are shown in Table 1.
  • Average Grain Diameter 1 and Volume Percentage 1 indicate the average grain diameter of the R 2 T 17 phase having an average grain diameter of 3 ⁇ m or less in the short axis direction and the volume percentage of the region containing the R 2 T 17 phase; Average Grain Diameter 2 and Volume.
  • Percentage 2 indicate the average grain diameter of the region containing an R 2 T 17 phase having an average grain diameter of 5 ⁇ m or more in the short axis direction and the volume percentage of the region containing the R 2 T 17 phase; and Average Grain Diameter 3 and Volume Percentage 3 indicate the average grain diameter of the R-rich phase having an average grain diameter of 3 ⁇ m or less in the short axis direction present in the region containing an R 2 T 17 phase having an average grain diameter of 3 ⁇ m or less in the short axis direction and the volume percentage of the region.
  • the obtained alloy flakes were heat-treated at 1,000°C for 2 hours and a backscattered electron image (BEI) of each alloy flake was photographed at a magnification of 350 by a scanning electron microscope (SEM), as a result, complete disappearance of the R 2 T 17 phase was confirmed.
  • BEI backscattered electron image
  • SEM scanning electron microscope
  • Raw materials were blended to give the same composition as in Example 1, and melting and casting by the SC method were performed in the same manner as in Example 1.
  • the peripheral velocity of the roll at the casting was 0.8 m/sec
  • the average molten alloy supply rate to the casting roll was 13.0 g/sec per 1-cm width
  • the average temperature of the alloy on separating from the casting roll, measured by a radiation thermometer was 630°C
  • the difference between the maximum temperature and the minimum temperature of the measured values was 160°C. Since the melting point of the R 2 T 14 B phase of this alloy is about 1,170°C, the difference from the average separation temperature is 540°C.
  • the average cooling rate of the R-T-B type alloy on the casting roll was 920°C/sec and the average thickness was 0.23 mm.
  • the obtained alloy flakes were evaluated in the same manner as in Example 1, and the results are shown in Table 1. Incidentally, in Comparative Example 1, the R 2 T 17 phase-containing region could not be confirmed.
  • the peripheral velocity of the roll at the casting was 1.3 m/sec
  • the average molten alloy supply rate to the casting roll was 28 g/sec per 1-cm width
  • the average temperature of the alloy on separating from the casting roll, measured by a radiation thermometer was 850°C
  • the difference between the maximum temperature and the minimum temperature of the measured values was 20°C.
  • the melting point of the R 2 T 14 B phase of this alloy is about 1,140°C
  • the difference from the average separation temperature is 290°C.
  • the average cooling rate of the R-T-B type alloy on the casting roll was 1,060°C/sec and the average thickness was 0.29 mm.
  • the obtained alloy flakes were evaluated in the same manner as in Example 1, and the results are shown in Table 1.
  • the composition of the R-T-B type alloy of Example 2 is formulated not to allow for precipitation of the R 2 T 17 phase and in Example 2, the R 2 T 17 phase-containing region could not be confirmed.
  • Raw materials were blended to give the same composition as in Example 1, and melting and casting by the SC method were performed in the same manner as in Example 1.
  • the peripheral velocity of the roll at the casting was 0.8 m/sec
  • the average molten alloy supply rate to the casting roll was 13.0 g/sec per 1-cm width
  • the average temperature of the alloy on separating from the casting roll, measured by a radiation thermometer was 620°C
  • the difference between the maximum temperature and the minimum temperature of the measured values was 180°C. Since the melting point of the R 2 T 14 B phase of this alloy is about 1,140°C, the difference from the average separation temperature is 520°C.
  • the average cooling rate of the R-T-B type alloy on the casting roll was 930°C/sec and the average thickness was 0.23 mm.
  • the obtained alloy flakes were evaluated in the same manner as in Example 1, and the results are shown in Table 1. Incidentally, in Comparative Example 2, the R 2 T 17 phase-containing region could not be confirmed.
  • Raw materials were blended to give the same composition as in Example 1, and melting and casting by the SC method were performed in the same manner as in Example 1.
  • the peripheral velocity of the roll at the casting was 0.8 m/sec
  • the average molten alloy supply rate to the casting roll was 70 g/sec per 1-cm width
  • the average temperature of the alloy on separating from the casting roll, measured by a radiation thermometer was 1,000°C
  • the difference between the maximum temperature and the minimum temperature of the measured values was 250°C. Since the melting point of the R 2 T 14 B phase of this alloy is about 1,170°C, the difference from the average separation temperature is 170°C.
  • the average cooling rate of the R-T-B type alloy on the casting roll was 290°C/sec and the average thickness was 1.2 mm.
  • the obtained alloy flakes were evaluated in the same manner as in Example 1, and the results are shown in Table 1.
  • Comparative Example 3 the presence of a slight amount of the R 2 T 17 phase-containing region was confirmed even after the alloy flake was heat-treated at 1,000°C for 2 hours similarly to Example 1. This is caused because the grain size of the R 2 T 17 phase present before heat treatment is large and a long time is necessary for the phase to disappear.
  • the R 2 T 17 phase is not stably present at a temperature of 1,170°C or less, which is the melting point of the R 2 T 14 B phase.
  • the alloy flake obtained in Example 1 was subjected to hydrogen cracking and pulverization by a jet mill.
  • the conditions in the hydrogen absorption step as the pre-step of the hydrogen cracking step were a 100% hydrogen atmosphere, a pressure of 2 atm, and a holding time of 1 hour.
  • the temperature of the metal strip at the initiation of a hydrogen absorption reaction was 25°C.
  • the conditions in the dehydrogenation step as the post-step were an in-vacuum atmosphere of 0.133 hPa, 500°C and a holding time of 1 hour.
  • the obtained powder material was press-shaped by a shaping machine in a transverse magnetic field in a 100% nitrogen atmosphere.
  • the shaping pressure was 0.8 t/cm 2 and the magnetic field in the die cavity was set to 15 kOe.
  • the resulting powder compact was sintered by holding it in vacuum of 1.33 ⁇ 10 -5 hPa at 500°C for 1 hour, then in vacuum of 1.33 ⁇ 10 -5 hPa at 800°C for 2 hours, and further in vacuum of 1.33 ⁇ 10 -5 hPa at 1,030°C for 2 hours.
  • the sintering density was 7.7 g/cm 3 or more and this was a sufficiently large density.
  • This sintered body was further heat-treated at 530°C for 1° hour in an argon atmosphere to produce a sintered magnet.
  • Example 3 The magnetic characteristics of this sintered body of Example 3 were measured by a direct current BH curve tracer, and the results are shown in Table 2. [Table 2] Br T iHc kA/m (BH)max kJ/m 3 SQ (%) Example 3 1.16 2680 260 91 Example 4 1.45 1247 403 92 Comparative Example 4 1.16 2551 259 91 Comparative Example 5 1.45 1068 403 91 Comparative Example 6 1.1 2425 234 90
  • Example 2 Using the alloy flake obtained in Example 2, a sintered magnet was produced by the same method as in Example 3. The magnetic characteristics of this sintered magnet of Example 4 were measured by a direct current BH curve tracer, and the results are shown in Table 2.
  • the alloy flake obtained in Comparative Example 2 was ground by the same method as in Example 3 to obtain a fine powder.
  • the magnetic characteristics of the obtained sintered magnet of Comparative Example 5 were measured by a direct current BH curve tracer, and the results are shown in Table 2.
  • the alloy flake obtained in Comparative Example 3 was ground by the same method as in Example 3 to obtain a fine powder.
  • the magnetic characteristics of the obtained sintered magnet of Comparative Example 6 were measured by a direct current BH curve tracer, and the results are shown in Table 2.
  • Example 4 using the alloy of Example 2 which has a composition containing no heavy rare earth and not allowing for precipitation of an R 2 T 17 phase and is produced by the production method of an R-T-B type alloy flake of the present invention, the coercive force is large as compared with Comparative Example 5 where the average temperature difference exceeds 300°C.
  • the cause of this is still being studied, but one presumable reason therefor is that by virtue of the low solidification rate, the number of crystal defects is smaller in the alloy of Example 2.

