WO2002093591A2 - Iron-based rare earth alloy nanocomposite magnet and method for producing the same - Google Patents

Abstract

An iron-based rare earth alloy nanocomposite magnet has a composition represented by (Fe1-mTm)¿100-x-y-z?QxRyTiz where T is Co and/or Ni, Q is B and/or C and R is rare earth element(s) including substantially no La or Ce. x, y, z and m satisfy 10 at%<x≤17 at%, 7 at%≤y<10 at%, 0.5 at%≤z≤6 at% and 0≤m≤0.5, respectively. The magnet includes crystal grains of an R2T14Q type compound having an average grain size of 20 nm to 200 nm and a ferromagnetic iron-based boride that exists in a grain boundary between the crystal grains of the R2T14Q type compound. The boride is dispersed in, or present in the form of a film over, the grain boundary to cover the surface of the crystal grains of the R2T14Q type compound at least partially.

Description

DESCRIPTION

IRON-BASED RARE EARTH ALLOY NANOCOMPOSITE MAGNET AND

METHOD FOR PRODUCING THE SAME

TECHNICAL FIELD

The present invention generally relates to a method for producing a

permanent magnet that is applicable for use in motors and actuators of various

types. More particularly, the present invention relates to a method for producing

an iron-based rare earth alloy nanocomposite magnet including multiple

ferromagnetic phases.

BACKGROUND ART

Recently, it has become more and more necessary to further improve the

performance of, and further reduce the size and weight of, consumer electronic

appliances, office automation appliances and various other types of electric

equipment. For these purposes, a permanent magnet for use in each of these

appliances is required to maximize its performance to weight ratio when operated

as a magnetic circuit. For example, a permanent magnet with a remanence B_r of

about 0.5 T or more is now in high demand. Hard ferrite magnets have been used widely because magnets of this type are relatively inexpensive. However, the

hard ferrite magnets cannot achieve the high remanence B_r of about 0.5 T or

more.

An Sm-Co based magnet, produced by a powder metallurgical process,

is currently known as a typical permanent magnet that achieves the high

remanence B_r of about 0.5 T or more. However, the Sm-Co based magnet is

expensive, because Sm and Co are both expensive materials.

Examples of other high-remanence magnets include an Nd-Fe-B based

sintered magnet produced by a powder metallurgical process and an Nd-Fe-B

based rapidly solidified magnet produced by a melt quenching process. An

Nd-Fe-B based sintered magnet is disclosed in Japanese Laid-Open Publication

No. 59-46008, for example, and an Nd-Fe-B based rapidly solidified magnet is

disclosed in Japanese Laid-Open Publication No. 60-9852, for instance.

The Nd-Fe-B based sintered magnet is mainly composed of relatively

inexpensive Fe (typically at about 60 wt% to about 70 wt% of the total weight),

and is much less expensive than the Sm-Co based magnet. Nevertheless, it is

still expensive to produce the Nd-Fe-B based sintered magnet. This is partly

because huge equipment and a great number of manufacturing and processing

steps are required to separate and purify, or to obtain by reduction reaction, Nd, which usually accounts for about 10 at% to about 15 at% of the magnet. Also, a

powder metallurgical process normally requires a relatively large number of

manufacturing and processing steps by its nature.

Compared to an Nd-Fe-B based sintered magnet formed by a powder

metallurgical process, an Nd-Fe-B based rapidly solidified magnet can be

produced at a lower cost by a melt quenching process. This is because an

Nd-Fe-B based rapidly solidified magnet can be produced through relatively

simple process steps of melting, melt quenching and heat treating. However, to

obtain a permanent magnet of melt-quenched materials in bulk, a bonded magnet

should be formed by compounding a magnet powder, made from a rapidly

solidified alloy, with a resin binder. Accordingly, the magnet powder normally

accounts for at most about 80 volume % of the molded bonded magnet. Also, a

rapidly solidified alloy, formed by a melt quenching process, is magnetically

isotropic.

For these reasons, an Nd-Fe-B based rapidly solidified magnet produced

by a melt quenching process has a remanence B_r lower than that of a

magnetically anisotropic Nd-Fe-B based sintered magnet produced by a powder

metallurgical process.

As disclosed in Japanese Laid-Open Publication No. 1-7502, a technique of adding, in combination, at least one element selected from the group consisting

of Zr, Nb, Mo, Hf, Ta and W and at least one more element selected from the

group consisting of Ti, V and Cr to the material alloy effectively improves the

magnetic properties of an Nd-Fe-B based rapidly solidified magnet. When these

elements are added to the material alloy, the magnet has increased coercivity Hc

and anticorrosiveness. However, the only known effective method of improving

the remanence B_r is increasing the density of the bonded magnet. Also, where an

Nd-Fe-B based rapidly solidified magnet includes a rare earth element at about 6

at% or more, a melt spinning process, in which a melt of its material alloy is

ejected against a chill roller, has often been used in the prior art to rapidly cool

and solidify the material alloy at an increased rate.

As for an Nd-Fe-B based rapidly solidified magnet, an alternative magnet

material was proposed by R. Coehoorn et al., in J. de Phys, C8, 1998, pp.

669-670. The Coehoorn material has a composition including a rare earth

element at a relatively low mole fraction (i.e., around Nd3.sFe77.2B19, where the

subscripts are indicated in atomic percentages) and an Fe₃B type compound as

its main phase. This permanent magnet material is obtained by heating and

crystallizing an amorphous alloy that has been prepared by a melt quenching

process. Also, the crystallized material has a metastable structure in which soft

magnetic Fe₃B and hard magnetic Nd2Feι₄B phases coexist and in which crystal grains of very small sizes (typically on the order of several nanometers) are

distributed finely and uniformly as a composite of these two crystalline phases.

For that reason, a magnet made from such a material is called a "nanocomposite

magnet". It was reported that such a nanocomposite magnet has a remanence

B_r as high as about 1 T or more. But the coercivity Hcj thereof is relatively low,

i.e., from about 160 kA/m to about 240 kA/m. Accordingly, this permanent magnet

material is applicable only when the operating point of the magnet is about 1 or

more.

It has been proposed that various metal elements be added to the material

alloy of a nanocomposite magnet to improve the magnetic properties thereof.

See, for example, Japanese Laid-Open Publication No. 3-261104, United States

Patent No. 4,836,868, Japanese Laid-Open Publication No. 7-122412, PCT

International Publication No. WO 003/03403 and W. C. Chan et. al., "The Effects

of Refractory Metals on the Magnetic Properties of a -Fe R2Fei4B-type

Nanocomposites", IEEE Trans. Magn. No. 5, INTERMAG. 99, Kyongiu, Korea,

pp. 3265-3267, 1999. However, none of these proposed techniques are reliable

enough to always obtain a sufficient "characteristic value per cost". More

specifically, none of the nanocomposite magnets produced by these techniques

realize coercivity high enough to actually use it in various applications. Thus,

none of these magnets can exhibit commercially viable magnetic properties. DISCLOSURE OF INVENTION

In order to overcome the problems described above, preferred

embodiments of the present invention provide a method for producing an

iron-based alloy permanent magnet with excellent magnetic properties at a low

cost and also provide a permanent magnet that achieves a coercivity H_cj that is

high enough to actually use the magnet in various applications (e.g., H_cj≥about

550 kA/m) while maintaining a remanence B_r of about 0.80 T or more.

An iron-based rare earth alloy nanocomposite magnet according to a

preferred embodiment of the present invention has a composition represented

by the general formula (Feι_-mTm)ιoo-x-y-zQχR_yTi_z, where T is at least one element

selected from the group consisting of Co and Ni, Q is at least one element

selected from the group consisting of B and C, and R is at least one rare earth

element including substantially no La or Ce. The mole fractions x, y, z and m

preferably satisfy the inequalities of: 10 at%<x≤17 at%; 7 at%≤y<10 at%;

0.5 at%≤z≤ 6 at%; and O≤ m≤O.5, respectively. The magnet preferably

includes crystal grains of an R2Tι₄Q type compound having an average grain size

of about 20 nm to about 200 nm and a ferromagnetic iron-based boride that exists

in a grain boundary between the crystal grains of the R2T14Q type compound. The

ferromagnetic iron-based boride is preferably dispersed in the grain boundary or present in the form of a film over the grain boundary to cover the surface of the

crystal grains of the R2T14Q type compound at least partially.

In a preferred embodiment of the present invention, the mole fractions x, y

and z satisfy the inequalities of: 10 at%<x≤15 at%; 7 at%≤y≤9.3 at%; and

1.5 at%≤z≤5 at%.

In another preferred embodiment, the magnet includes crystalline phases,

including the R2Tι₄Q type compound and the ferromagnetic iron-based boride, at

about 95 vol% or more in total and amorphous phases at about 5 vol% or less.

