JP2021180059A - Rare earth compound, and magnetic recording medium using the same, and method for producing the same - Google Patents

Abstract

To provide a new compound having magnetic properties applicable to a magnetic recording medium by using a rare earth compound having a ThMn12 type structure, a magnetic storage medium using the same, and a method of producing the same.SOLUTION: A rare earth compound according to the embodiment of the present invention has overall composition represented by formula 1: (R1(1-x)R2x)aTbMc (in the above formula, R1 is at least one type of element selected from a group consisting of Sm, Pm, Er, Tm, and Yb, R2 is at least one type of element selected from a group consisting of Zr, Y, La, Ce, Pr, Nd, Eu, Gd, Tb, Dy, Ho, and Lu, T is at least one type of element selected from a group consisting of Fe, Co, and Ni, M is boron, x is a number from 0 to 0.5, a is a number of 6.0 to 13.7 atomic%, c is a number more than 0 atomic% and less than or equal to 12 atomic%, and b is a number represented by 100-a-c atomic%), and has an amorphous grain boundary phase existing between main phases by using a ThMn12 type of crystal phase as the main phase.SELECTED DRAWING: None

Description

The present invention relates to a rare earth compound, a magnetic recording medium using the rare earth compound, and a method for producing the same.

Due to cloud computing and the digitization of home appliances, the amount of information handled in the world is increasing explosively. The most important device for storing such data is a hard disk drive (HDD). HDDs have the advantages of low cost, large capacity, and non-volatile memory, and are therefore considered to be used as storage devices for large capacity data for the time being.

The part of the HDD that records information is called a magnetic recording medium, and the medium currently used has a fine structure in which CoCrPt alloy fine particles having a size of about 8 nm are uniformly dispersed in the non-magnetic matrix of _{SiO 2.} have. The current recording density is about 1.5 Tbit / in ² , but it is expected to achieve a density exceeding 4 Tbit / in ^2. Further fine particle formation of the CoCrPt alloy is required to increase the density of the medium. However, the magnetic anisotropy of the CoCrPt alloy is small, 4Tbit / in stochastic magnetization reversal due to thermal energy occurs ² 5nm less size required to implement the _{(K u V ≦ k B T} , K _{u V} anisotropic energy, k B _T of the ferromagnetic fine particles for thermal energy), the thermal disturbance problems. To overcome this, it is necessary to use a material having a large K _u to medium.

What is attracting attention as such a material is a FePt having an L1 ₀ ordered structure. Since FePt has a ^{magnetic anisotropy of 7 × 10 7} erg / cc, which is an order of magnitude higher than that of the CoCrPt alloy, it is possible to reduce the size to 4 nm. However, at the same time as reducing the size of the fine particles, the magnetization reversal magnetic field becomes a large value of 3T or more, which exceeds the magnetic field (about 1.5T) that can be generated by the magnetic head for writing, and writing cannot be performed by the current recording method. As a new magnetic recording method that solves this problem, heat-assisted magnetic recording (HAMR) is performed by locally raising the temperature of the medium to the curry point (the temperature at which ferromagnetism disappears) only during writing. ) Has been proposed. FePt is attracting attention as a HAMR medium, and there is an example in which Seagate has demonstrated a recording density ^{of 1.5 Tbit / in 2.}

Y. Hirayama, et al., Scr. Mater. 138 (2017) 62-65. D. Ogawa et al., Scr. Mater. 164 (2019) 140.

Although HAMR is expected to be put into practical use at an early stage, there is a problem in reliability because the operation of locally raising the temperature of the medium to near the Curie point (~ 700K) is repeated. Therefore, there is an urgent need to develop a medium material that has a low Curie point and exhibits high magnetic anisotropy.

Therefore, the present invention uses _{a rare earth compound having a ThMn 12-} type structure represented by a _{Sm (Fe 0.8} Co _0.2 ) ₁₂ compound, and a novel compound having magnetic properties applicable to a magnetic recording medium. It is an object of the present invention to provide a magnetic recording medium using this and a method for manufacturing them.

The Sm (Fe _0.8 Co _0.2 ) ₁₂ compound _{having a ThMn 12-} type structure is expected as a candidate material for permanent magnets exceeding the _{Nd 2} Fe _{14 B magnet.} For example, in Non-Patent Document 1, the Sm (Fe _0.8 Co _0.2 ) ₁₂ compound has a saturation magnetization of 1.78T, an anisotropic magnetic field of 12T, and a Curie temperature of 859K, all of which are Nd ₂ Fe. It has been reported that the value exceeds _{14 B.} Further, in Non-Patent Document 2, a coercive force of 0.87 T is obtained by diffusing a low melting point eutectic alloy such as Cu—Ga or Mg—Zn _{into 12 grain boundaries of Sm (Fe 0.8} Co _0.2 ). Has been reported to have been obtained.

As a result of diligent studies, the present inventors have found that the above-mentioned problems can be solved by containing a predetermined amount of boron (B) in the composition of a rare earth compound having _{a ThMn 12} type structure, and the present invention has been developed. It came to be completed.

That is, the gist of the present invention is the following [1] to [10].

[1] The overall composition is a ^{group consisting of formula 1: (R 1} _(1-x) R ² _x ) _a T _b M _c (in the above formula, R ¹ is Sm, Pm, Er, Tm, and Yb. at least one element more selective, at least 1 ^{R 2} is, Zr, Y, La, Ce , Pr, Nd, Eu, Gd, Tb, Dy, Ho, and selected from the group consisting of Lu An element of the species, T is at least one element selected from the group consisting of Fe, Co, and Ni, M is boron, x is a number from 0 to 0.5, a. Is a number of 6.0 to 13.7 atomic%, c is a number larger than 0 atomic% and 12 atomic% or less, and b is a number represented by 100-a-c atomic%. ), _{A rare earth compound having a ThMn 12} type crystal phase as a main phase and having an amorphous grain boundary phase existing between the main phases.
[2] The rare earth compound according to [1], wherein the absolute value of the difference between the boron content in the main phase and the boron content in the grain boundary phase is 1.0 atomic% or more.
[3] The rare earth compound according to [1] or [2], ^{wherein R 1 contains at least Sm.}
[4] The rare earth compound according to any one of [1] to [3], which contains substantially none of Ti, V, Mo, Nb, Cr, and W.
[5] The rare earth compound according to any one of [1] to [4], wherein T is Fe and Co.
[6] The rare earth compound according to any one of [1] to [5], wherein at least one selected from the group consisting of a crystal orientation and an easy magnetization axis is preferentially oriented along a predetermined direction.
Rare earth compound [7] The rare earth compound according to any one of [1] to [6], wherein a is 6.0 to 10.0 atomic%.
[8] A magnetic recording medium using the rare earth compound according to any one of [1] to [7] in the magnetic layer.
[9] The magnetic recording medium according to [8], wherein the crystal orientation of the magnetic layer is preferentially oriented in the [001] direction.
[10] A method for producing a magnetic layer containing the rare earth compound according to any one of [1] to [7] on a substrate by a sputtering method, wherein the substrate is heated to 250 to 400 ° C. ^{A method in which a target containing elements (R 1} , R ² , T, and M) constituting a rare earth compound is sputtered and then cooled at a temperature lower than the substrate heating temperature.

According to the present invention, by adding a predetermined amount of boron (B) to the composition of a rare earth compound having _{a ThMn 12 type structure, a ThMn 12} type crystal phase is used as a main phase, and an amorphous grain boundary phase is formed between the crystal phases. Can form a granular structure in which is present.
Therefore, according to the present invention, it is possible to provide a rare earth compound having an excellent coercive force, and the rare earth compound is used as a medium material having a low Curie point and exhibiting a high magnetic anisotropy, and the magnetism of a magnetic recording medium is used. Can be used for layers.
Further, according to the present invention, the magnetic layer containing this rare earth compound can be easily produced by a sputtering method.