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Claims (3)

  1. Verfahren zum Herstellen von Flocken einer Legierung vom R-T-B-Typ mit einer mittleren Dicke von 0,1 bis 1 mm durch ein Dünnbandgießverfahren, welches das Festlegen der mittleren Rate der Zuführung der geschmolzenen Legierung zu der Formwalze auf 10 g/sec. bis 100 g/sec. pro 1 cm Breite umfasst, wobei R mindestens ein Seltenerdelement, einschließlich Y ist, wobei R hauptsächlich Nd ist und ein Teil durch Pr, Dy und/oder Tb ersetzt ist,
    T mindestens ein Element ist, das aus der Gruppe der Übergangselemente ausgewählt ist, wobei 80 bis 99,9 Gew.-% des Gesamtgehalts dieser Übergangselemente in der Legierung durch Fe gebildet wird, und
    B Bor ist,
    wobei der Volumenprozentsatz der Region, die eine R2T17-Phase mit einem mittleren Korndurchmesser von 3 µm oder weniger in Richtung der kurzen Achse enthält, 0,5 bis 10 % ist.
  2. Verfahren nach Anspruch 1, wobei die mittlere Kühlrate der Legierung vom R-T-B-Typ auf der Formwalze 500 bis 3000°C/sec ist.
  3. Verfahren nach Anspruch 1 oder 2, wobei die mittlere Temperatur der Legierung vom R-T-B-Typ beim Abtrennen von der Formwalze 100 bis 400°C niedriger ist als die Verfestigungstemperatur der R2T14B-Phase in einem Gleichgewichtszustand der Legierung vom R-T-B-Typ.
EP06020850A 2005-10-31 2006-10-04 R-T-B-Legierung, Verfahren zur Herstellung von Flocken der R-T-B-Legierung, feines Pulver für R-T-B- Seltenerd-Dauermagnete und R-T-B Seltenerd-Dauermagnete Active EP1780736B1 (de)

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JP5572673B2 (ja) * 2011-07-08 2014-08-13 昭和電工株式会社 Ｒ−ｔ−ｂ系希土類焼結磁石用合金、ｒ−ｔ−ｂ系希土類焼結磁石用合金の製造方法、ｒ−ｔ−ｂ系希土類焼結磁石用合金材料、ｒ−ｔ−ｂ系希土類焼結磁石、ｒ−ｔ−ｂ系希土類焼結磁石の製造方法およびモーター
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