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ISSN 2432-0250

Vol.44 No.5
JournalJournal of the Magnetics Society of Japan
Electronic Journal URL: https://www.jstage.jst.go.jp/browse/msjmag

2020

Thin Films, Fine Particles, Multilayers, Superlattices
Magneto-optical Imaging Using Bismuth-substituted Iron Garnet Films Prepared by Metal‒Organic Decomposition
発行（公社）日本磁気学会 〠101‒0052 東京都千代田区神田小川町2丁目8番地
三井住友海上小川町ビル5階☎(03)5281-0106
印刷（株）国 際 文 献 社 〠169‒0075 東京都新宿区高田馬場3の8の8☎(03)3362-9741

T. Ishibashi…108

Magnetic Recording
MnGa Film with (001) Texture Fabricated on Thermally Oxidized Si Substrate Using CoGa Buﬀer Layer
Y. Miwa, D. Oshima, T. Kato, and S. Iwata…117

Magnetic Recording
Media Design for Three-Dimensional Heat-Assisted Magnetic Recording

広告◇東英工業株式会社T. Kobayashi, Y. Nakatani, and Y. Fujiwara

…122

44

HP: http://www.magnetics.jp/ e-mail: msj@bj.wakwak.com
Electronic Journal: http://www.jstage.jst.go.jp/browse/msjmag

5

2020

INDEX

Journal of the Magnetics Society of Japan
Vol. 44, No. 5
Electronic Journal URL: https://www.jstage.jst.go.jp/browse/msjmag

CONTENTS
Thin Films, Fine Particles, Multilayers, Superlattices
Magneto-optical Imaging Using Bismuth-substituted Iron Garnet Films Prepared by
Metal–Organic Decomposition………………………………………………………………………T. Ishibashi

108

Magnetic Recording
MnGa Film with (001) Texture Fabricated on Thermally Oxidized Si Substrate Using CoGa Buffer Layer…
…………………………………………………………………………Y. Miwa, D. Oshima, T. Kato, and S. Iwata

117

Magnetic Recording
Media Design for Three-Dimensional Heat-Assisted Magnetic Recording……………………………………………………………………… T. Kobayashi, Y. Nakatani, and Y. Fujiwara

Board of Directors of The Magnetics Society of Japan
President:
K. Nakagawa
Vice Presidents:
S. Sugimoto, S. Matsunuma
Directors, General Affairs:
Y. Miyamoto, H. Saito, H. Yuasa
Directors, Treasurer:
K. Ishiyama, H. Takahashi
Directors, Planning:
S. Nakagawa, T. Kondo
Directors, Editorial:
T. Ono, T. Kato
Directors, Public Relations:S. Greaves, S. Sakurada
Directors, International Affairs:
M. Nakano, H. Yanagihara
Specially Appointed Director, Gender Equality:
F. Akagi
Specially Appointed Director, Societies Collaborations: K. Fujisaki
Specially Appointed Director, International Conferences: Y. Miyamoto
Auditors:
R. Nakatani, Y. Takano

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J. Magn. Soc. Jpn., 44, 108-116 (2020)<Review>

Magneto-optical Imaging Using Bismuth-substituted Iron Garnet Films
Prepared by Metal–Organic Decomposition
Takayuki Ishibashi

Nagaoka University of Technology, 1603-1 Kamitomioka, Nagaoka, Niigata 940-2188, Japan

Magneto-optical (MO) imaging using an MO material is one of the most important imaging techniques formeasuring magnetic field distributions. The technique is attracting attention because of its ability to visualize
magnetic fields and current distributions at a wide range of temperatures from cryogenic to room temperature or
higher, a wide range of length scales from micrometers to several tens of centimeters, and a wide range of frequencies
from DC to GHz. Nevertheless, further development of films with good MO properties is required in order to extend
the range of applications of MO imaging. In this paper, developments related to bismuth-substituted iron garnet films,
R3−xBixFe5−yGayO12 (R = Y, Nd; 1 ≤ x ≤ 3), prepared by metal–organic decomposition for MO imaging plates are reviewed.
In addition, an MO imaging plate with an electroluminescent backlight, which enables wide-area MO imaging, andthe MO color imaging technique, which indicates the strength of a magnetic field using colors, are reviewed.
Key words: magneto-optical imaging, Faraday effect, Bi-substituted iron garnet, metal-organic decomposition
1. Introduction