It is a schematic diagram which shows the layer structure of the magnetic recording medium which concerns on embodiment of this invention. It is the magnetization curve in the in-plane (In-Plane (IP)) and the plane (Out-of-Plane (OOP) direction) of the thin films of Example 1 and Example 2 having a film thickness of 100 nm. BF-TEM images of in-plane and cross-sections of the thin films of Example 1 and Example 2; (a) Example 1 and (b) Example 2. HAADF-STEM images of the thin films of Examples 1 and 2; (a) and (b) Example 1, (c) and (d) Example 2. In-plane and cross-sectional STEM-EDS images of the thin films of Example 1 and Example 2; (a) Example 1 and (b) Example 2. (A) and (b) are the analysis results of the membrane of Example 1 by 3DAP. (A) and (b) are the analysis results of the membrane of Example 2 by 3DAP. It is a schematic diagram which shows the fine structure of the thin film of Example 1 and Example 2. FIG. Examples 3 (a) and (b) to (e) are Sm (Fe _0.8 Co _0.2 ) ₁₂ B film out-of-plane (Out-of-Plane) XRD patterns. Examples 3 (a) and (f) to (i) are Sm (Fe _0.8 Co _0.2 ) ₁₂ B film Out-of-Plane XRD patterns. 3 is an in-plane and perpendicular-to-plane magnetization curve of the films of Examples 3 (a) and (b) to (e). 3 is an in-plane and perpendicular-to-plane magnetization curve of the films of Examples 3 (a) and (f) to (i). 3 is a diagram showing the magnetic properties of the films of Examples 3 (c) to (i); (a) saturation magnetization (M _s ), (b) coercive force (H _c ), (c) residual magnetization ratio at zero magnetic field. ( _Mr / M _s ), (d) Vertical anisotropy ( _Ku ). It is an Out-of-Plane XRD pattern of the thin films of Examples 4, 7 to 10 and Example 12 in which the thickness of the film is changed from 20 nm to 200 nm. It is a change of the lattice constants a, c, and c / a of the thin films of Examples 4 to 12 in which the film thickness (t _{SFC) was changed.} 4 is a magnetization curve of the thin films of Examples 4, 7 to 10 and Example 12 in which the thickness of the film is changed. Magnetization of thin film with constant B content (M _s ), residual magnetization ratio in zero magnetic field ( _Mr / M _s ), coercive force (H _c ), vertical anisotropy ( _Ku ) film thickness Dependency on (t _SFC). It is the temperature dependence of the coercive force. It is an Out-of-Plane XRD pattern _{of Sm (Fe 0.8} Co _0.2 ) ₁₂ B film of Examples 13 and 15 to 21 in which the film thickness is 50 nm. It is the relationship between the lattice constants a, c, and c / a (vertical axis) of the thin films of Examples 13 to 18 and Example 21 in which the thickness of the film is constant, and the content of B (horizontal axis). 6 is an in-plane and perpendicular-to-plane magnetization curve of the thin films of Examples 13 and 15 to 21 in which the B content is changed while the film thickness is constant. B of thin film magnetization (M _s ) with constant film thickness, residual magnetization ratio in zero magnetic field ( _Mr / M _s ), coercive force (H _c ), and vertical anisotropy ( _Ku ) Dependence on content.

Hereinafter, embodiments of the present invention will be described in detail.
The description of the constituent elements described below may be based on a representative embodiment of the present invention, but the present invention is not limited to such embodiments.
In the present specification, the numerical range represented by using "~" means a range including the numerical values before and after "~" as the lower limit value and the upper limit value unless otherwise specified.

[Rare earth compounds]
The rare earth compound according to the embodiment of the present invention (hereinafter, also referred to as “rare earth compound of the present embodiment”) has an overall composition of the formula 1: (R ¹ _(1-x) R ² _x ) _a T _b M _c. It is represented by a ThMn ₁₂ type crystal phase as a main phase, and has an amorphous grain boundary phase existing between the main phases.
Here, in the above formula, R ¹ is at least one element selected from the group consisting of Sm, Pm, Er, Tm, and Yb, and R ² is Zr, La, Ce, Pr, Nd. , Eu, Gd, Tb, Dy, Ho, and at least one element selected from the group consisting of Lu, and T is at least one element selected from the group consisting of Fe, Co, and Ni. M is boron, x is a number from 0 to 0.5 (0 ≦ x ≦ 0.5), and a is a number from 6.0 to 13.7 atomic% (6.0). ≤a≤13.7), c is a number greater than 0 atomic% and 12 atomic% or less (0 <c≤12), and b is a number represented by 100-ac atomic%.

Hereinafter, the composition of the rare earth compound of the present embodiment will be described in detail.

[Overall composition]
The overall composition of the rare earth compound of this embodiment is represented by (R ¹ _(1-x) R ² _x ) _a T _b M _c .
In the present specification, the "overall composition" is determined by an ICP emission spectrophotometer (ICP-OES), and the measuring method thereof is as described in Examples described later.

^{・ R 1} in the overall composition
In formula 1, R ¹ is at least one element selected from the group consisting of Sm (samarium), Pm (promethium), Er (erbium), Tm (thulium), and Yb (ytterbium) (hereinafter, "" It is also called "specific rare earth element"). As ^{the specific rare earth element of R 1} , one kind may be used alone, or two or more kinds may be used in combination.

The specific rare earth element is a rare earth element having a Stevens factor. The Stevens factor is a physical quantity related to the electrification density (shape) of 4f electrons in the inner shell of a rare earth element. If this is negative, it will be a shape that contracts with respect to the axis of symmetry, and if it is positive, it will be a shape that extends from spherical symmetry. Since the 4f electron cloud receives a crystal field from surrounding ions and its stable direction is determined, the shape of the electron cloud determines the direction of magnetic anisotropy.
Since the rare earth compound of the present embodiment contains an element in which the Stevens factor is positive, it has excellent magnetic properties.

Among them, in that the rare earth compound having an effect of better present invention is obtained, as the R ^1, at least 1 Stevens factor positive matrix Sm, Yb, Tm, and selected from the group consisting of Pm Species are preferred, at least one selected from the group consisting of Sm, Yb, and Tm is more preferred, and even more preferably at least one selected from the group consisting of Sm, Yb.

Among them, Sm is presumed to have a positive Stevens factor and a ferromagnetic bond with an atom represented by T described later in the ground state, and when R ¹ contains Sm, the obtained rare earth compound is more excellent. It also has the effect of the present invention.
Therefore, rare earth compound of the present embodiment, as R ^1, it is preferable to contain at least Sm.

^{・ R 2} in the overall composition
In formula 1, R ² is Zr (zyrium), Y (yttrium), La (lantern), Ce (cerium), Pr (placeodim), Nd (neopium), Eu (europium), Gd (gadolinium), Tb ( It is at least one selected from the group consisting of terbium), Dy (dysprosium), Ho (hormium), and Lu (lutettrium) (hereinafter, also referred to as “specific element”). Specific element of R ² may be used alone or in combination of two or more thereof.

R ² is a component that contributes to the stabilization of ThMn ₁₂ type crystal phase having the rare earth compound of the present embodiment.
Among, Y, and, Ce (among others, Ce (IV) is preferred.), Although it itself is an element having no magnetism, by rare earth compound contains such as R ^2, obtained rare earth The compound has excellent stability.

Gd and Zr have a function of further improving the stability of the obtained rare earth compound, and the effect is more remarkable ^{especially when R 1 contains Sm.} That is, when the rare earth compound of the present embodiment contains at least Sm as ^{R 1} , the rare earth compound of the present embodiment is at least one specific element selected from the group consisting of Gd and Zr as ^{R 2.} Is preferably contained.

La, Ce, Pr, Nd, Eu, Gd, Tb, Dy, Ho, and Lu have a negative or zero Stevens factor, and are of the ThMn ₁₂ type crystal phase of the rare earth compound of the present embodiment. Contributes to stabilization. On the other hand, since the R ² element ^{also has an action of weakening the magnetic anisotropy derived from the R 1} element, it is preferable to adjust the content thereof from both the stability of the crystal phase and the magnetic characteristics.

^{Specifically, the ratio of the content of R 2 based on} the atomic% (at%) to the total content ^{of R 1} and R ² in the rare earth compound of the present embodiment (number represented by x in the formula). Is 0.5 or less, preferably 0.4 or less, and more preferably 0.3 or less.
Although not particularly limited as lower limit, rare earth compound of the present embodiment may not contain R ^2. That is, in the rare earth compound of the present embodiment, 0 ≦ x ≦ 0.5, 0 ≦ x ≦ 0.4 is preferable, and 0 ≦ x ≦ 0.3 is more preferable.

^{The total content a of R 1} and R ² on an atomic% basis in the overall composition is 6.0 ≦ a ≦ 13.7. In terms of obtaining a rare earth compound having a better effect of the present invention, a is preferably 6.0 ≦ a ≦ 10.