More particularly, the magnet preferably includes the R2T14Q type

compound at about 65 vol% to about 85 vol%.

In still another preferred embodiment, the crystal grains of the R2T Q type

compound have a Ti concentration of about 2 at% or less while a Ti concentration

in the grain boundary between the crystal grains of the R2Tι Q type compound is

higher than the Ti concentration inside the crystal grains of the R2Tι Q type

compound.

Specifically, the Ti concentration in the grain boundary between the

crystal grains of the R2Tι Q type compound is preferably about 7 at% or more.

In yet another preferred embodiment, the ferromagnetic iron-based boride

has an average size of about 50 nm or less as measured along the thickness of

the grain boundary. In yet another preferred embodiment, the ferromagnetic iron-based boride

is present in the form of a film having an average thickness of about 20 nm or

less over the grain boundary between the crystal grains of the R2T14Q type

compound.

In yet another preferred embodiment, the ferromagnetic iron-based boride

is dispersed in the grain boundary between the crystal grains of the R2T14Q type

compound and has an average major axis length of about 1 nm to about 50 nm.

In yet another preferred embodiment, on an arbitrary cross section of the

magnet, the crystal grains of the R2T14Q type compound have an average size

greater than an average size of the ferromagnetic iron-based boride.

In yet another preferred embodiment, the mole fractions x and z satisfy the

inequalities of 10 at%<x≤14 at% and 0.5 at%≤z≤4 at%.

In yet another preferred embodiment, the iron-based boride includes Fe₃B

In yet another preferred embodiment, the magnet is in the shape of a thin

strip having a thickness of about 10 μm to about 300 μ.m.

In yet another preferred embodiment, the magnet has been pulverized into

powder particles. In this particular preferred embodiment, the powder particles

preferably have a mean particle size of about 30 m to about 250 μm.

In yet another preferred embodiment, the magnet has hard magnetic properties as represented by a remanence B_r of about 0.80 T or more, a

maximum energy product (BH)max of about 100 kJ/m³ or more and a coercivity

Hcj of about 480 kA/m or more.

More specifically, the magnet preferably has hard magnetic properties as

represented by a remanence B_r of about 0.85 T or more and a maximum energy

product (BH)max of about 120 kJ/m³ or more.

A bonded magnet according to a preferred embodiment of the present

invention is obtained by molding a magnet powder, comprising a powder of the

iron-based rare earth alloy magnet according to any of the preferred

embodiments of the present invention described above, with a resin binder.

BRIEF DESCRIPTION OF DRAWINGS

FIGS. 1A and 1 B schematically illustrate the structure of an iron-based rare

earth alloy nanocomposite magnet according to a preferred embodiment of the

present invention.

FIG. 2A is a cross-sectional view illustrating an overall arrangement for a

melt quenching machine for use to make a rapidly solidified alloy for the

iron-based rare earth alloy magnet according to the preferred embodiment of the

present invention; and

FIG. 2B illustrates part of the machine shown in FIG. 2A, where a melt is quenched and rapidly solidified, to a larger scale.

FIG. 3 is a graph showing the demagnetization curve of a sample

representing an example of preferred embodiments of the present invention.

FIG. 4 is a graph showing the X-ray diffraction patterns of the sample that

has been annealed.

FIG. 5 is a photograph that was taken for an example of preferred

embodiments of the present invention at a power of about 125,000 by observing

the micro metal structure of the annealed sample with a transmission electron

microscope.

FIG. 6 is a photograph showing an FIM image that was taken for the

example of preferred embodiments of the present invention by observing the

metal structure of a needle-like sample with an atom probe field-ion microscope

(APFIM).

FIG. 7 is a graph showing the cumulative concentration profiles of Nd, B

and Ti as measured in the depth direction for sample No. 2.

FIG. 8A is a graph showing relationships between the Ti concentrations

and the Nd concentrations in the regions Nos. 1 to 12 shown in FIG. 7; and

FIG. 8B is a graph showing relationships between the Ti concentrations

and the B concentrations in the regions Nos. 1 to 12 shown in FIG. 7. BEST MODE FOR CARRYING OUT THE INVENTION

An iron-based rare earth alloy magnet according to a preferred

embodiment of the present invention is a nanocomposite magnet formed by

rapidly cooling and solidifying a melt of a rare earth-iron-boron based material

alloy including Ti. This rapidly solidified alloy includes nanocrystalline phases. If

necessary, the rapidly solidified alloy is heated and further crystallized.

Generally speaking, if the mole fraction of a rare earth element R is set

lower than about 10 at%, then an R₂Feι₄B phase, contributing to hard magnetic

properties, decreases its volume percentage. In addition, an a -Fe phase

nucleates earlier than the R2Feι B phase, and easily increases its grain size

excessively. Since the α-Fe phase has high magnetization, the magnetization of

the resultant magnet increases as a whole. However, since the magnet includes

the excessively large α-Fe phase, the loop squareness of its demagnetization

curve deteriorates and its coercivity also decreases. In the prior art, one tried to

increase the coercivity by decreasing the size of each and every crystalline phase

(including the α-Fe phase that easily has an excessively large size) with the

addition of some metal element. However, the resultant coercivity was still

insufficient.

The present inventors discovered and confirmed via experiments that if an

appropriate amount of Ti is added to a material alloy including a rare earth element R at less than about 10 at% and B at about 10 at% to about 17 at%, then

crystal grains of an R2Tι₄Q type compound and then an iron-based boride can be

nucleated earlier and faster than the Fe phase in the molten alloy being cooled

and solidified with the nucleation and excessive grain growth of the α -Fe

prevented. The present inventors also discovered that by nucleating and growing

the hard magnetic phase faster and earlier than the soft magnetic phase in this

manner, nuclei of the iron-based boride with ferromagnetic properties can be

created in the grain boundary between the R2TuQ crystals. Accordingly, if the

ferromagnetic iron-based boride is grown in the crystallization process of the

material alloy so as to cover the surface of the R2Tι Q crystal grains without

increasing the interfacial energy, then the iron-based boride, grown from the

multiple nuclei, are dispersed temporarily but will soon be partially combined

together on the surface of the R2T Q crystal grains. As a result, the iron-based

boride will form a sort of film or layer in the end. In this manner, the surface of the

R2Tι Q crystal grains is covered with the film of the iron-based boride at least

partially.

Preferred embodiments of the present invention provide a nanocomposite

structure in which the R2Tι Q crystal grains serving as a hard magnetic phase are

separated from each other by a thin film of the iron-based boride serving as a soft

magnetic phase (with an average thickness of about 20 nm or less) and/or fine particles thereof (with a major-axis size of about 1 nm to about 50 nm). That is to

say, along the grain boundary of the R2T14Q crystal grains, the soft and hard

magnetic phases are magnetically coupled together through exchange

interactions. The exchange coupling between the hard and soft magnetic phases

is produced mainly in their interface. Since the soft magnetic phase is present in

the form of a film to cover the hard magnetic phase, the magnetic moment of most

of the soft phase is magnetically constrained and therefore excellent magnetic

properties are achievable. As a result, a structure that can function as an

exchange spring magnet is formed. The present inventors also discovered that

this magnetic coupling has been weakened by a feeble magnetic phase or a

non-magnetic phase existing in the grain boundary.

FIGS. 1A and 1B schematically illustrate the structure of a nanocomposite

magnet according to a preferred embodiment of the present invention.

Specifically, FIG. 1A is a cross-sectional view schematically illustrating the

R2Feι₄B and grain boundary phases in the magnet of this preferred embodiment.

As shown in FIG. 1 A, an iron-based boride (Fe-B) is present in the grain boundary

phase. FIG. 1 B is a perspective view illustrating the R₂Feι₄B phase and the

iron-based boride. As shown in FIG. 1B, the iron-based boride is finely dispersed,

and/or combined together like a film or layer, in the grain boundary, thereby

partially covering the surface of the R2Feι₄B phase. In other words, the iron-based boride existing in the grain boundary of the R2FeuB phase is partly

continued and partly discontinued.

In this manner, according to preferred embodiments of the present

invention, an iron-based boride with ferromagnetic properties is grown in the grain

boundary or sub-boundary between the R2T14Q crystal grains so as to cover the

surface of the R2T14Q crystal grains at least partially. As a result, a unique

nanocomposite magnet structure, in which a fine (or thin) iron-based boride is

present in the grain boundary or sub-boundary between the R2T14Q crystal grains,

can be obtained. The term " grain boundary" used herein refers not only to

" grain boundary" in the strict sense but also to " sub-boundary" .

The results of experiments conducted by the present inventors revealed

that particularly when the iron-based boride with ferromagnetic properties formed

a thin film over the surface of R2T14Q crystal grains in the grain boundary thereof,

a nanocomposite magnet with excellent magnetic properties could be obtained.