2. MO imaging using Bi-substituted iron garnet

Magneto-optical (MO) imaging has played an
important role in both basic research on magnetic
materials and applications such as magnetic recording
technologies. The polarizing microscope developed in the
1950s has been used to visualize magnetic domains inmagnetic materials via MO effects1). Magnetic domains
become visible under a polarizing microscope because thepolarization plane of the light from a sample is rotated asa result of the MO Kerr effect2) in reflection or the
Faraday effect3) in transmission.
MO imaging techniques that utilize an MO material
measure the magnetic field distribution in a magnetic
material, whereas the aforementioned MO microscopy
technique measures the magnetization. We refer to theplate with MO materials that enable visualization ofmagnetic fields as an MO imaging plate. The magnetic
field distribution at a location on the MO imaging plateis visualized in real-time using the Faraday effect of theMO materials, as shown in Fig. 1. The MO imaging platesconsist of an MO material (typically a bismuth (Bi)substituted iron garnet film prepared on a transparent
substrate as described below), a mirror layer, and aprotection layer. The spatial resolution in MO imaging
can be as high as the optical diffraction limit, although it
is adversely affected by the dispersive characteristics ofthe magnetic field. However, MO imaging would enable
the measurement samples with a large area (larger thanseveral tens of square centimeters) if large MO imaging
plates become available. Given these features, MO
imaging is promising as a nondestructive measurement
technique for characterizing electronic devices, detecting
impurities and defects in metals. Moreover, it isconsidered a promising technique for characterizing
magnetic fields with frequencies as high as the
microwave range4),5).

Alers introduced the MO imaging technique usingparamagnetic cerium phosphate glass in 1959 to study
the intermediate state in superconductors. The glass waslater replaced by Eu-chalcogenides such as EuSe or EuS,
which were used in a narrow temperature range in which
both a large Faraday rotation angle and paramagnetism
were utilized6),7). MO imaging techniques using Bisubstituted iron garnets were first reported in 19898).
This approach enables the measurement of magnetic
fields in a wide temperature range (from extremely low
temperatures to a high temperature of ~100°C) using MO
imaging9),10). In the early stages of the development of
MO imaging, the magnetic domain size in garnet films
with perpendicular magnetic anisotropy limited thespatial resolution to the lengths of the magnetic domains.
The use of Bi-substituted iron garnet, which exhibits inplane magnetization, has led to a substantial
improvement of the spatial resolution11). Consequently,
the real-time observation of single vortices insuperconductors has been realized12)–14).
The magnetic garnets have a garnet structure withcubic space group Ia3d. Their structure includes threetetrahedral and two octahedral sites for Fe3+ ions andthree dodecahedral sites for R3+ rare-earth ions in one
unit cell described by the chemical formula R3Fe5O12.
Each of the Fe3+ ions in the two sites has a magnetic
moment of 5 µB with directions opposite to each other,resulting in a ferrimagnet with a relatively weak
spontaneous magnetization. Most magnetic garnets have
similar Curie temperatures of approximately 550 K,
although some rare-earth elements have a magnetic
moment induced by the moment of Fe3+, aligned parallelor antiparallel to it depending on the rare-earth element.
For MO imaging, yttrium or lutetium iron garnets havebeen used because Y3+ and Lu3+ do not have magnetic

108

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INDEXepitaxial films prepared on Gd3Ga5O12 (GGG) single
crystals with a lattice constant of 1.2383 nm have beenused as substrates. The saturation magnetic field wascontrolled by substituting nonmagnetic elements such as
Ga and Al for Fe, resulting in a decrease in the
demagnetization field. In addition, magnetostrictive
effects, which are affected by the lattice matching andwhose values differ with the kind of rare-earth element,
must also be considered for controlling the magnetic
anisotropy.
To obtain Bi-substituted iron garnet films with a large
proportion of Bi (1.5 ≤ x ≤ 3), researchers have used awide variety of film growth techniques capable of growing
films under nonequilibrium conditions, including
sputtering, pulsed laser deposition29)–31), ion-beam
sputtering29), and the sol–gel method17),32). The resultant
films, however, often included secondary phases or
exhibited cracks. However, highly Bi-substituted irongarnet films suitable for optical applications have beensuccessfully grown by metal–organic decomposition(MOD)33),34), as described in section 3.
3. Preparation of Bi-substituted iron garnet films
by MOD