・ T in the overall composition
In formula 1, T is at least one element selected from the group consisting of Fe (iron), Co (cobalt), and Ni (nickel). These are classified as iron group elements and have common properties in that they exhibit ferromagnetism at normal temperature and pressure. Therefore, Fe, Co, and Ni as T can be substituted with each other, and the iron group element may be used alone or in combination of two or more as T.
The content of T in the overall composition is not particularly limited as long as it satisfies 100-ac in relation to a described above and c described later, but is generally preferably 50 to 95 atomic%, preferably 60 to 85. Atomic% is more preferred.

T preferably contains at least one element selected from the group consisting of Fe and Co, in that a rare earth compound having a more excellent effect of the present invention can be obtained, and Fe and Co are used. It is preferable to use them together.
When Co is contained as T, the magnetization of the rare earth compound is further improved, and the Curie temperature is further increased. In other words, in the rare earth compound of the present embodiment, the Curie point can be controlled without impairing the magnetic properties by adjusting the content of Co as T.

T is preferably Fe and Co. That is, the overall composition of the rare earth compound of the present embodiment is preferably represented by ^{the formula 2: (R 1} _(1-x) R ² _x ) _a (Fe _p Co _1-p ) _b M _c. At this time, p is a number of 0.5 to 1.0.
Among them, in that a rare earth compound having a more excellent effect of the present invention can be obtained, the overall composition of the rare earth compound of the present embodiment is as follows: Formula 3: (Sm _1-x R ² _x ) _a (Fe _p Co _1). it is more preferably represented by _{-p) b} M _c. At this time, p is a number of 0.5 to 1.0, preferably 0.9 to 1.0.

・ M in the overall composition
M is boron (boron element, B). The content of boron in the overall composition is preferably larger than 0 atomic% and preferably 12 atomic% or less.

Conventionally, by adding a light element such as boron to the 1-12 phase (ThMn ₁₂ type crystal phase), the coercive force of the rare earth compound obtained due to the change in magnetic anisotropy in the plane is lowered. It has been thought that it would be done.
However, the present inventors have examined the possibility of various elements from a multifaceted viewpoint without being bound by the above-mentioned common general technical knowledge. As a result of these efforts, it was found that a rare earth compound having an overall composition represented by the formula 1 containing a predetermined amount of boron has an excellent coercive force, and completed the present invention.

Further, the rare earth compound of the present embodiment is surprisingly composed of crystals such as α-Fe (α- (Fe, Co)) phase in addition to the 1-12 phase in the structure, as shown in Examples described later. Despite having a phase, it has an excellent coercive force. In general, the mixture of the above other crystal phases in the structure of a rare earth compound having a phase 1-12 as a main phase may cause a decrease in coercive force or a decrease in the angularity of a demagnetization curve. However, due to the diligent studies of the present inventors, the rare earth compounds of the present embodiment may be unpredictable by those skilled in the art even when the above-mentioned other crystal phases are contained. It was revealed that it has excellent properties.

The reason why the excellent coercive force is exhibited even when the crystal phase has an α-Fe (α- (Fe, Co)) phase in addition to the 1-12 phase is not always clear, but the structure shown in Examples described later. From the results of the analysis, the present inventors infer as follows. That is, by adding boron, an amorphous phase (B-concentrated phase) having a high boron concentration is formed at the grain boundary of the crystal phase having the 1-12 phase as the main phase, and α-Fe (α- (Fe, Co)). Even if a soft magnetic phase such as the phase) is formed, the movement of the magnetic wall is restricted by the non-magnetic boron compound contained in the amorphous grain boundary phase (B-concentrated phase), or the magnetic wall between the amorphous phase and the main phase. It is considered that the coercive force increased because the energy difference was steep.

Further, from the viewpoint of more reliably obtaining the effect of improving the coercive force by the amorphous grain boundary phase (B-concentrated phase), the content of boron in the main phase (1-12 phase) and the amorphous grain boundary phase (the amorphous grain boundary phase). The absolute value of the difference in the content of boron in the B-enriched phase) is preferably 1.0 atomic% or more. Further, the absolute value of the difference in the content of the boron may be a larger value depending on the content of boron in the overall composition, and may be, for example, 1.5 atomic% or more. It may be 0 atom% or more, 3.0 atom% or more, 4.0 atom% or more, 5.0 atom% or more, 7.5 atom% or more. It may be% or more, and may be 10 atomic% or more.

More specifically, the content of boron in the main phase (1-12 phase) is more than 0 atomic%. Therefore, the content of boron in the amorphous grain boundary phase (B-concentrated phase) is preferably more than 1.0 atomic%. Further, the content of boron in the amorphous grain boundary phase (B-concentrated phase) may be 1.5 atomic% or more, 2.0 atomic% or more, or 3.0 atoms. % Or more, 4.0 atomic% or more, 5.0 atomic% or more, 7.5 atomic% or more, 10 atomic% or more. You may.

The upper limit of the boron content in the main phase (1-12 phase) is a value obtained by dividing 1.0 from half the value based on the boron content (atomic%) in the overall composition. The value is preferably less than. For example, in the embodiment in which the boron content in the overall composition of the rare earth compound of the present embodiment is 12 atomic%, the boron content in the main phase (1-12 phase) is preferably less than 5.0 atomic%. 4.0 atomic% or less is more preferable, 3.0 atomic% or less is further preferable, 2.0 atomic% or less is particularly preferable, 1.0 atomic% or less is most preferable, and 0.5 atomic% or less is more most preferable. preferable.

The boron content in the amorphous grain boundary phase (B-concentrated phase) is preferably a value equal to or more than half of the boron content (atomic%) in the overall composition. For example, in the embodiment in which the content of boron in the overall composition of the rare earth compound of the present embodiment is 12 atomic%, the content of boron in the amorphous grain boundary phase (B-concentrated phase) is 6.0 atoms. % Or more is preferable, 6.5 atomic% or more is more preferable, 7.0 atomic% or more is further preferable, 8.0 atomic% or more is particularly preferable, 9.0 atomic% or more is most preferable, and 10 atomic% or more is more preferable. More preferably, 10.5 atomic% or more is even more preferable. It should be understood that the upper limit of the boron content in the amorphous grain boundary phase (B-concentrated phase) is not particularly limited and may vary depending on the boron content in the overall composition.

Further, the rare earth compound of the present embodiment is Ti (tungsten), V (vanadium), Mo (molybdenum), Nb (niobium), Cr (chromium) in that a rare earth compound having a better effect of the present invention can be obtained. ) And W (tungsten) (hereinafter, also referred to as “excluded element”) are preferably not substantially contained. Rare earth compounds that are substantially free of the above excluded elements have a better maximum magnetic energy product (BH) _max and coercive force.

In the present specification, "substantially free" means that the content of the excluded element is 0.1 atomic% or less of all atoms when the elements contained in the main phase are analyzed by ICP-OES analysis. It is more preferably 0.01 atomic% or less, and further preferably 0.001 atomic% or less.
When the main phase contains two or more kinds of excluded elements, it is preferable that the total of the two or more kinds of excluded elements is within the above numerical range.

The rare earth compound of the present embodiment may have another crystal phase as long as it has a _{ThMn 12 type crystal phase.} The other crystal phase is not particularly limited, but for example, when the overall composition is _{represented by Sm (Fe 0.8} Co _0.2 ) ₁₂ B, α-Fe, α- (Fe, Co), Sm (FeCo). ) ₇ and Sm ₂ (FeCo) ₁₇ phase and the like.
The rare earth compound of the present embodiment has a ChMn ₁₂ type crystal phase as a main phase, and α-Fe, α- (Fe, Co), Sm (FeCo) _{7 in that a better effect of the present invention can be easily obtained.} , And Sm ₂ (FeCo) ₁₇ phases and the like are preferable as other crystal phases.
In the present specification, the main phase means the phase having the highest peak intensity detected in the X-ray diffraction measurement of the rare earth compound.

In the rare earth compound of the present embodiment, the crystal orientation and the orientation state of the easy magnetization axis are not particularly limited, but from the crystal orientation and the easy magnetization axis in that the maximum energy product (BH) _max tends to be larger. It is preferable that at least one selected from the group is preferentially oriented along a predetermined direction. A rare earth compound in which at least one selected from the group consisting of a crystal orientation and an easy magnetization axis is preferentially oriented along a predetermined direction is also referred to as a rare earth compound having magnetic anisotropy in the present specification.
That is, the rare earth compound of the present embodiment is preferably a rare earth compound having magnetic anisotropy.