In preferred embodiments of the present invention, the soft magnetic phase,

existing thinly in the grain boundary, preferably accounts for about 10 vol% to

about 40 vol% of the entire magnet.

A structure like this is realized by nucleating and growing the R2T14Q crystal

grains faster and earlier than crystal grains of other soft magnetic phases through

the addition of Ti and then by precipitating the iron-based boride in the gap (or in the grain boundary) between the R2T14Q crystal grains. That is to say, it is not

until the R2Tι Q crystal grains have been sufficiently nucleated and grown that the

soft magnetic phase is allowed to grow in the grain boundary thereof.

Accordingly, the grain growth of the soft magnetic phase in the grain boundary

portion is constrained by the R2T14Q crystal grains. The crystal lattice of the soft

magnetic phase that has been formed in this manner does not match with that of

the hard magnetic phase. The structure of this nanocomposite magnet is different

from that of an Fe3B/R₂Fei₄B type nanocomposite magnet in this respect also.

Specifically, in an Fe₃B/R2Feι₄B type nanocomposite magnet, including about 2

at% to about 6 at% of rare earth element R and about 15 at% to about 20 at% of

boron (B), soft magnetic Fe₃B phase and hard magnetic R2Feι₄B phase are

formed through a phase transformation process. That is to say, the soft magnetic

Fβ3B phase precipitates before the hard magnetic R2Feι₄B phase nucleates. In

that case, the soft magnetic phase cannot be present in the film shape. Thus, the

crystal lattice of Fe₃B is observed as partially matching with that of R2Feι₄B.

In the preferred embodiments of the present invention, as R2T14Q crystal

grains nucleate and grow, Ti, which has been distributed substantially uniformly in

a molten alloy, is forced into, and condensed in, the grain boundary of the R2T14Q

crystal grains. The present inventors believe the reason is that Ti cannot exist in a

chemically stabilized state inside the R2T14Q crystal grains. According to the experimental data collected by the present inventors, the Ti condensed in the

grain boundary phase existed at a concentration of about 5 at% to about 30 at%

in the grain boundary. However, it was not clear exactly in what form Ti existed

there.

Hereinafter, an iron-based rare earth alloy magnet according to preferred

embodiments of the present invention will be described in further detail.

In various preferred embodiments of the present invention, the iron-based

rare earth alloy magnet preferably has a composition represented by the

general formula: (Feι-mTm)ιoo-x-y-zQχRyTi_z. In this formula, T is preferably at least

one element selected from the group consisting of Co and Ni, Q is preferably at

least one element selected from the group consisting of B and C, and R is

preferably at least one rare earth element including substantially no La or Ce.

The mole fractions x, y, z and m preferably satisfy the inequalities of 10 at%<x

≤17 at%, 7 at%≤y<10 at%, 0.5 at%≤z≤6 at% and O≤m≤O.5, respectively.

This alloy magnet preferably includes crystal grains of an R2T14Q type compound

and a ferromagnetic iron-based boride that have been magnetically coupled

together through exchange interaction. The ferromagnetic iron-based boride has

preferably grown to form a film in the grain boundary between the R2T14Q crystal

grains and cover the surface of the R2Tι₄Q crystal grains as a whole. The iron-based rare earth alloy magnet according to this preferred

embodiment of the present invention includes a rare earth element at less than

about 10 at%. However, the iron-based boride with ferromagnetic properties has

precipitated in the grain boundary of the main phase. Accordingly, the magnet

can exhibit comparable (or even increased) magnetization (or remanence B_r)

and improved loop squareness of the demagnetization curve compared to a

magnet that does not contain Ti.

In the iron-based rare earth alloy magnet according to the preferred

embodiment of the present invention, the crystal grains of its main phase with

ferromagnetic properties are thinly covered with the iron-based boride as a soft

magnetic phase. Accordingly, the crystal grains of these two main phases are

magnetically coupled together through exchange interaction. Consequently, the

alloy as a whole can exhibit excellent loop squareness in its demagnetization

curve.

The iron-based rare earth alloy magnet according to this preferred

embodiment of the present invention preferably includes iron-based borides and

α-Fe that have saturation magnetizations approximately equal to, or even higher

than, that of the R2T14Q (typically, Nd₂Feι₄B) phase. Examples of the iron-based

borides include Fe₃B (with a saturation magnetization of about 1.5 T) and Fe₂3B6 (with a saturation magnetization of about 1.6 T). In this case, the Nd2Feι₄B phase

has a saturation magnetization of about 1.6 T and the α -Fe phase has a

saturation magnetization of about 2.1 T.

Normally, where the mole fraction x of B is greater than about 10 at% and

the mole fraction y of the rare earth element R is about 6 at% to about 8 at%,

R2Fβ23B3 is produced unless Ti is added. However, even when a material alloy

with such a composition is used, the addition of Ti can produce R2Feι₄B, Fe23Bδ

and Fe₃B, not R2Fe23B₃, in the preferred embodiment of the present invention.

These iron-based borides contribute to increasing the magnetization. It should be

noted that "FesB" herein includes Fe3.δB, which is hard to distinguish from Fe₃B.

In various preferred embodiments of the present invention, the rapidly

solidified alloy has either a structure in which almost no α -Fe phase with an

excessively large grain size has precipitated but an R2Tι Q type phase with a

very small grain size exists instead or a structure in which an R2Tι₄Q type phase

with a very small grain size and an amorphous phase coexist. As used herein,

the term "amorphous phase" means not only a phase in which the atomic

arrangement is sufficiently disordered, but also a phase including embryos for

crystallization, extremely small crystalline regions (size: several nanometers or

less), and/or atomic clusters. More specifically, the term "amorphous phase" refers to any phase having a crystal structure that cannot be defined by X-ray

diffraction analysis or TEM observation. Stated otherwise, any phase having a

crystal structure clearly identifiable by X-ray diffraction analysis or TEM

observation will be herein referred to as a "crystalline phase".

In the prior art, if a molten alloy with a composition similar to that of the

preferred embodiments of the present invention (i.e., a composition including all

the elements included in the inventive composition but Ti) is cooled relatively

slowly, the resultant alloy will have a structure in which a lot of α-Fe phase has

grown coarsely. Thus, when the alloy is heated and crystallized after that, the

-Fe phase will increase its grain size excessively. Once soft magnetic phases,

including the α-Fe phase, have grown too much, the magnetic properties of the

alloy deteriorate significantly, thus making it virtually impossible to produce a

quality permanent magnet from such an alloy.

Only when Ti is added to the material alloy, the hard magnetic phase will

nucleate and grow faster and earlier than any other phase and then an

iron-based boride with ferromagnetic properties will precipitate in the grain

boundary between the crystal grains of the main phase. Thereafter, the

iron-based boride precipitated will soon be partially combined together to form a

continuous film. As a result, a structure in which the surface of the crystal grains of the main phase is thinly covered with that film is formed.

If any of the other metal elements (e.g., Nb, V, Cr, etc.), excluding Ti, is

added, the grain growth of the α-Fe phase advances remarkably in a relatively

high temperature range in which the α -Fe phase grows rapidly, and the

magnetization direction of the α-Fe phase cannot be effectively constrained by

the exchange coupling between the α-Fe and hard magnetic phases. As a result,

the demagnetization curve will have greatly deteriorated loop squareness. It

should be noted that even if Nb, Mo or W is added instead of Ti, good hard

magnetic properties, including superior loop squareness of the demagnetization

curve, are achievable by thermally annealing the alloy in a relatively low

temperature range where no α-Fe phase precipitates. In an alloy that has been

annealed at such a low temperature, however, R2Feι₄B crystalline particles would

be dispersed in non-magnetic amorphous phases and the alloy does not have

large remanence B_r. Also, if the alloy is annealed at an even higher temperature,

then the α-Fe phase nucleates and grows out of the amorphous phases. Unlike

the situation where Ti is added, the α-Fe phase rapidly grows and increases its

grain size excessively after its nucleation. As a result, the magnetization direction

of the α-Fe phase cannot be effectively constrained anymore by the exchange

coupling between the α-Fe and hard magnetic phases, and the loop squareness of the demagnetization curve deteriorates considerably.

On the other hand, where V or Cr is added instead of Ti, the magnetic

moment of these additive metal elements is coupled anti-ferromagnetically to

the magnetic moment of Fe to form a solid solution, thus decreasing the

magnetization significantly.

In contrast, where Ti is added to the material alloy, the crystallization

kinetics of the α-Fe phase is slowed down, i.e., it takes a longer time for the α

-Fe phase to nucleate and grow. Thus, the present inventors believe that the

Nd2Fei4B phase starts to nucleate and grow before the α-Fe phase has grown

coarsely. For that reason, where Ti is added, crystal grains of the Nd2Fe B

phase can be grown sufficiently and distributed uniformly before the α-Fe phase

has grown too much.