Fig.. 1 (a) A schematic of MO imaging using an MO
imaging plate and (b) the magnetic-field dependenceof the Faraday rotation of the MO imaging plate.

moments and, consequently, there is no magnetic
compensation temperature and less temperaturedependence of the magnetizations.
One of the most important properties ofBisubstituted iron garnets is their excellent MO properties,
which can be improved by substituting Bi3+ for R3+. The
Faraday rotation for Y3−xBixFe5O12 (Bi:YIG) increases
from 4°/µm for x = 1 to 30°/µm for x = 3 (Bi3Fe5O12)15)–17).
Because of their excellent MO properties,
Bisubstituted iron garnets have been used in various MO
applications, including optical isolators18), spatial light
modulators19),20), and MO imaging plates. Moreover, they
are promising materials for studying fast magnetization
reversals using a femtosecond laser21), spin-wave
phenomena22)–24), the spin Seebeck effect25)–27), andterahertz emission28).
In addition to the large Faraday rotation, the magnetic
anisotropy of an MO material must be controlled for MO
imaging. In-plane magnetic anisotropy with a linear
magnetic-field dependence of Faraday rotation issuitable for MO imaging. From the viewpoint ofmagnetocrystalline anisotropy, (001)- or (011)-orientated

3.1 MOD method
MOD is a film-growth method for oxides and offers
several advantages:- The process is simple and does not require a vacuum
system.
- Chemical composition can be precisely controlled.
- MOD solutions are uniform at the molecular leveland stable for several years.
- It enables the easy preparation of large-scale films.
Figure 2 shows a typical MOD procedure consisting of
spin-coating of a MOD solution (3000 rpm, 60 s), drying(100°C, 10–30 min), pre-annealing to decompose theorganic materials (~400°C, 10–30 min), and annealing to
induce crystallization (700°C, 3 h). All of these thermal
treatments are performed in air. We used MOD solutions
prepared by Kojundo Chemical Co., Ltd. in our study.
The starting materials for the MOD solutions consist of

Fig. 2 A typical procedure in the MOD process.

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solutions of metal octylates, including those of Bi, Y, Fe,
Nd, and Ga. They are mixed to obtain the desired
chemical proportion. The total concentration of octylates
in the MOD liquids is usually fixed at 3–4%. The typical
thickness of a single coated film is 30–40 nm.
3.2 Bi:YIG films
Figure 3 shows XRD patterns for Y0.5Bi2.5Fe5O12(Bi2.5:YIG) and Bi3Fe5O12 (BIG) films prepared on GGG(001). Only the 004 and 008 peaks are observed to the leftof the peaks for the GGG substrate for both the YBIG and
the BIG film, indicating that the Bi2.5:YIG and BIG films
were both grown along the [100] direction. However, thefilms exhibited a granular structure consisting of grains
smaller than 100 nm. Fortunately, because theirmagnetic domains are as small as their grains, thesematerials are suitable as materials for MO imaging inwhich the spatial resolution is limited by the magnetic
domain structure. In fact, large zigzag domain structures,
which have been observed in garnet films grown by theliquid-phase epitaxy, have not been observed in MODgrown garnet films.
Figure 4 shows the magnetic-field dependence of the
Faraday rotation for Bi:YIG films with 1.5 ≤ x ≤ 3 on

GGG(001) substrates, measured at a photon energy of 3
eV. A linear magnetic-field dependence of the Faraday
rotation is observed below the saturation field ofapproximately 2 kOe for all of the samples represented
in the figure.
Moreover, off-diagonal components of permittivity
tensors were obtained35) for the first time, as shown in
Fig. 5. The permittivity tensor 𝜀𝜀̃ was defined as𝜀𝜀%𝜀𝜀̃ = $𝑖𝑖 ∙ 𝜀𝜀)0−𝑖𝑖 ∙ 𝜀𝜀)
𝜀𝜀%
0

0
0 ,.
𝜀𝜀+

(1)

These results indicate that Bi:YIG films prepared by
the MOD method exhibited excellent MO properties andsufficient optical quality for optical analysis using
ellipsometry.
3.3 Bi:NIG on GGG
As previously mentioned, Y3+ or Lu3+, which haverelatively small ionic radii, have been used in Bisubstituted iron garnets to minimize lattice mismatch
between the garnet films and the GGG substrates.
However, other rare-earth elements could be used if
lattice matching is not required. In one study, Nd3+ waschosen because its ionic radius is similar to that of Bi3+and its magnetic moment is parallel to the net magnetic
moment of Fe3+ in garnets, and Nd0.5Bi2.5Fe5−yGayO12(Bi2.5Ga:NIG) films were successfully prepared by
MOD36),37). Figure 6 shows magnetization curves of
Bi2.5Ga:NIG films prepared on GGG(001) substrates,

Fig. 3 XRD patterns for (a) B2.5:YIG and (b) BIG films
prepared on GGG(100) substrates.