For example, in _{Sm (Fe 0.8} Co _0.2 ) ₁₂ , which is a magnetic compound having _{a ThMn 12} type crystal phase, the [001] direction is the easy magnetization axis, and the easy magnetization axis is preferentially oriented in a predetermined direction. If so, a better maximum energy product (BH) _max is obtained.
The rare earth compound of the present embodiment has a ThMn ₁₂ type crystal phase, and when the crystal orientation and / or the easy magnetization axis is preferentially oriented in a predetermined direction, the effect of the present invention is more excellent. It is easy to obtain a rare earth compound having.

In addition, in this specification, "priority orientation" means in-plane measurement (In-Plane XRD) of XRD (X-ray diffraction), or ordinary thin film X-ray diffraction (out-of-plane). In XRD), the peak intensity of the diffraction line from (00L) _{of the ThMn 12} type crystal phase with respect to the peak intensity of the phase (second phase) other than the _{detected ThMn 12 type crystal phase is 1 or more.} It means, and it is preferable that it is 2 or more.
The rare earth compound of the present embodiment may have a _{phase other than the ThMn 12} type crystal phase (such as the α-Fe phase already described), and in this case, the peak derived from the α-Fe phase is the ThMn ₁₂ type. It is not included in the peak derived from the predetermined crystal orientation of the crystal phase of.

[Use of rare earth compounds]
Since the rare earth compound of the present embodiment has an excellent coercive force, it can be preferably used for the magnetic layer of a magnetic recording medium as a medium material having a low Curie point and exhibiting high magnetic anisotropy.

[Magnetic recording medium]
FIG. 1 is a schematic diagram showing a layered structure of a magnetic recording medium according to an embodiment of the present invention (hereinafter, also referred to as “magnetic recording medium of the present embodiment”).
In FIG. 1, the magnetic recording medium 10 (perpendicular magnetic recording medium) of the present embodiment includes a heat absorption layer 30, a base layer 40, and a magnetic layer 50 in this order on a substrate 20.

The substrate 20 is not limited, but a MgO single crystal substrate or a glass substrate is preferably used. Examples of the glass for the substrate include aluminosilicate glass, aluminoborosilicate glass, soda-time glass and the like, and among them, aluminosilicate glass is preferable. Further, amorphous glass and crystallized glass can be used. It is preferable to use chemically strengthened glass because of its high rigidity. In the present embodiment, the surface roughness of the main surface of the substrate _{is preferably 10 nm or less in R max} and 0.3 nm or less in Ra.

In the heat-assisted magnetic recording device, the writing area temporarily becomes a high temperature of 600 K or more due to the laser when writing to the magnetic recording medium. Therefore, in the magnetic recording medium of the present embodiment, the heat absorption layer 30 is provided on the substrate 20. It is preferable to provide it. As the material of the heat absorption layer 30, a material having high thermal conductivity can be used, and specific examples thereof include metals such as Ta, Cu (including a laminated film of Ta and Cu), and NiTa alloy. It is appropriate that the film thickness of the heat absorption layer 30 is in the range of about 50 to 100 nm. The heat absorption layer 30 can be formed by using a sputtering method.

An underlayer 40 (orientation control layer) is provided on the substrate 20, preferably on the substrate 20 on which the heat absorption layer 30 is provided. _{The underlayer 40 has a vertical orientation of the easy-to-magnetize axis of the ThMn 12} type structure in the magnetic layer 50 (vertical magnetic recording layer) formed on the base layer 40 (the crystal orientation is oriented in the direction perpendicular to the substrate surface) and crystals. It is used to suitably control grain boundary segregation for forming a granular structure in which the main phase (1-12 phase) is surrounded by an amorphous grain boundary phase, and the like.

Further, in the present embodiment, the underlayer 40 is preferably in a form capable of reducing lattice defects due to lattice mismatch of the upper magnetic layer 50 and improving crystallinity. That is, the material of the base layer 40 is more preferably the same as the lattice constant of the magnetic layer 50, and the material of the magnetic layer 50 has a lattice constant mismatch of 10% or less with the rare earth compound having _{a ThMn 12 type structure.} Especially suitable. Since the lattice mismatch of the base layer 40 with the rare earth compound in the magnetic layer 50 is within the above range, the crystal orientation of the magnetic layer 50 made of a material containing the rare earth compound as a main component is disturbed by the base layer 40. The effect of suppressing and improving the fine structure is well exhibited.

The material of the base layer 40 may be a single crystal or a polycrystal, but is preferably a single crystal in that a magnetic recording medium having a better effect of the present invention can be obtained. By using the base layer 40 which is a single crystal, the magnetic layer 50 formed on the base layer 40 is more likely to grow epitaxially, and as a result, a magnetic recording medium having a better effect of the present invention can be obtained.

Specifically, in the present embodiment, a film having a two-layer structure of MgO / X or Mg—Ti—O / X is preferably used as the base layer 40. Here, X is at least one metal or compound selected from the group consisting of Pt, TiN, Ag, Au, Cr, Mo, V, Ti, Ta, Nb, and W.

Further, in the present embodiment, CrRu, Pt, Cr, V may be used as the base layer 40, and in this case, the base layer 40 may be a single layer or may be composed of a plurality of layers. In the case of multiple layers, not only the same material combination but also different materials can be combined. In that case, it may be any of a solid solution, a eutectic (eutectic mixture), an intermetallic compound, and a mixture thereof.

The film thickness of the base layer 40 is not particularly limited, but it is desirable that the film thickness is the minimum necessary for controlling the structure of the magnetic layer 50. For example, the film thickness may be in the range of about 5 to 30 nm as a whole. Appropriate. The base layer 40 can be formed by using a sputtering method.

The magnetic layer 50 is made of a material containing the rare earth compound according to the above-described embodiment of the present invention as a main component. The magnetic layer 50 may contain the rare earth compound, and its preferred form is as described above. Further, in the present embodiment, it is preferable that the crystal orientation of the magnetic layer 50 is preferentially oriented in the [001] direction.

In this embodiment, the magnetic layer 50 can be formed by using a sputtering method. Hereinafter, the form of forming the magnetic layer 50 by the sputtering method will be described in detail.

^{The pressure in the chamber of the film forming apparatus during sputtering is not particularly limited, but is preferably 10-6} Pa or less, preferably ^10-8 Pa or less, from the viewpoint of further reducing the mixing of unintended components in the obtained magnetic layer 50. Pa or less is more preferable.

Further, when the heat absorption layer 30, the base layer 40, and the magnetic layer 50 are laminated on the substrate 20 by the sputtering method, it is preferable to clean the surface of the substrate 20 in advance. The method for cleaning the surface of the substrate 20 is not particularly limited, and examples thereof include a method of sputtering the substrate 20 itself. According to the above, the oxide film formed on the substrate 20 and organic substances can be removed. Further, the surface of the substrate 20 may be cleaned by heating the substrate 20 to a predetermined temperature (for example, about 600 to 800 ° C.), holding the substrate 20 for a predetermined time (for example, about 10 to 30 minutes), and performing a heat treatment. ..

The sputtering method is not particularly limited, but a magnetron sputtering method that enables sputtering in a lower pressure Ar atmosphere is preferable. Here, by adjusting the thickness of the target material, the reduction of the leakage flux of the magnetron sputtering can be further suppressed, and the sputtering can be made easier. The power source for sputtering can be either DC or RF, and can be appropriately selected depending on the target material.

The substrate heating temperature (heating temperature of the substrate 20), the film formation rate, and the film formation time when forming the magnetic layer 50 are not particularly limited, and may be appropriately adjusted according to the required thickness of the magnetic layer 50. .. The film formation rate can be adjusted by the power of sputtering and / or the time.
The substrate heating temperature is preferably 250 to 400 ° C, more preferably 300 to 350 ° C, from the viewpoint of more efficiently obtaining the magnetic recording medium of the present embodiment having the above-mentioned structure.
Further, after the formation of the magnetic layer 50 (lamination of the magnetic layer 50), it is preferable to cool the magnetic layer 50 at a temperature lower than the substrate heating temperature. Thereby, the effect of the magnetic recording medium of the present embodiment can be obtained more reliably. The cooling means and the cooling time when cooling at a temperature lower than the substrate heating temperature are not particularly limited. For example, after the magnetic layer 50 is formed, the substrate 20 may be naturally cooled without being heated and cooled. In the case of the above natural cooling, the temperature of the substrate 20 varies depending on the value of the substrate heating temperature, the atmosphere in the chamber of the film forming apparatus, and the like. It drops to about 50-70 ° C.
Further, after the magnetic layer 50 is formed, the substrate 20 may be cooled by setting the temperature of the substrate heating means to a temperature lower than the substrate heating temperature and holding the magnetic layer 50 for a predetermined time. In this case, the set temperature of the substrate heating means may be any temperature as long as it is lower than the substrate heating temperature, and the holding time may be any time.
Further, the cooling by the substrate heating means described above and the natural cooling may be combined.
Further, after the magnetic layer 50 is formed, the substrate 20 may be cooled by using a cooling medium such as nitrogen gas, whereby the magnetic recording medium of the present embodiment can be obtained in a shorter production time. In the case of cooling using the above cooling medium, the temperature of the substrate 20 drops to about 20 ° C. in about several minutes to 10 minutes, for example, when nitrogen gas is used, although it depends on the type of cooling medium and the value of the substrate heating temperature. ..
Regarding the above-mentioned cooling conditions, the average cooling rate (° C./° C.) obtained by dividing the difference between the substrate heating temperature at the time of film formation of the magnetic layer 50 and the substrate temperature at the end of cooling (the temperature of the substrate 20) by the cooling time. As a rough guideline for minutes), it may be about 3 to 85 ° C./min.