Accordingly, only when Ti is added, the crystal grain coarsening of the

α -Fe phase is minimized as intended, and therefore, iron-based borides with

ferromagnetic properties can be obtained. Furthermore, Ti, as well as B and C,

plays an important role as an element that delays the crystallization of Fe initial

crystals (i.e., y -Fe that will be transformed into α -Fe) and thereby facilitates

the production of a supercooled liquid during the melt quenching process. Thus,

even if a melt of the material alloy including Ti is rapidly cooled and solidified at a relatively low cooling rate of about 10² °C/sec to about 10⁵ °C/sec, a rapidly

solidified alloy, in which the α -Fe phase with an excessively large grain size has

not precipitated but which includes the nanocrystalline R2Fe ₄B phase at about 60

vol% or more (and sometimes iron-based borides as well), can be obtained.

Preferred composition

Q may be either B (boron) only or a combination of B and C (carbon).

The ratio of C to Q is preferably about 0.25 or less.

If the mole fraction x of Q is about 10 at% or less, then it is difficult to

make the desired rapidly solidified alloy, in which the nanocrystalline R2Feι₄B and

amorphous phases coexist, at the low cooling rate of about 10² °C/sec to about

10⁴ °C/sec. Also, even if the rapidly solidified alloy is annealed after that, the

resultant Hcj of the alloy will be less than about 400 kA/m. In addition, a strip

casting process, which is one of the most cost-effective techniques among

various rapid cooling methods, cannot be adopted in that case, and the price of

the resultant permanent magnet product increases unintentionally. On the other

hand, if the mole fraction x of Q exceeds about 17 at%, then the iron-based boride

will start to nucleate almost simultaneously with the R2Feι₄B phase, and will grow

excessively in the end. As a result, the desired nanocomposite structure, in which

the iron-based boride phase is distributed uniformly in, or present in the form of a film over, the grain boundary of the R2Feι₄B phase, cannot be obtained, and the

resultant magnetic properties deteriorate.

In view of these considerations, the mole fraction x of Q is preferably

greater than about 10 at% and equal to or less than about 17 at%. A more

preferable upper limit of x is about 16 at% and an even more preferable upper

limit of x is about 15 at%.

The (atomic) ratio p of C to Q is preferably from about 0 to about 0.25. To

achieve the effects expected from the additive C, the C ratio p is preferably

equal to or greater than about 0.01. The reason is as follows. If p is much

smaller than about 0.01 , then almost no expected effects are achievable even if

C is added. On the other hand, if p is far greater than about 0.25, then the

volume percentage of the α-Fe phase produced increases too much, thereby

causing deterioration of the resultant magnetic properties. The lower limit of the

ratio p is preferably about 0.02, while the upper limit thereof is preferably about

0.20. More preferably, the ratio p is from about 0.02 to about 0.17.

R is at least one element selected from the rare earth elements (including

yttrium). Preferably, R includes substantially no La or Ce. This is because if La or

Ce is included, R (typically Nd) included in the R2Fei4B phase is replaced with La

or Ce, thus decreasing the coercivity and deteriorating the loop squareness of the demagnetization curve. However, the magnetic properties will not be affected so

seriously if a very small percentage (i.e., about 0.5 at% or less) of La or Ce exists

as an inevitably contained impurity. Therefore, the phrase "substantially no La

(Ce)" or "substantially excluding La (Ce)" herein means that the content of La

(Ce) is about 0.5 at% or less. More specifically, R preferably includes Pr or Nd as

an indispensable element, a portion of which may be replaced with Dy and/or Tb.

If the mole fraction y of R is less than about 6 at%, then fine grains with the

nanocrystalline R₂Fei4B structure, which is needed for realizing the coercivity, do

not crystallize sufficiently and the desired high coercivity Hcj of about 480 kA/m or

more cannot be obtained. On the other hand, if the mole fraction y of R is equal to

or greater than about 10 at%, then the percentage of the iron-based borides with

ferromagnetic properties and α -Fe phase decreases but that of the R-rich

non-magnetic phases increases instead. As a result, the intended nanocompo¬

site structure cannot be formed and the magnetization drops. For these reasons,

the mole fraction y of the rare earth element R is preferably equal to or greater

than about 6 at% but less than about 10 at% (e.g., from about 7 at% to about 9.5

at%), more preferably from about 7.5 at% to about 9.3 at%, and even more pref¬

erably from about 8.0 at% to about 9.0 at%. Although the mole fraction y of the

rare earth element R is low in the preferred embodiments of the present invention,

the additive Ti enables the R₂FeuB phase to nucleate and grow faster and earlier than any other phase. Accordingly, R included in the molten alloy can be used

effectively to produce the R₂Fei4B phase and the concentration of R will be low in

the grain boundary portion of the R2FeuB phase. As a result, the concentration of

R in the grain boundary phase becomes about 0.5 at% or less, which is much

lower than the R concentration of about 11 at% in the hard magnetic phase. In

this manner, R can be used effectively to form the hard magnetic phase (e.g., the

R2Fe B phase) in the preferred embodiments of the present invention. Therefore,

even though the R mole fraction y is less than about 10 at% and the hard mag¬

netic phase (e.g., the R₂Fei4B phase) accounts for about 65 at% to about 85 at%

of the entire alloy, that hard magnetic phase still can be coupled magnetically with

the soft magnetic phase existing in the grain boundary through exchange interac¬

tion. As a result, excellent hard magnetic properties are achieved. It should be

noted that the volume percentage of each constituent phase such as the R2Fei4B

phase is herein measured by Mossbauer spectroscopy.

To achieve the effects described above, Ti is indispensable. The additive

Ti increases the coercivity Hcj, remanence B_r and maximum energy product

(BH)max and improves the loop squareness of the demagnetization curve.

If the mole fraction z of Ti is less than about 0.5 at%, then those effects

expected from the additive Ti cannot be achieved fully. Nevertheless, if the mole fraction z of Ti exceeds about 6 at%, then the volume percentage of the

non-magnetic phases, remaining in the alloy even after the alloy has been heated

and crystallized, increases and the remanence B_r likely drops. Also, if the mole

fraction z of Ti exceeds about 6 at%, ΗB2 is produced in the molten alloy, thus

making it difficult to carry out the melt quenching process as intended. In view of

these considerations, the mole fraction z of Ti is preferably from about 0.5 at% to

about 6 at%. The lower limit of a more preferable z range is about 1.0 at% and

the upper limit thereof is about 5 at%. The upper limit of an even more preferable

z range is about 4 at%.

Also, the higher the mole fraction x of Q, the more likely the amorphous

phases, including an excessive percentage of Q (e.g., boron), are formed.

Accordingly, the mole fraction z of Ti is preferably set higher because of this

reason also. Ti has a strong affinity for B and is condensed in the film-like grain

boundary phase. However, if the ratio of the mole fraction z of Ti to the mole

fraction x of B is too high, then Ti will not be present in the Fe₃B grain boundary

phase anymore but will be incorporated into the R2FeuB compound, thus

possibly decreasing the magnetization. Nevertheless, if the z/x ratio is too low,

then non-magnetic B-rich amorphous phases will be produced profusely. The

present inventors discovered and confirmed via experiments that the mole fractions x and z are preferably controlled to satisfy the inequality of 0.05≤z/x≤

0.4, more preferably to satisfy the inequality of 0.1≤z/x≤0.35 and even more

preferably to satisfy the inequality of 0.13≤z/x≤0.3.

The balance of the material alloy, other than the elements B, C, R and Ti,

may be Fe alone. Alternatively, at least one transition metal element T selected

from the group consisting of Co and Ni may be substituted for a portion of Fe,

because the desired hard magnetic properties are achievable in that case also.

However, if more than about 50% of Fe is replaced with T, then a high remanence

B_r of about 0.7 T or more cannot be obtained. For that reason, the percentage of

Fe replaced is preferably from about 0% to about 50%. Also, by substituting Co

for a portion of Fe, the loop squareness of the demagnetization curve improves

and the Curie temperature of the R2Fei4B phase increases, thus increasing the

thermal resistance of the alloy. The percentage of Fe replaceable with Co is

preferably from about 0.5% to about 40%.

To achieve various desired advantages and effects, metal element(s) M

may be added at a mole fraction of about 0 at% to about 10 at%. M is at least

one element selected from the group consisting of Al, Si, V, Cr, Mn, Cu, Zn, Ga,

Zr, Nb, Mo, Hf, Ta, W, Pt, Pb, Au and Ag.

Hereinafter, specific preferred embodiments of the present invention will be described with reference to the accompanying drawings.

Melt quenching machine

In this preferred embodiment, a material alloy is prepared using a melt

quenching machine such as that shown in FIGS. 2A and 2B. The alloy

preparation process is performed within an inert atmosphere to prevent the

material alloy, which includes rare earth element R and Fe that are easily

oxidizable, from being oxidized. The inert gas may be either a rare gas of

helium or argon, for example, or nitrogen. A rare gas of helium or argon is

preferred to nitrogen, because nitrogen reacts with the rare earth element R

relatively easily.