Fig. 4 Faraday rotation hysteresis loops of Y3-

xBixFe5O12 films on GGG substrates, as measured at 3
eV.35)

110

Fig. 5 (a) Real and (b) imaginary parts of offdiagonal elements of the permittivity tensor for Y3xBixFe5O12 (x = 1.5, 2, 2.5, 3.0) prepared on
GGG(001) substrates 35).

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INDEX

with decreasing Nd content, suggesting a large
contribution of Nd3+ to the Ku. Aplesnin et al.38) reported
on the temperature dependence of the magnetostrictionin Bi2.5:NIG films. They found a large temperaturedependence of the magnetostriction effect in Bi2.5:NIG,
indicating that the magnetic anisotropy could be variedwith temperature.



3.4 Bi,Ga:NIG on glass substrate
Increasing the measurement area in MO imaging
necessitated the development of highly Bi-substituted
iron garnet films on glass substrates. However, thecrystallization of highly Bi-substituted iron garnets with
Bi contents in the range 2 ≤ x ≤ 3, which are metastablematerials, was found to be difficult on glass substrates.
Okuda successfully grew a polycrystalline BIG film on apolycrystalline GGG layer prepared on a fused-quartz
substrate at annealing temperatures ranging from 700 to
800°C. However, some secondary phases existed and theannealing temperature was too high for general glass
substrates. Meanwhile, our group successfully preparedpolycrystalline BIG films on Bi1:YIG buffer layers39) andBi2.5Ga:NIG on Bi1:NIG buffer layers40),41). These films
were prepared at annealing temperatures of 600–
700ºC42), which were sufficiently low to enable the use ofglass substrates. A photograph of a Bi2.5:NIG film
prepared on a glass substrate (Eagle XG, Corning) with





Fig. 6 Magnetization curves for Nd0.5Bi2.5Fe5−yGayO12
films with (a) y = 0, (b) 0.5 and (c) 1.0 on GGG(001)substrates36).

Nd0.5Bi2.5Fe5O12 film withφ 6 inch size glass sub.

Fig. 8 Photo of a ⌀6 inch Bi2.5NIG film.

Fig. 7 Uniaxial magnetic anisotropy Ku andmagnetocrystalline anisotropy K1 in Y3-xBixFe5O12 andNd3-xBixFe5O12 films.

showing that the magnetic anisotropy was controlled bysubstituting Ga3+ for Fe3+. The effect of Nd3+ on themagnetic anisotropy was carefully analyzed using
ferromagnetic resonance (FMR) measurements37). Figure
7 shows the uniaxial magnetic anisotropy Ku andmagnetocrystalline anisotropy K1 in Bi:YIG and Bi:NIG.
The values of K1 were found not to vary with the Bi
content for either the Bi:YIG or the Bi:NIG films, and thevalue of Ku in Bi:YIG also did not varied with the Y
content. By contrast, the value of Ku in Bi:NIG decreased

Fig.
9
Faradayhysteresis
loops
for
Nd0.5Bi2.5Fe5−yGayO12 films with y = 0, 0.5 and 1 and a
thickness of 450 nm prepared on glass substrates.

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Fig. 10 Polar angle dependence of the FMR field for
Nd0.5Bi2.5Fe5−yGayO12 thin films on glass substrates.