Further, it is preferable to provide a protective layer (not shown) on the magnetic layer 50 (perpendicular magnetic recording layer). By providing the protective layer, the surface of the magnetic recording medium can be protected from the magnetic head floating and flying on the magnetic recording medium. As the material of the protective layer, for example, a carbon-based protective layer is suitable. The film thickness of the protective layer is preferably about 3 to 7 nm. The protective layer can be formed by, for example, a plasma CVD method or a sputtering method.

Further, it is preferable to further provide a lubricating layer (not shown) on the protective layer. By providing the lubricating layer, wear between the magnetic head and the magnetic recording medium can be suppressed, and the durability of the magnetic recording medium can be improved. As the material of the lubricating layer, for example, a perfluoropolyether (PFPE) -based compound is preferably used. The lubricating layer can be formed, for example, by a dip coating method. The thickness of the lubricating layer is preferably about 0 to 10 nm.

Hereinafter, the present invention will be described in more detail based on Examples.
The materials, amounts used, ratios, treatment contents, treatment procedures, etc. shown in the following examples can be appropriately changed as long as they do not deviate from the gist of the present invention. Therefore, the scope of the present invention should not be construed as limiting by the examples shown below.

[Preparation of an epitaxial thin film containing a rare earth compound whose overall composition is Sm (Fe _0.8 Co _0.2 ) _{12 B]}
The DC magnetron sputtering method was used to prepare the thin film.
Under an Ar atmosphere of 0.167 Pa (1.30 mTorr), the pressure in the chamber of the DC magnetron sputtering device compatible with ultra high vacuum (UHV) was set to ^10-8 Pa or less, and the MgO (100) single crystal substrate was 20 at 700 ° C. The surface was cleaned by heat treatment for a minute. Then, at a substrate temperature of 325 ° C., Sm (Fe _0.8 Co _0.2 ) ₁₂ and V (thickness = 20 nm) having a small lattice mismatch are arranged on this MgO (100) single crystal substrate as a base layer. (Hereinafter, the MgO (100) single crystal substrate on which the above-mentioned base layer is arranged is also simply referred to as a "substrate").
_{Subsequently, the Sm, Fe, Fe 50} Co ₅₀ , and Fe ₈₀ B ₂₀ targets are simultaneously sputtered at the same substrate temperature of 325 ° C. to form a Sm (Fe _0.8 Co _0.2 ) ₁₂ B layer. It was a film. Further, V (thickness = 10 nm) was deposited as a cap layer to prevent oxidation.

The relationship between the setting conditions of the DC magnetron sputtering apparatus and the film formation rate was measured in advance using a step meter, and using this result, the amount of boron introduced in each sample film was estimated from the film formation rate.
That is, even if the films have the same thickness, the amount of boron introduced can be controlled by controlling the film formation rate, and if the film formation rate is constant, the samples having a constant boron content and different film thicknesses. Can also be formed. _{According to the preliminary experiments by the present inventors, Sm (Fe 0.8} Co _0.2 ) having a thickness of 5 nm or more on the substrate using the above materials, film forming conditions, and film forming apparatus. ₁₂ It has been confirmed that the B layer can be formed. Further, if necessary, _{the thickness of the Sm (Fe 0.8} Co _0.2 ) ₁₂ B layer can be set to less than 5 nm.

ICP-OES analysis was performed by the following procedure.
First, a sample is taken in a quartz beaker, 5 ml of a 1: 1 (volume) solution of nitric acid and water, 10 ml of a 1: 1 (volume) solution of hydrochloric acid and water, and 1: 1 of sulfuric acid and water. (Volume) 3 ml of the solution was added, heated at 120 ° C. for 30 minutes to dissolve, allowed to cool, and then set to 100 ml. The content of each element in this solution was measured by an ICP-OES apparatus "720-ES ICP-OES" manufactured by Agilent.

(Examples 1 to 2)
Boron introduction of (B) (target) amount was varied from 0 vol% to 1.5 vol%, to prepare a _{Sm (Fe 0.8 Co 0.2) 12} B film having a thickness of 100nm on the substrate .. In the following, a sample having a B content of 0% by volume (Sm (Fe _0.8 Co _0.2 ) ₁₂ membranes; Example 1) and a sample having a B content of 0.5% by volume (Sm). The characteristics of (Fe _0.8 Co _0.2 ) ₁₂ B _0.5 film; Example 2) will be described.

2 (a) and 2 (b) show in-plane (In-Plane (IP)) and straight (Out-of-Plane) (In-Plane (IP)) and in-plane (Out-Plane (IP)) of the thin films of Example 1 and Example 2 having a film thickness of 100 nm, respectively. The magnetization curve in the OOP)) direction is shown. A superconducting quantum interference magnetic flux meter (SQUID) was used to measure the magnetic characteristics, and a magnetic field was applied in the range of ± 7T (70 kOe) at room temperature. The same applies to FIGS. 11, 12, 16 and 21 which will be described later.

According to FIG. 2A, the film of Example 1 in which B is 0% by volume shows strong vertical anisotropy. The coercive force was 0.1T and the residual magnetization at zero magnetic field was 0.2T.
According to FIG. 2 (b), the film of Example 2 in which B is 0.5% by volume shows a strong vertical anisotropy comparable to that of Example 1 and 1.2T significantly higher than that of Example 1. The coercive force of was shown. The residual magnetization in the zero magnetic field was 1.50 T. In addition, the temperature coefficient (β) of the coercive force of the film of Example 2 in the temperature range of 300 to 500 K (27 to 227 ° C) was calculated to be −0.22% / ° C. Since the temperature coefficient (β) of the coercive force of the previously reported anisotropic Nd-Fe-B magnet is about -0.4 to -0.6% / ° C., it was obtained for the film of Example 2. The absolute value of the temperature coefficient is much smaller than that of the conventional Nd-Fe-B magnet, which means that the film containing the rare earth compound according to the embodiment of the present invention is the conventional Nd-Fe-B magnet. It is shown that the thermal stability of the coercive force is superior to that of the above.

Regarding the magnetization curve of the film of Example 2 (FIG. 2 (b)), the demagnetization rate is set to 0.83 and the demagnetization rate is corrected based on the result of the microstructure analysis of the thin film described later. The value of residual magnetization at zero magnetic field (1.50T) obtained from the film of Example 2 is the value (1.61T) of the previously reported anisotropic Nd ₂ Fe ₁₄ B / FeCo nanocomposite thin film. Although it is slightly smaller than the above, it is the largest value among _{the anisotropic Sm (Fe 0.8} Co _0.2 ) ₁₂ films as far as the present inventors know. Further, from the XRD diffraction pattern (data not shown), in the film of Example 1 in which B is 0% by volume, the lattice constants a, c, and c / a are a = 0.8524 nm and c =, respectively. In the film of Example 2 in which 0.4811 nm, c / a = 0.564, and B is 0.5% by volume, a = 0.8656 nm, c = 0.4755 nm, and c / a = 0.549. rice field.

3 (a) and 3 (b) show in-plane and cross-sectional transmission electron microscope images (BF-TEM images) of the thin films of Examples 1 and 2, respectively.