The machine shown in FIG. 2A includes a material alloy melting chamber

1 and a material alloy quenching chamber 2, in which a vacuum or an inert

atmosphere is created at an adjustable pressure. Specifically, FIG. 2A illustrates

an overall arrangement of the machine, while FIG. 2B illustrates a part of the

machine to a larger scale.

As shown in FIG. 2A, the melting chamber 1 includes a melting crucible 3,

a melt reservoir 4 with a teeming nozzle 5 at the bottom and an airtight mixed

material feeder 8. A material alloy 20, which has been prepared to have a desired magnet alloy composition and supplied from the feeder 8, is melted in the

melting crucible 3 at an elevated temperature. A melt 21 of the material alloy 20 is

poured into the reservoir 4, which is provided with a heater (not shown) for holding

the temperature of the melt 21 teemed therefrom at a predetermined level.

The quenching chamber 2 includes a rotating chill roller 7 for rapidly

cooling and solidifying the melt 21 that has been poured through the teeming

nozzle 5.

In this machine, the atmosphere and pressure inside the melting and

quenching chambers 1 and 2 are controlled to prescribed ranges. For that

purpose, atmospheric gas inlet ports 1 b, 2b and 8b and outlet ports 1a, 2a and

8a are provided at appropriate positions of the machine. In particular, the gas

outlet port 2a is connected to a pump to control the absolute pressure inside the

quenching chamber 2 within a range from about 30 kPa to around the

atmospheric pressure.

The melting crucible 3 may be tilted to a desired angle to pour the melt 21

through a funnel 6 into the reservoir 4. The melt 21 in the reservoir 4 is heated

by the heater (not shown).

The teeming nozzle 5 of the reservoir 4 is positioned on the boundary wall between the melting and quenching chambers 1 and 2 to pour the melt 21

from the reservoir 4 onto the surface of the chill roller 7, which is located under

the nozzle 5. The orifice diameter of the nozzle 5 may be from about 0.5 mm to

about 2.0 mm, for example. If the viscosity of the melt 21 is high, then the melt 21

cannot flow through the nozzle 5 easily. In this preferred embodiment, however,

the pressure inside the quenching chamber 2 is kept lower than the pressure

inside the melting chamber 1. Accordingly, there exists an appropriate pressure

difference between the melting and quenching chambers 1 and 2, and the melt 21

can be teemed smoothly.

To achieve a good thermal conductivity, the chill roller 7 may be made of

Al alloy, Cu alloy, carbon steel, brass, W, Mo or bronze. However, the roller 7 is

preferably made of Cu, because Cu realizes a sufficient mechanical strength at

a reasonable cost. The diameter of the roller 7 may be about 300 mm to about

500 mm, for instance. The water-cooling capability of a water cooler provided

inside the roller 7 is calculated and controlled based on the latent heat of

solidification and the volume of the melt teemed per unit time.

The machine shown in FIGS. 2A and 2B can rapidly solidify about 10 kg of

material alloy in about 10 to about 20 minutes, for example. The rapidly

solidified alloy obtained in this manner is in the form of a thin strip (or alloy ribbon) 22 with a thickness of about 10 μm to about 300 m and a width of

about 2 mm to about 3 mm.

Melt quenching process

First, the melt 21 of the material alloy, which is represented by the general

formula described above, is prepared and stored in the reservoir 4 of the

melting chamber 1 shown in FIG. 2A. Next, the melt 21 is poured through the

teeming nozzle 5 onto the water-cooled chill roller 7 to come into contact with,

and be rapidly cooled and solidified by, the roller 7 within a reduced-pressure Ar

atmosphere. In this case, an appropriate rapid solidification technique that

makes the cooling rate controllable precisely needs to be adopted.

In this preferred embodiment, the resultant magnet can easily have its

desired microstructure by nucleating the R2T14Q type phase as a hard magnetic

phase in a high-temperature range where crystal nuclei are created very often

while the molten alloy is being cooled. Accordingly, the R2T14Q type phase can be

included at about 60 vol% or more in the as-cast (i.e., yet to be annealed) rapidly

solidified alloy. To obtain such a structure, the molten alloy is preferably

quenched at a rate of about 1 X 10² °C/sec to about 1 X 10⁸ °C/sec, more

preferably about 1 X 10² °C/sec to about 1 10⁶ °C/sec. An interval during which the molten alloy 21 is quenched by the chill roller

7 is equivalent to an interval between a point in time the molten alloy 21 reaches

the circumference of the rotating chill roller 7 and a point in time the quenched

alloy 22 leaves the roller 7. In the meantime, the alloy has its temperature

decreased to become a supercooled liquid. Thereafter, the supercooled alloy

leaves the roller 7 and travels within the inert atmosphere. While the thin-strip

alloy is traveling, the alloy has its heat dissipated into the atmospheric gas. As a

result, the temperature of the alloy further drops. According to the preferred

embodiment of the present invention, the pressure of the atmospheric gas is

controlled to a range of about 30 kPa to around the atmospheric pressure. Thus,

the heat of the alloy can be dissipated into the atmospheric gas even more

effectively, and the R2T14Q type phase such as an Nd2Feι₄B phase can be

nucleated and grown finely and uniformly in the alloy. It should be noted that

unless an appropriate amount of Ti has been added to the material alloy, the α

-Fe phase nucleates and grows faster and earlier in the rapidly solidified alloy,

thus deteriorating the resultant magnet properties.

In this preferred embodiment, the surface velocity of the roller 7 is adjusted

to a range of about 10 m/sec to about 30 m/sec and the pressure of the

atmospheric gas is set equal to about 30 kPa or more to increase the secondary cooling effects caused by the atmospheric gas. In this manner, a rapidly

solidified alloy, including about 60 vol% or more of R2T14Q type phase with an

average grain size of as small as about 50 nm or less, is prepared.

According to various preferred embodiments of the present invention, the

technique of rapidly cooling the melt is not limited to the single roller melt-spinning

method described above. Examples of other applicable techniques include twin

roller method, gas atomization method, strip casting method that requires no flow

rate control using nozzle or orifice, and rapid cooling technique utilizing the roller

and atomization methods in combination.

Among these rapid cooling techniques, the strip casting method results in a

relatively low cooling rate of about 10² °C/sec to about 10⁵ °C/sec. According to

this preferred embodiment, by adding an appropriate volume of Ti to the material

alloy, a rapidly solidified alloy, most of which has a structure including no Fe initial

crystals, can be obtained even by the strip casting method. The process cost of

the strip casting method can be about half or less of any other rapid cooling

method. Accordingly, in preparing a large quantity of rapidly solidified alloy, the

strip casting method is much more effective than the single roller method, and is

suitably applicable for use in mass production. However, if no Ti is added to the

material alloy or if Mn, Mo, Ta and/or W are/is added thereto instead of Ti, then a metal structure including a lot of Fe initial crystals will be produced even in the

rapidly solidified alloy prepared by the strip casting method. Consequently, the

desired microstructure cannot be obtained.

Heat treatment

In this preferred embodiment, heat treatment (annealing) is conducted

within an argon atmosphere. Preferably, the alloy is heated at a temperature rise

rate of about 5 °C/sec to about 20 °C/sec, held at a temperature of about 550 °C to

about 850 °C for a period of time from about 30 seconds to about 20 minutes and

then cooled to around room temperature. This heat treatment results in

nucleation and/or crystal growth of metastable phases in a remaining amorphous

phase, thus forming a nanocomposite nanocrystalline structure. According to the

preferred embodiment of the present invention, the nanocrystalline R₂Feι₄B

phase (e.g., Nd2Feι₄B phase) already accounts for about 60 volume % or more of

the as-cast rapidly solidified alloy yet to be annealed. Thus, when the heat

treatment is conducted under the conditions specified above, soft magnetic

phases other than the nanocrystalline Nd2FeuB phase (such as α-Fe and other

crystalline phases) will not increase their sizes too much and will be distributed

finely and uniformly in the grain boundary between the nanocrystalline Nd2Feι B

grains. When the heat treatment is finished, the nanocrystalline R2Fe B (e.g., Nd2Fei4B) phase will account for about 65 vol% to about 85 vol% of the annealed

alloy.

If the heat treatment temperature is lower than about 550 °C, then a lot of

amorphous phases may remain even after the heat treatment and the resultant

coercivity may not reach the desired level depending on the conditions of the

rapid cooling process. On the other hand, if the heat treatment temperature

exceeds about 850 °C, the grain growth of the respective constituent phases may

advance too much, thus decreasing the remanence B_r and deteriorating the loop

squareness of the demagnetization curve. For these reasons, the heat treatment

temperature is preferably about 550 °C to about 850 °C, more preferably about

570 °C to about 820 °C.