In FIG. 3A (film of Example 1), it can be seen from the upper in-plane image that the thin film is continuous and from the lower cross-sectional image that it has a columnar structure. No grain boundary phase is observed between the particles having a columnar structure.
On the other hand, in FIG. 3 (b) (film of Example 2), columnar SmFe _12- based crystal grains were confirmed. The average particle size of the columnar grains was about 40 to 50 nm, and the ratio of height to width (average particle size) was about 2.5: 1 (average aspect ratio of about 2.5). Further, it can be seen that the structure is such that the columnar grains are used as the main phase and grain boundary phases having a certain width (about 1 nm or more) exist between the main phases.
From these BF-TEM images, in the film containing the rare earth compound according to the embodiment of the present invention, a predetermined amount of boron is added to obtain a ThMn ₁₂ type crystal phase (1-12 phase) as the main phase. It was suggested that a grain boundary phase having a certain thickness was formed between the columnar crystal phases as the columnar growth occurred, and that such a fine structure exerted a high coercive force.

FIGS. 4 (a) to 4 (d) show images of a high-angle annular dark-field scanning transmission electron microscope (HAADF-STEM). 4 (a) and 4 (b) are _{HAADF-STEM images of the thin film (Sm (Fe 0.8} Co _0.2 ) ₁₂ film) of Example 1, and FIGS. 4 (c) and 4 (d) are examples. It is a HAADF-STEM image of the thin film of 2 (Sm (Fe _0.8 Co _0.2 ) ₁₂ B _{0.5 film).}

According to FIG. 4A, in the film of Example 1, no _{clear phase was observed between the two Sm (Fe, Co) 12} crystal grains, and the two crystal grains were in direct contact with each other. I understand. FIG. 4B is a high-magnification image obtained along the crystal zone axis of [100] of the Sm (Fe, Co) _{12 crystal grains in the film of Example 1 (direction on the paper surface of FIG. 4A).} .. Since contrast occurs in the HAADF-STEM image due to the number of atomic atoms of the elements contained in the sample, what is seen as a brighter spot in the image corresponds to the atomic column of the heavier element, ie Sm (Fe). , Co) It is considered to be Sm in _{12 compounds.} Further, by projecting the nanobeam electron diffraction pattern (upper right) in the region indicated by the symbol i in FIG. 4 (b) and the result of the atomic resolution STEM-EDS mapping, Sm (Fe, Co) in the film of Example 1 is projected. ) ₁₂ crystal grains were shown to have a _{ThMn 12 type structure.} However, in the film of Example 1, as shown by the symbol ii in FIG. 4 (b), there are regions where the results of the nanobeam electron diffraction pattern (upper left) and the atomic resolution STEM-EDS mapping are different.

According to FIG. 4 (c), in the film of Example 2, an amorphous grain boundary phase (GB) having a thickness of about 3 nm was confirmed between _{two Sm (Fe, Co) 12 crystal grains.} FIG. 4D is a high-magnification image obtained along the crystal zone axis of [100] of the Sm (Fe, Co) _{12 crystal grains in the film of Example 2 (direction on the paper surface of FIG. 4C).} .. From the nanobeam electron diffraction patterns (upper right and upper left) inserted in FIG. 4 (d), the two Sm (Fe, Co) ₁₂ crystal grains separated by the grain boundary phase both have a ThMn ₁₂ type structure. The misalignment is as small as about 2 °, and _{the [001] direction (c-axis) of the Sm (Fe 0.8} Co _0.2 ) ₁₂ crystal grains that are close to each other is upward (upward on the paper surface in FIG. 4 (d)). It was shown to be oriented to.

5 (a) and 5 (b) show the results (STEM-EDS image) of the in-plane and cross-sectional scanning transmission electron microscope-energy dispersive X-ray analysis of the thin films of Examples 1 and 2, respectively. rice field.

According to FIG. 5A, it can be seen that in the film of Example 1, Sm, Fe and Co are uniformly distributed in the film. In addition, although a portion having a partially different concentration of Sm is observed, the portion is due to the existence of a second phase different from the main phase, and the second phase is shown in FIG. 4 (b). From the HAADF-STEM image, it is considered to be either _{a Th 2} Mn ₁₇ type phase (Sm ₂ (FeCo) ₁₇ phase) or a TbCu ₇ type phase (Sm (FeCo) _{7 phase).}
According to FIG. 5 (b), a non-uniform Co distribution was confirmed in the film of Example 2, and the portion where the Co concentration was low was explained with reference to the BF-TEM image shown in FIG. 3 (b). It was confirmed that it is the part corresponding to the grain boundary phase. Due to the detection limit of the energy dispersive X-ray analyzer, information on the distribution of B is not obtained from the STEM-EDS image.

6 (a) and 6 (b) show the results of analysis of the membrane of Example 1 by a three-dimensional atom probe (3DAP). The 3DAP map shown in FIG. 6A was obtained by scanning the probe parallel to the plane of the thin film. FIG. 6B shows a 3DAP map of Sm in the rectangular parallelepiped region in FIG. 6A (upper row), a composition profile of Sm, Co and Fe in the region (middle row), and Sm shown in the middle row. It is an enlarged version of the composition profile (lower).

According to FIGS. 6A and 6B, in the film of Example 1, the content ratio of Co and Fe based on the atomic% (at%) is 1: 4, and the composition ratio predicted from the film forming conditions. It was the value of. Regarding the Sm content, as explained with reference to the STEM-EDS image shown in FIG. 5 (a), a portion where the Sm concentration is partially different is observed, and the Sm content is about 10 in this portion. Although it was ~ 11 atomic%, the average content of Sm was about 7.95 atomic%, and this value is an anisotropic Sm (Fe _0.8 Co _0.2 ) _{12 having a ThMn 12 type structure.} It is consistent with the composition of Sm in the phase.

7 (a) and 7 (b) show the analysis result of the membrane of Example 2 by a three-dimensional atom probe (3DAP). The 3DAP map shown in FIG. 7A was obtained by scanning the probe parallel to the plane of the thin film. That is, the paper surface penetration direction in FIG. 7A is the plane perpendicular direction of the thin film. FIG. 7 (b) is a composition profile of the constituent elements (Sm, Co, Fe and B) in the region shown by the rectangular parallelepiped in FIG. 7 (a).

According to FIG. 7 (a), it can be seen that a large amount of B is distributed in the portion where the concentration of Co is low, which is explained with reference to the STEM-EDS image shown in FIG. 5 (b). This can also be clearly confirmed from the composition profile based on atomic% (at%) shown in FIG. 7 (b).
More specifically, according to FIG. 7 (b), B is distributed in both the main phase (1-12 phase) and the grain boundary phase (GB), and is particularly unevenly distributed in the grain boundary phase (uneven distribution). I understand that. In other words, in the film of Example 2, it was confirmed that a grain boundary phase having a distinctly different elemental composition from the main phase was present, and that the grain boundary phase was a B-enriched phase. The content of B in the grain boundary phase calculated based on the composition profile shown in FIG. 7 (b) was about 10.3 atomic%.

The average content of Sm in the main phase 1-12 phase (ThMn ₁₂ type crystal phase) was about 8.0 atomic%, which was consistent with the stoichiometric composition of the 1-12 phase. On the other hand, the average content of Sm in the grain boundary phase was slightly lower than that in the 1-12 phase, which was about 5.8 atomic%. As mentioned above, B is also distributed in the main phase, but its amount is negligibly low. The Co content was about 19.8 atomic% in the central portion of the 1-12 phase, but was about 10.6 atomic% in the grain boundary phase, which is the B-enriched phase.

As a result of ICP-OES analysis on the membrane of Example 2, the overall composition of the membrane was Sm _7.7 Fe _72.1 Co _16.5 B _3.7 (atomic%), and the main phase 1-12 phase. And the composition of the amorphous grain boundary phase is Sm _8.0 Fe _71.9 Co _19.8 B _0.3 and Sm _5.8 Fe _73.3 Co _10.6 B _10.3 (both atomic%). )Met. Further, since the total content of Fe and Co in the grain boundary phase is about 84%, it is considered that the grain boundary phase is ferromagnetic. In addition, as confirmed for Nd-Fe-B magnets, it is considered that the grain boundary phase can act as a barrier for domain wall movement, which is a factor showing a high coercive force of 1.2T. Then, in the film of Example 2, since all the phases are preferentially oriented in the [001] direction on the underlayer of V, it is considered that 1.50 T of residual magnetization was obtained.

FIG. 8 shows a schematic diagram of the fine structure of the thin films of Examples 1 and 2 derived from the results of the above-mentioned characteristic analysis. The left side of FIG. 8 is the thin film of Example 1 (Sm (Fe _0.8 Co _0.2 ) ₁₂ film, coercive force (μ ₀ H _c = 0.1 T)), and the right side is the thin film of Example 2 (Sm). (Fe _0.8 Co _0.2 ) ₁₂ B _0.5 film, coercive force (μ ₀ H _c = 1.2 T)).

In the film of Example 1, there may be a second phase having a Sm content partially different from that of the main phase, but in addition to the second phase, it is structurally and structurally clearly distinguished from the 1-12 phases of the main phase. There is no such phase.

On the other hand, in the film of Example 2, since it is a Sm (Fe, Co) ₁₂ film containing B, anisotropic crystal grains having a _{ThMn 12} type structure as the main phase (1-12 phase) are formed. The main phase is surrounded by an amorphous grain boundary phase having a certain thickness, and the grain boundary phase is a B-concentrated phase. Specifically, the film of Example 2 is a Sm (Fe, Co) ₁₂ film containing 3.7 atomic% B, so that a columnar 1-12 phase having an average grain size of about 40 to 50 nm is formed. The 1-12 phase is surrounded by an amorphous grain boundary phase having an average thickness of about 3 nm, and the grain boundary phase has a B concentration of about 10.3 atomic%. It is a chemical phase. By having such a fine structure, the film of Example 2 exhibited a coercive force of 1.2 T, which was not obtained with _{the conventional Sm (Fe, Co) 12 film containing no B.}

(Example 3)
Next, using the same material as in Example 2 and the film forming apparatus, a sample having a thickness of 100 nm and a B content of 0.5% by volume (Sm (Fe _0.8 Co _{0.). 2} ) ₁₂ B _0.5 film) was prepared. In Example 3, as the film forming conditions, _{the substrate temperature at the time of forming the Sm (Fe 0.8} Co _0.2 ) ₁₂ B layer is 350 ° C., and the temperature conditions after depositing the V cap layer (10 nm) are set. A total of 9 types of samples were prepared by setting as described in (a) to (i) below.
(A) Natural cooling (cooling to 60 ° C in 30 minutes)
(B) Cooling with nitrogen gas (cooling to 20 ° C in 4 minutes, then holding for 30 minutes)
(C) After holding at 200 ° C for 60 minutes, natural cooling (cooling to 50 ° C in 30 minutes)
(D) After holding at 250 ° C. for 60 minutes, natural cooling (cooling to 60 ° C. in 30 minutes)
(E) After holding at 300 ° C. for 60 minutes, natural cooling (cooling to 65 ° C. in 30 minutes)
(F) After holding at 350 ° C for 60 minutes, natural cooling (cooling to 65 ° C in 30 minutes)
(G) After holding at 400 ° C for 60 minutes, natural cooling (cooling to 70 ° C in 30 minutes)
(H) After holding at 450 ° C. for 60 minutes, natural cooling (cooling to 70 ° C. in 30 minutes)
(I) After holding at 500 ° C for 60 minutes, natural cooling (cooling to 80 ° C in 30 minutes)

Regarding the temperature conditions of (a) to (i) above, in the following, the condition (a) is set as a normal process (Normal process), and the conditions (b) to (i) are Sm (Fe _0.8 Co _{0. 2) 12} (based on 350 ° C.), of less than the temperature ((b) the substrate temperature for forming the layer B ~ (e)) the cooling conditions (cooling), more than the temperature ((f) ~ (I)) will be classified as heating conditions (Annealing), and the characteristics of the samples prepared under each condition will be described.

FIG. 9 shows the out-of-plane XRD pattern of the film obtained by the cooling conditions (Cooling) of the conditions (a) and the conditions (b) to (e).

According to FIG. 9, in each sample, a diffraction pattern derived from (002) and (004) (indicated by a downward white triangle mark in the figure) was observed, and the (001) plane was strongly oriented ThMn ₁₂ type. It was found that the crystal phase of Sm (Fe _0.8 Co _0.2 ) ₁₂ B layer is the main phase, that is, the crystal orientation of the Sm (Fe 0.8 Co 0.2) 12 B layer is preferentially oriented in the [001] direction. Moreover, in the comparison between the condition (a) and the conditions (b) to (e), no significant change was observed in the peak intensity and position of the main phase.

FIG. 10 shows the Out-of-Plane XRD pattern of the film obtained by the conditions (a) and the heating conditions (Annealing) of the conditions (f) to (i).

According to FIG. 10, in each sample, a diffraction pattern derived from (002) and (004) (indicated by a downward white triangle mark in the figure) was observed, and the (001) plane was strongly oriented ThMn ₁₂ type. It was found that the crystal phase of Sm (Fe _0.8 Co _0.2 ) ₁₂ B layer is the main phase, that is, the crystal orientation of the Sm (Fe 0.8 Co 0.2) 12 B layer is preferentially oriented in the [001] direction. Moreover, in the comparison between the condition (a) and the conditions (f) to (i), no significant change was observed in the intensity and position of the peak of the main phase. On the other hand, in the samples under the conditions (f) to (i), _{the peak of the second phase (α-Fe, Sm (FeCo) 7} , Sm ₂ (FeCo) ₁₇ phase, etc.) (indicated by the downward black triangle mark in the figure). There was a tendency for the peak intensity of to increase. In addition, the peak intensity of V used for the base layer and the cap layer tended to decrease. This suggests that the V layer is diffused in the main phase under the heating conditions (Annealing).

FIG. 11 shows the in-plane and anti-plane magnetization curves of the film obtained by the conditions (a) and the cooling conditions (Cooling) of the conditions (b) to (e).

According to FIG. 11, it was confirmed that all the samples had a coercive force of 1.16T (expressed as 11.6 kOe using the oersted unit Oe in the figure).

FIG. 12 shows the in-plane and in-plane magnetization curves of the film obtained by the conditions (a) and the heating conditions (Annealing) of the conditions (f) to (i).

According to FIG. 12, in the samples of the conditions (f) to (i), the coercive force value is lower than the value (11.6 kOe) obtained in the sample of the condition (a), and the higher the holding temperature, the lower the value. Was a low result.

13 (a) to 13 (d) show the magnetic properties (saturation magnetization, coercive force, residual magnetization ratio at zero magnetic field, and vertical anisotropy) of the samples under the conditions (c) to (i), respectively. Indicated.

According to FIG. 13 (a), the saturation magnetization ( _Ms ) tended to increase linearly as the holding temperature increased.
According to FIG. 13 (b), _{it was confirmed that the coercive force (H c} ) decreases under the heating condition (Annealing) in which the holding temperature is 350 ° C. or higher.
According to FIG. 13 (c), _{it was confirmed that the residual magnetization ratio (Mr} / M _s ) in a zero magnetic field decreases under heating conditions (Annealing) in which the holding temperature is 350 ° C. or higher, similar to the coercive force. rice field.
According to FIG. 13 (d), the perpendicular anisotropy _{(K u),} although the holding temperature is increased at 350 ° C. or more heating conditions (Annealing), the holding temperature is decreased at the 500 ° C. Conditions (i) Was confirmed.

From these results and the results of the coercive force under each condition described with reference to FIGS. 11 and 12, the above-mentioned amorphous grain boundary phase (B-concentrated phase) is Sm (Fe _0.8 Co _{0). .2)} it is formed has been suggested during the formation of the 12 _B layer. Further, in the method for manufacturing a magnetic recording medium according to an embodiment of the present invention, when a magnet layer containing a rare earth compound according to the embodiment of the present invention is formed on a substrate by a sputtering method, the temperature of the substrate is set to a predetermined temperature. It was confirmed that it is desirable to cool the target at a temperature lower than the heating temperature of the substrate after sputtering the target containing the elements constituting the rare earth compound.

(Examples 4 to 12)
Next, using the same material as in Example 2 and the film forming apparatus, the thickness of the film was changed from 20 nm to 200 nm, and the sample having a B content of 5.0% by volume (Sm (Fe _{0)). .8} Co _0.2 ) ₁₂ B film) was prepared.

(Examples 13 to 21)
Further, using the same material as in Example 2 and the film forming apparatus, the amount of introduction (target) of boron (B) is changed from 0% by volume to 8.0% by volume, and the sample has a thickness of 50 nm. (Sm (Fe _0.8 Co _0.2 ) ₁₂ B film) was prepared.

Table 1 below summarizes the amount of boron (B) introduced (target) in the samples of Examples 4 to 21 and the results of elemental analysis by ICP-OES.

In Table 1, "-" indicates that ICP-OES measurement was not performed. In addition, "at%" in the table represents atomic%.

FIG. 14 shows the Out-of-Plane XRD patterns of the thin films of Examples 4, 7 to 10 and Example 12 in which the thickness of the film was changed from 20 nm to 200 nm.
The intensity of the diffraction lines from (002) and (004) (indicated by the downward white triangle mark in the figure) increases as the thickness of the film increases. On the other hand, the intensity of the diffraction line (indicated by the downward black triangle mark in the figure) from the second phase (α-Fe, Sm (FeCo) ₇ , Sm ₂ (FeCo) _{17) observed near 65 ° is} It did not change depending on the thickness of the film.