In this preferred embodiment of the present invention, a sufficient amount

of Nd₂Feι₄B phase has crystallized uniformly and finely in the rapidly solidified

alloy. Accordingly, even if the rapidly solidified alloy is not annealed, the solidified

alloy itself can exhibit good enough magnet properties. That is to say, the heat

treatment for crystallization is not indispensable for the present invention.

However, to further improve the magnet properties, the heat treatment is

preferably conducted. In addition, even though the heat treatment is carried out at

lower temperatures than the conventional process, the magnet properties still can be improved sufficiently.

To prevent the alloy from being oxidized, the heat treatment is preferably

conducted within an inert gas (e.g., Ar or N2 gas) atmosphere having a pressure

of about 50 kPa or less. Alternatively, the heat treatment may also be carried out

in a vacuum of about 1.0 kPa or less.

Before the heat treatment, the rapidly solidified alloy may include

metastable phases such as Fe₃B, Fe₂3B6 and R₂Fe23B3 phases in addition to the

R₂Fei4B and amorphous phases. In that case, when the heat treatment is

finished, the R₂Fe23B₃ phase will have disappeared. Instead, crystal grains of an

iron-based boride (e.g., Fe23Be), showing a saturation magnetization

approximately equal to, or even higher than, that of the R2Fei4B phase, or α-Fe

phase can be grown.

The resultant magnet alloy that has gone through this heat treatment

includes an R2T14Q type phase (e.g., R2FeuB phase) at about 65 vol% to about

85 vol%. More specifically, if the mole fraction y of R is about 9 at%, then the

R₂Tι₄Q type phase is about 75 vol% of the entire alloy. But if the mole fraction y

of R is decreased to about 8 at%, then the R2Tι₄Q type phase accounts for about

68 vol% of the entire alloy. On the other hand, the magnet also includes soft

magnetic phases at about 10 vol% to about 35 vol%. Also, the crystalline phases, including the R2T14Q type compound and the

ferromagnetic iron-based boride, account for about 95 vol% or more in total, while

the amorphous phases account for about 5 vol% or less of the entire alloy.

In the preferred embodiments of the present invention, even if the soft

magnetic phases such as the iron-based boride remain in the resultant annealed

alloy, those soft magnetic phases are present in the shape of a thin film that

covers the hard magnetic phases. Accordingly, excellent magnetic properties are

still achievable because the soft and hard magnetic phases are magnetically

coupled together through exchange interaction.

According to the preferred embodiments of the present invention, the grain

boundary phases are mostly made of an iron-based boride (e.g., Fe₃B) with

ferromagnetic properties and also include α -Fe phase with ferromagnetic

properties and other additional phases. More specifically, the iron-based boride

accounts for about 70 vol% or more of the grain boundary phases. However,

almost no rare earth elements R such as Nd are present in the grain boundary

phases. Instead, the rare earth elements R are effectively used for producing the

hard magnetic phase. A nanocrystalline structure like this cannot be obtained

unless an appropriate amount of Ti is added to a composition having an R mole

fraction y of less than about 10 at% and a Q mole fraction x of greater than about 10 at%. Stated otherwise, if any other metal element were added instead of Ti,

the resultant grain boundary phases, if any, would be amorphous phases with low

magnetization. Thus, it would be difficult to make a nanocomposite magnet with

good properties from such an alloy. Also, even if Ti is added to a material alloy

having a Q mole fraction x of about 10 at% or less, no soft magnetic phases with

high magnetization will be formed in the grain boundary. Thus, the resultant

magnet cannot be a nanocomposite magnet that exhibits excellent magnet

properties through the exchange coupling of its composite phases.

After the heat treatment, the R2T14Q type phase needs to have an average

crystal grain size of about 300 nm or less, which is a single magnetic domain size.

The average grain size of the R₂Tι₄Q type phase is preferably about 20 nm to

about 200 nm, more preferably about 20 nm to about 100 nm.

On the other hand, if the thin film of the iron-based boride has an average

thickness of greater than about 50 nm, then the exchange interaction among the

respective constituent phases weakens, thus deteriorating the loop squareness of

the demagnetization curve and decreasing (BH)maχ. This is why the average size

of the iron-based boride phase (i.e., the average thickness of the film) as

measured in the thickness direction of the grain boundary is preferably about 50

nm or less, more preferably about 30 nm or less and even more preferably about 20 nm or less.

It should be noted that the thin strip of the rapidly solidified alloy may be

coarsely cut or pulverized before subjected to the heat treatment. After the heat

treatment, the resultant magnetic alloy is finely pulverized to obtain a magnet

powder. Then, various types of bonded magnets can be made from this magnet

powder by performing known process steps on the powder. In making a bonded

magnet, the magnet powder of the iron-based rare earth alloy is compounded with

an epoxy or nylon resin binder and then molded into a desired shape. At this time,

a magnet powder of any other type (e.g., an Sm-Fe-N type magnet powder or

hard ferrite magnet powder) may be mixed with the nanocomposite magnet

powder.

Using the resin-resultant bonded magnet, motors, actuators and other

rotating machines can be produced.

When the magnet powder according to this preferred embodiment of the

present invention is used for an injection-molded bonded magnet, the magnet

powder is preferably pulverized to- have a mean particle size of about 200 Min or

less, more preferably about 30 m to about 150 Mm. On the other hand, if this

magnet powder is used for a compacted bonded magnet, the magnet powder is

preferably pulverized to have a mean particle size of about 300 Mm or less, more preferably about 30 M m to about 250 M m and even more preferably about 50 M

m to about 200 Mm with a bimodal size distribution.

It should be noted that if the magnet powder obtained in this manner is

subjected to a surface treatment (e.g., coupling treatment, conversion coating or

plating), then the powder for a bonded magnet can have its moldability

improved no- matter how the powder is molded. In addition, the resultant

bonded magnet can have increased anticorrosiveness and thermal resistance.

Alternatively, after a bonded magnet has been once formed by molding the

magnet powder into a desired shape, the surface of the magnet may also be

treated, e.g., covered with a plastic or conversion coating or plated. This is

because the anticorrosiveness and thermal resistance of the bonded magnet

can also be increased as in the situation where the magnet powder is subjected

to the surface treatment.

Examples

EXAMPLE 1

The respective materials B, Fe, Ti and Nd with purities of about 99.5% or

more were weighed so as to have an alloy composition represented by

Nd9Fe7e.7B10.3Ti2 (where the subscripts are indicated in atomic percentages) and a total weight of about 30 g. Then, the mixture was injected into a crucible of quartz.

The quartz crucible had an orifice with a diameter of about 0.8 mm at the

bottom. Accordingly, the alloy including these materials was melted in the quartz

crucible so as to be a melt of the material alloy, which was then poured down

through the orifice. The material alloy was melted by an induction heating method

within an argon atmosphere at a pressure of about 35 kPa. In this specific

example of preferred embodiments of the present invention, the temperature of

the melt was set to about 1500 °C.

The surface of the melt was pressurized with Ar gas at about 26.7 kPa,

thereby ejecting the melt against the outer circumference of a copper chill roller,

which was located about 0.7 mm under the orifice. The roller was rotated at a

high velocity while being cooled inside so that the outer circumference

temperature thereof would be held around room temperature. Accordingly, the

melt, which had been poured down through the orifice, came into contact with the

surface of the chill roller to have its heat dissipated therefrom while being forced to

rapidly move on the rotating chill roller. The melt was continuously expelled

through the orifice onto the surface of the roller. Thus, the rapidly solidified alloy

was in the shape of an elongated thin strip (or ribbon) with a width of about 2 mm

to about 3 mm and a thickness of about 20 M m to about 50 M m. In the (single) roller method adopted in this example, the cooling rate is

determined by the circumference velocity of the roller and the weight of the melt

poured per unit time, which depends on the diameter (or cross-sectional area) of

the orifice and the pressure on the melt. In the present example, the feeding rate

was about 0.5 kg/min to about 1 kg/min and the roller surface velocity was about

20 m/sec.

The rapidly solidified alloy structure that had been obtained by this rapid

cooling process was analyzed by a CuKα characteristic X-ray. As a result,

diffraction peaks representing Nd2Feι B phase were barely recognizable in the

halo pattern. Thus, the present inventors confirmed that a nanocrystalline

Nd2Feι₄B phase was present in the amorphous phases of the rapidly solidified

alloy.

Next, this rapidly solidified alloy was annealed within an argon gas.

Specifically, the rapidly solidified alloy was held at about 660 °C for about 10

minutes and then cooled to room temperature. Thereafter, the magnetic

properties of the annealed alloy were measured with a vibrating sample

magnetometer (VSM). The results are shown in the following Table 1 : Table 1

FIG. 3 shows the demagnetization curve of this sample.

When the annealed alloy structure was analyzed by a CuKα characteristic

X-ray, the halo pattern had disappeared but diffraction peaks representing

Nd2Fei4B, Fβ23B6 and α-Fe phases were observed. FIG. 4 shows the powder

X-ray diffraction patterns of the annealed alloy.

Next, the annealed micro metal structure was analyzed with a transmission

electron microscope (TEM). As a result, crystal grains with an average grain size

of about 150 nm and fine crystal grains with an average grain size of about 20 nm

were observed. The fine crystal grains were present in the grain boundary

between the former crystal grains. FIG. 5 is a TEM photograph showing a

dark-field image of the annealed alloy at a power of about 125,000.

Next, the alloy compositions of the crystal grains that had been observed

with the TEM were analyzed by a TEM-EDX. As a result, the crystal grains with the average grain size of about 150 nm were identified as Nd2Fei4B. However,

the phase of the fine crystal grains with the average grain size of about 20 nm,

existing in the grain boundary between the Nd2Fei4B crystal grains, was not

identifiable. Thus, it can be seen that the iron-based boride was either dispersed

as very small particles in the grain boundary or present in the form of a thin film (or

layer).

Next, the annealed alloy was polished mechanically and processed into a

prismatic rod-like sample. Furthermore, the end of this rod-like sample was

sharpened into a needle shape by an electrolytic polishing technique. Then, the

metal structure of this needle-like sample was analyzed with an atom probe

field-ion microscope (APFIM). As a result, the present inventors identified not

only an Nd2Fei4B phase as main phase but also an iron-based boride and a very

small amount of Fe. The present inventors also discovered that the concentration

of Ti in the iron-based boride was about three times as high as that of Ti in the

Nd₂Fei4B phase.

The cumulative numbers of ions counted by the APFIM analysis were

plotted. When the total number of ions plotted reached about 12,700, the APFIM

analysis was stopped to take an FIM image of the end of the needle-like sample.

FIG. 6 shows this FIM image. The center of the image shown in FIG. 6 corresponds to a probe hole. The

FIM image obtained shows that bright island-like regions are dispersed in the dark

region. Considering the conditions under which this image was taken, the dark

matrix may be regarded as representing the Nd₂Feι B phase while the bright

island-like regions may be regarded as representing the iron-based boride phase.

That is to say, it can be seen that the boride phase was finely dispersed in the

annealed alloy. However, the composition of the iron-based boride phase of the

very small size was not analyzable with the TEM.

To detect the three-dimensional distribution of the iron-based boride more

accurately, a number of FIM images of the needle-like sample need to be taken

successively at relatively short intervals in the depth direction thereof. Then, the

FIM images obtained need to be processed so as to be arranged in the depth

direction.

The TEM and FIM images both showed that the iron-based boride phase

had a cross-sectional size of about 10 nm. Accordingly, if the iron-based boride

had been present inside the Nd2Feι₄B crystal grains, then its phase should have

been identified with the TEM. However, no such phases were identifiable

according to the results of experiments conducted by the present inventors.

These results showed that the iron-based boride was not present inside the Nd2Feι B crystal grains but either finely dispersed in the grain boundary between

the Nd2Feι₄B crystal grains or present in the form of a partially continuous film

covering the Nd2Feι₄B crystal grains.

If the iron-based boride in the shape of such a porous film is observed as a

TEM or FIM image, then the iron-based boride will be observed as mutually

separated, fine crystal grains having a size of about 1 nm to about 20 nm on an

arbitrary cross section of the Nd2Fei4B crystal grains and the iron-based boride.

In the preferred embodiments of the present invention, while the molten

alloy is being rapidly solidified, the R₂Feι₄B phase crystallizes first, and the

iron-based boride phase crystallizes next. Accordingly, the present inventors

believe that the iron-based boride phase will nucleate and grow from the surface

of R2Fei4B crystal grains, which are dispersed in the amorphous matrix, as its

non-uniform nuclei. The present inventors also believe that the iron-based boride

phase, which has nucleated on the surface of the R2Fei4B crystal grains, would

grow in such a manner as to cover the surface of the R2Fe B crystal grains to

prevent increase in interface energy. As a result, portions of the iron-based

boride phase would be combined together to form a film that covers the R₂Fe B

crystal grains at least partially.

In this manner, preferred embodiments of the present invention provide a novel nanocomposite magnet structure in which R₂Feι₄B crystal grains with hard

magnetic properties are covered with a film of an iron-based boride with soft

magnetic properties. The present inventors believe that this magnet exhibit

excellent magnetic properties due to this unique structure.

EXAMPLE 2

For each of samples Nos. 1 to 8 shown in the following Table 2, the

respective materials B, C, Fe, Ti, V, Cr and Nd with purities of about 99.5% or

more were weighed so that each sample had a total weight of about 30 g. Then,

the mixture was injected into a crucible of quartz. In Table 2, "bal" means the

balance.

Table 2

The quartz crucible had an orifice with a diameter of about 0.8 mm at the

bottom. Accordingly, the alloy including these materials was melted in the quartz

crucible so as to be a melt of the alloy, which was then poured down through the

orifice. The material alloy was melted by an induction heating method within an

argon atmosphere. In this second specific example, the temperature of the melt

was set to about 1400 °C. The surface of the melt was pressurized with Ar gas at about 30 kPa,

thereby ejecting the melt against the outer circumference of a copper chill roller,

which was located under the orifice. In the present example, the feeding rate was

about 0.4 kg/min and the roller surface velocity was about 20 m/sec. In this

manner, a thin-strip rapidly solidified alloy with a width of about 1.0 mm and a

thickness of about 50 Mm was obtained.

This rapidly solidified alloy structure was analyzed by a powder XRD

analysis. As a result, the present inventors confirmed that the rapidly solidified

alloy was made up of amorphous phases.

Next, the rapidly solidified alloy was cut to a length of about 20 mm and

then annealed in Ar gas. This annealing process was conducted by holding each

sample at the temperature shown in Table 2 for about 10 minutes.

The annealed alloy structure was subjected to a powder XRD analysis. As

a result, diffraction peaks representing Nd2Feι₄B, Fβ3B and Fe23Bδ phases were

observed. Also, when the metal structure of the alloy was analyzed with a

transmission electron microscope, it was confirmed that the Nd2Fe B and FesB

phases coexisted in the alloy. Specifically, the Nd₂Feι₄B phase existed as crystal

grains with an average grain size of about 50 nm to about 150 nm, while the FesB

phase existed in the grain boundary zone of the Nd2FeuB phase. The grain boundary zone had a thickness of several nm to about 20 nm.

The magnetic properties of the respective samples were measured with a

vibrating sample magnetometer (VSM). The results are also shown in Table 2.

Next, the concentration of Ti in the Nd2Fei4B phase was measured by an

APFIM analysis. Specifically, the annealed rapidly solidified alloy was polished

mechanically and processed into a prismatic rod-like sample. Furthermore, the

end of this rod-like sample was sharpened into a needle shape by an electrolytic

polishing technique. Then, the metal structure of the needle-like sample was

analyzed with the APFIM analysis.

FIG. 7 shows the cumulative concentration profiles of Nd, B and Ti as

measured in the depth direction of sample No. 2. In the graph shown in FIG. 7,

the cumulative numbers of ions as counted by the APFIM analysis are plotted in

the depth direction. Specifically, a range in which the number of atoms counted

increases in the depth direction corresponds to a region where the atoms exist.

On the other hand, a range in which the number of atoms counted is constant in

the depth direction corresponds to a region where the atoms do not exist.

The graph shown in FIG. 7 has a number of nodes at each of which the

gradient of the curve representing the count of the Nd, B or Ti ions changes. The range in which the concentrations of Nd, B and Ti were measured was divided

into regions Nos. 1 to 12 at these nodes, thereby calculating the concentrations of

Nd, B and Ti in these regions Nos. 1 to 12. The results of the calculations are

shown in FIGS. 8A and 8B. Specifically, the ordinate of the graph shown in FIG.

8A represents the Ti concentration and the abscissa thereof represents the Nd

concentration. On the other hand, the ordinate of the graph shown in FIG. 8B

represents the Ti concentration and the abscissa thereof represents the B

concentration.

As can be seen from FIG. 8A, in the regions where the Nd concentrations

are from about 8 at% to about 14 at%, the Ti concentrations are about 2 at% or

less. Since the Nd2Feι₄B phase exists in the regions having the Nd

concentrations of about 8 at% to about 14 at%, it can be seen that the

concentration of Ti in the Nd2Feι₄B phase is about 2 at% or less.