FIG. 15 shows changes in the lattice constants a, c, and c / a of the thin films of Examples 4 to 12 in which the film _{thickness (t SFC) was changed.} a increased to 8.70 Å at 30 nm, then showed a constant value of 8.70 Å up to 80 nm, and then decreased.
On the other hand, c decreased to 4.75 Å at 30 nm and then showed a constant value of 4.75 Å. c / a decreased to 30 nm, but then showed a constant value.

FIG. 16 shows the magnetization curves of the thin films of Examples 4, 7 to 10 and Example 12 in which the thickness of the film was changed. Both were perpendicularly magnetized films. The coercive force in the vertical direction increases with the increase in the thickness of the film, and shows a maximum value of 0.94 T at 100 nm. This value was the largest value among the films of Examples 4 to 21.

When the thickness of the film is 100 nm, the magnetization value is 1.55 T, which is a relatively high value. When the B concentration was further increased, the coercive force decreased.

FIG. 17 shows the magnetization of a thin film with a constant B content (M _s ), the residual magnetization ratio in a zero magnetic field ( _Mr / M _s ), the coercive force (H _c ), and the vertical anisotropy ( _Ku ). Dependence on the thickness of the film (t _SFC ) was shown.

FIG. 18 shows the temperature dependence of the coercive force. In the figure, the horizontal axis represents temperature (K) and the vertical axis represents coercive force (T). For the measurement of the temperature dependence of the coercive force, a SQUID-VSM equipped with a sample vibration type magnetometer was used, and the temperature was changed in the range of 50 to 700 K.
The black circle plot shows the sample obtained by adjusting the film formation conditions so that the boron content is 5% by volume, and the black square plot shows Cu diffusion _{into Sm (Fe 0.8} Co _0.2 ) _12. The black triangular plot is Sm (Fe _0.8 Co _0.2 ) ₁₂ with Cu-Ga diffusion.
A total of three black solid lines and two black broken lines are commercial Nd ₂ Fe ₁₄ B magnets (NMX-36 (Dy content 8%), NMX-43 (Dy content 4%), NMX-52 (Dy). Free)) shows the temperature dependence. According to FIG. 18, it can be seen that the Sm (FeCo) ₁₂ magnet exhibits a smaller temperature dependence than the _{Nd 2} Fe _{14 B magnet.} Among the Sm (FeCo) ₁₂ magnets, the sample containing 5% by volume of boron shows a high coercive force over the entire temperature range, and the temperature coefficient (β) of the coercive force in the temperature range of 300 to 700 K is the lowest value (-0. It was 15% / K).

The thin film of Example 2 was described with reference to the microstructure as described with reference to FIGS. 4 (c) and 4 (d), and FIGS. 5 (b), 7 (a) and (b). The distribution and composition profile of such constituent elements were also confirmed from the HAADF-STEM image of the thin film of Example 9, the nanobeam electron diffraction pattern, the STEM-EDS image, and the analysis results by 3DAP (data not shown).

Figure 19 showed the Out-of-Plane XRD pattern of _{Sm (Fe 0.8 Co 0.2) 12} B film of Example 13 and Example 15-21 were the 50nm thick film. From the results of Out-of-Plane XRD, diffraction patterns derived from (002) and (004) (indicated by downward white triangles in the figure) were observed, and the (001) plane of the _{ThMn 12 type crystal phase was observed.} It was found that the crystals were strongly oriented, that is, _{the crystal orientation of the Sm (Fe 0.8} Co _0.2 ) ₁₂ B layer was preferentially oriented in the [001] direction.

Here, when the content of B is increased (002), the diffraction line from (004) shifts to the high angle side, and is caused by _{α-Fe, Sm (FeCo) 7} , and Sm ₂ (FeCo) _17. A peak (indicated by a downward black triangle in the figure) was observed near 2θ = 65 °. Moreover, the strength increased with the increase of the content of B.

In FIG. 20, the lattice constants a, c, and c / a (vertical axis) and the content of B (horizontal axis) of the thin films of Examples 13 to 18 and Example 21 in which the film thickness is constant (50 nm) are shown. The relationship with was shown. a showed an almost constant value at about 8.56 Å (angstrom) up to a content of B up to 4% by volume, but increased sharply when B exceeded 4% by volume. On the other hand, c and c / a decreased monotonically with the increase in the content of B.

FIG. 21 shows the in-plane and in-plane magnetization curves of the thin films of Examples 13 and 15 to 21 in which the B content was changed while the film thickness was constant.
According to FIG. 21, when B is 0% by volume, it shows strong vertical anisotropy. The coercive force is 0.14T.
According to FIG. 21, the coercive force in the direction perpendicular to the plane increased with the increase in the content of B, and the coercive force of 0.75 T was shown when the content of B was 5.0% by volume. Moreover, in the above, it showed a high magnetization of 1.59T. Furthermore, the in-plane component also increased.

Further, when the content of B was increased, the coercive force in the direction perpendicular to the plane decreased, and the direction in which the magnetization was easy changed into the plane.
In FIG. 22, the magnetization of the thin film having a constant film thickness (M _s ), the residual magnetization ratio in a zero magnetic field ( _Mr / M _s ), the coercive force (H _c ), and the vertical anisotropy (K) are shown. _{The dependence of u} ) on the B content was shown. In addition, FIG. 22 also shows the results of the samples prepared under the film forming conditions different from the B introduction target amount of Examples 13 to 21.

Since the rare earth compound of the present invention has an excellent coercive force, it is particularly suitable as a medium material used for a perpendicular magnetic recording medium for a perpendicular magnetic recording disk mounted on a magnetic disk device such as an HDD. In addition, as a medium material for discrete track media (DTM) and bit-patterned media (BPM), which are promising media for achieving ultra-high recording densities that exceed the information recording densities of current perpendicular magnetic recording media. Alternatively, it is particularly preferably used as a medium material for heat-assisted magnetic recording (HAMR) that can achieve an ultra-high recording density that further exceeds the information recording density by the perpendicular magnetic recording method.

10 Magnetic recording medium (perpendicular magnetic recording medium)
20 Substrate 30 Heat absorption layer 40 Underlayer (orientation control layer)
50 Magnetic layer (perpendicular magnetic recording layer)

Claims (10)

The overall composition is Formula 1: (R ¹ _(1-x) R ² _x ) _a T _b M _c
(In the above formula, R ¹ is at least one element selected from the group consisting of Sm, Pm, Er, Tm, and Yb, and R ² is Zr, Y, La, Ce, Pr, Nd. , Eu, Gd, Tb, Dy, Ho, and at least one element selected from the group consisting of Lu, and T is at least one element selected from the group consisting of Fe, Co, and Ni. M is boron, x is a number from 0 to 0.5, a is a number from 6.0 to 13.7 atomic%, and c is greater than 0 atomic% and 12 atomic% or less. B is a number represented by 100-ac atomic%.)
A rare earth compound having a ThMn _12- type crystal phase as a main phase and having an amorphous grain boundary phase existing between the main phases. The rare earth compound according to claim 1, wherein the absolute value of the difference between the boron content in the main phase and the boron content in the grain boundary phase is 1.0 atomic% or more. The rare earth compound according to claim 1 or 2, wherein R ^{1 contains at least Sm.} The rare earth compound according to any one of claims 1 to 3, which contains substantially none of Ti, V, Mo, Nb, Cr, and W. The rare earth compound according to any one of claims 1 to 4, wherein T is Fe and Co. The rare earth compound according to any one of claims 1 to 5, wherein at least one selected from the group consisting of a crystal orientation and an easy magnetization axis is preferentially oriented along a predetermined direction. The rare earth compound according to any one of claims 1 to 6, wherein a is 6.0 to 10.0 atomic%. A magnetic recording medium using the rare earth compound according to any one of claims 1 to 7 in the magnetic layer. The magnetic recording medium according to claim 8, wherein the crystal orientation of the magnetic layer is preferentially oriented in the [001] direction. A method for producing a magnetic layer containing the rare earth compound according to any one of claims 1 to 7 on a substrate by a sputtering method.
^{The substrate is heated to 250 to 400 ° C. to sputter a target containing the elements (R 1} , R ² , T, and M) constituting the rare earth compound, and then cooled at a temperature lower than the substrate heating temperature. Method.

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