On the other hand, in the regions where the B concentrations are from

about 25 at% to about 35 at%, the Ti concentrations are about 7 at% or more as

can be seen from FIG. 8B. Since the iron-based boride exists in the regions

having the B concentrations of about 25 at% to about 35 at% (i.e., grain boundary

zones), it can be seen that the concentrations of Ti in the grain boundary zones

(or grain boundary phases) are about 7 at% or more. The concentrations of the additive metal(s) M (i.e., Ti, Cr and/or V) as

measured by the method described above for the respective samples are shown

in the following Table 3:

Table 3

As can be seen from Table 3, in each of the specific examples of preferred

embodiments of the present invention in which Ti was added, the concentration of

Ti in the crystal grains of the Nd2Fei4B type compound was about 2 at% or less.

Also, the concentration of Ti in the grain boundary between the Nd₂Fei4B crystal grains was greater than about 8 at%, which is much higher than that of Ti in the

Nd2Fei4B crystal grains themselves. Furthermore, the difference in Ti

concentration between the Nd₂Feι₄B (R₂FeuB) phase and the grain boundary

phases was as much as about 6 at% or more. On the other hand, in each of the

comparative examples in which Cr or V was added, the concentration of Cr or V in

the Nd2FeuB phase was greater than about 2 at%.

If the concentration of the additive M (e.g., Ti) in the Nd2Fei4B type

compound exceeds about 2 at%, then the magnetization of the Nd2Feι₄B crystal

grains decreases considerably. To avoid such a decrease in magnetization, the

concentration of Ti in the Nd2Fei4B crystal grains is preferably about 2 at% or less,

more preferably about 1.8 at% or less. The Ti concentration in the Nd₂Fe B

crystal grains is most preferably about 1.65 at% or less to further increase the

magnetization.

The concentration of Cr or V in the grain boundary zones of the

comparative examples was not so much different from the concentration of Ti in

the grain boundary zones of the examples of preferred embodiments of the

present invention. However, since no Ti was added in the comparative examples,

the structure of the comparative examples should be totally different from that of

the examples of preferred embodiments of the present invention. That is to say, the desired structure, in which the iron-based boride with high magnetization is

either finely dispersed in, or present in the form of a film over, the thin grain

boundary phases of the Nd₂Fei4B crystal grains, should not have been formed in

any of the comparative examples. The reason is that since no Ti was added in

the comparative examples, the crystal grains of the Nd₂Feι₄B type compound

cannot have nucleated faster and earlier than the α-Fe phase.

Also, in the comparative examples in which Cr or V was added instead of Ti,

the iron-based boride with high magnetization was not produced so much in the

grain boundary zones but an Nd₂Fe ₄B phase including a lot of Cr or V was

produced there instead. As a result, the magnetization decreased and the

remanence B_r of the resultant magnet was less than about 0.8 T.

The iron-based rare earth alloy magnet according to the preferred

embodiments of the present invention described above has a unique structure in

which the iron-based boride existing in the grain boundary zones covers the

crystal grains having hard magnetic properties. Thus, the iron-based rare earth

alloy magnet exhibits improved anticorrosiveness as well. Generally speaking, an

iron-based rare earth alloy magnet, including an R2T14Q type phase with hard

magnetic properties, shows inferior oxidation resistance and anticorrosiveness

unless treated in some way or other. This is because the R2T₁₄Q type phase includes an easily oxidizable rare earth element R at a high concentration.

However, in the preferred embodiments of the present invention, the crystal

grains of the R2T14Q type compound are covered with the iron-based boride in the

grain boundary zones with a low R concentration. Accordingly, the oxidation or

corrosion stops at the grain boundary zones and excessive oxidation or corrosion

can be prevented effectively. In addition, Ti, existing at a relatively high

concentration in the grain boundary zones, would also contribute to increasing the

oxidation resistance and anticorrosiveness. This is because Ti is a metal element

having high chemical stability. Furthermore, this structure can also adjust the

exchange coupling appropriately, thus realizing an excellent magnet in which

coercivity and remanence are well balanced.

INDUSTRIAL APPLICABILITY

According to various preferred embodiments of the present invention, a

melt of a material alloy, including an additive Ti, is rapidly cooled and solidified,

thereby realizing a permanent magnet that exhibits excellent magnetic

properties, including high coercivity and high magnetization, while reducing the

minimum required amount of a rare earth element to be included in the magnet.

Also, according to various preferred embodiments of the present invention, even if a rapidly solidified alloy is prepared by a melt quenching process at a

decreased cooling rate, the addition of Ti can effectively prevent the precipitation

of the α-Fe phase during the melt quenching process. Therefore, a strip casting

method, or a melt quenching process resulting in a relatively low cooling rate and

suitably applicable to mass production, can be adopted, thus reducing the

manufacturing cost advantageously.

It should be understood that the foregoing description is only illustrative of

the invention. Various alternatives and modifications can be devised by those

skilled in the art without departing from the invention. Accordingly, the present

invention is intended to embrace all such alternatives, modifications and

variances which fall within the scope of the appended claims.

Claims

1. An iron-based rare earth alloy nanocomposite magnet having a

composition represented by the general formula: (Feι-mT_m)ιoo-x-y-zQχRyTiz, where

T is at least one element selected from the group consisting of Co and Ni; Q is at

least one element selected from the group consisting of B and C; and R is at least

one rare earth element including substantially no La or Ce, the mole fractions x,

y, z and m satisfying the inequalities of:

10 at%<x≤17 at%;

7 at%≤y<10 at%;

0.5 at%≤z≤6 at%; and

O≤m≤O.5, respectively;

wherein the magnet comprises: crystal grains of an R2T14Q type compound

having an average grain size of about 20 nm to about 200 nm; and a

ferromagnetic iron-based boride that exists in a grain boundary between the

crystal grains of the R2T14Q type compound, and

wherein the ferromagnetic iron-based boride is dispersed in the grain

boundary or present in the form of a film over the grain boundary to cover the

surface of the crystal grains of the R2T14Q type compound at least partially.

2. The magnet of claim 1 , wherein the mole fractions x, y and z satisfy the inequalities of:

10 at%<x≤15 at%;

7 at%≤y≤9.3 at%; and

1.5 at%≤z≤5 at%.

3. The magnet of claim 1 or 2, wherein the magnet comprises: crystalline

phases, including the R2T14Q type compound and the ferromagnetic iron-based

boride, at about 95 vol% or more in total; and amorphous phases at about 5 vol%

or less.

4. The magnet of claim 3, wherein the magnet comprises the R2T14Q type

compound at about 65 vol% to about 85 vol%.

5. The magnet of one of claims 1 to 4, wherein the crystal grains of the

R2T14Q type compound have a Ti concentration of about 2 at% or less, and

wherein a Ti concentration in the grain boundary between the crystal

grains of the R₂T₁₄Q type compound is higher than the Ti concentration inside the

crystal grains of the R2T14Q type compound.

6. The magnet of claim 5, wherein the Ti concentration in the grain boundary between the crystal grains of the R2T14Q type compound is about 7

at% or more.

7. The magnet of one of claims 1 to 6, wherein the ferromagnetic

iron-based boride has an average size of about 50 nm or less as measured

along the thickness of the grain boundary.

8. The magnet of one of claims 1 to 7, wherein the ferromagnetic

iron-based boride is present in the form of a film having an average thickness of

about 20 nm or less over the grain boundary between the crystal grains of the

R2T14Q type compound.

9. The magnet of one of claims 1 to 7, wherein the ferromagnetic

iron-based boride exists in the grain boundary between the crystal grains of the

R2T14Q type compound and has an average major axis length of about 1 nm to

about 50 nm.

10. The magnet of one of claims 1 to 9, wherein on an arbitrary cross

section of the magnet, the crystal grains of the R₂T Q type compound have an

average size greater than an average size of the ferromagnetic iron-based boride.

11. The magnet of one of claims 1 to 10, wherein the mole fractions x and

z satisfy the inequalities of

10 at%<x≤14 at% and

0.5 at%≤z≤4 at%.

12. The magnet of one of claims 1 to 11, wherein the iron-based boride

comprises FesB and/or Fe 3Bδ.

13. The magnet of one of claims 1 to 12, wherein the magnet is in the

shape of a thin strip having a thickness of about 10 M m to about 300 M m.

14. The magnet of one of claims 1 to 12, wherein the magnet has been

pulverized into powder particles.

15. The magnet of claim 14, wherein the powder particles have a mean

particle size of about 30 M m to about 250 M m.

16. The magnet of one of claims 1 to 15, wherein the magnet has hard

magnetic properties as represented by a remanence B_r of about 0.80 T or more, a maximum energy product (BH)max of about 100 kJ/m³ or more and a coercivity

Hcj of about 480 kA/m or more.

17. The magnet of claim 16, wherein the magnet has hard magnetic

properties as represented by a remanence B_r of about 0.85 T or more and a

maximum energy product (BH)_maχ of about 120 kJ/m³ or more.

18. A bonded magnet obtained by molding a magnet powder, comprising

a powder of the iron-based rare earth alloy magnet according to claim 14 or 15,

with a resin binder.