Fig. 12 Frequency dependence of MO signals for
Nd0.5Bi2.5Fe5−yGayO12 thin films on glass
substrates with y = 0, 0.25, 0.5, 0.75 and 1 andBi1:YIG on GGG (111) substrate.

approximately the same as those measured via torquemeasurements41). The 𝑔𝑔 -factor values were obtainedusing the equation, 𝑔𝑔 = 𝛾𝛾ℎ⁄2𝜋𝜋𝜇𝜇K , where 𝜇𝜇K is the Bohrmagneton and ℎ is Planck’s constant, and are plotted asa function of the Ga content in Fig. 11. The 𝑔𝑔 -factor
slightly increased with increasing Ga content. Thisbehavior can be explained by the existence of the orbital
angular momentum 𝜇𝜇MNO in Nd3+ in addition to the spin
moment 𝜇𝜇PQRS, as given46) by
Fig. 11 The g-factor of Nd0.5Bi2.5Fe5−yGayO12 thin films
with y = 0, 0.25, 0.5, 0.75 and 1 prepared on glass
substrates.

a diameter of 6 inches is shown in Fig. 8.
The Bi2.5Ga:NIG films exhibited Faraday rotation
angles comparable to those for films on single-crystalline
GGG substrates, and the magnetic anisotropy wascontrolled by substituting Ga3+ for Fe3+, as shown in Fig.
943).
Magnetic anisotropy and high-frequency properties
have been studied by FMR measurements44). The
resonant magnetic field Hres measured for X-band (9.09
GHz) microwaves is plotted as a function of the polar
angle 𝛽𝛽 in Fig. 10 with the coordinate system used in the
FMR experiment. Open circles represent theexperimental data, and solid lines indicate fitted datausing the equation45)
𝜔𝜔 )𝐾𝐾B
0 3 = 4𝐻𝐻678 cos(𝛽𝛽 − 𝜃𝜃) + 2 0 − 2𝜋𝜋𝑀𝑀8 3 cos 2𝜃𝜃E
𝑀𝑀8
𝛾𝛾
𝐾𝐾B
× 4𝐻𝐻678 cos(𝛽𝛽 − 𝜃𝜃) + 2 0 − 2𝜋𝜋𝑀𝑀8 3 cos) 𝜃𝜃E,
𝑀𝑀8

(2)

where 𝜔𝜔 is the angular frequency of the microwaves, 𝜃𝜃is the angle between M and the [001] direction, and Ms isthe saturation magnetization. Ku and the gyromagnetic
ratio g were estimated by fitting the 𝛽𝛽 dependence ofHres. The Ku values obtained in this analysis were112

𝑔𝑔 = 2 T1 +

𝜇𝜇MNO
V.
𝜇𝜇PQRS(3)

Similar behavior of the 𝑔𝑔 -factor around thecompensation
point
has
been
reportedfor
Eu3Fe5−xGaxO12, although the sign of the 𝜇𝜇PQRS was
opposite to that in this case because the direction of themagnetic moment in Eu3+ was opposite to that in Nd3+.
From the viewpoint of the magnetic moment of each ion,
the effective gyromagnetic ratio 𝛾𝛾YZZ is expressed47) by𝛾𝛾YZZ =

𝑀𝑀[\ + (𝑀𝑀]Y (𝑡𝑡𝑡𝑡𝑡𝑡. ) − 𝑀𝑀]Y (𝑜𝑜𝑜𝑜𝑜𝑜. )),𝑀𝑀[\ 𝑀𝑀]Y (𝑡𝑡𝑡𝑡𝑡𝑡. ) − 𝑀𝑀]Y (𝑜𝑜𝑜𝑜𝑜𝑜. )+|𝛾𝛾[\ |
|𝛾𝛾]Y|(4)

where 𝑀𝑀[\ is the magnetic moment of Nd3+, 𝑀𝑀]Y (𝑡𝑡𝑡𝑡𝑡𝑡. )and 𝑀𝑀]Y (𝑜𝑜𝑜𝑜𝑜𝑜. ) are the magnetic moments of Fe3+ intetrahedral and octahedral sites, respectively, and 𝛾𝛾[\and 𝛾𝛾]Y are the gyromagnetic ratios of the rare-earth
sublattice and net Fe3+-ion sublattice, respectively. The
decrease of the 𝑔𝑔-factor with increasing Ga content is
explained by the decrease of the net magnetic moment
𝑀𝑀]Y (𝑡𝑡𝑡𝑡𝑡𝑡. ) − 𝑀𝑀]Y (𝑜𝑜𝑜𝑜𝑜𝑜. ) with increasing Ga substitution.
These results are consistent with the magnetic momentof the Nd3+ ions being parallel to the net magnetic
moment of the Fe3+ ions, and indicate that the Nd3+ ionsplay an important role in determining magnetic
properties such as the 𝛾𝛾YZZ even though the Nd content
was as low as 0.5. Because of the larger 𝛾𝛾YZZ , themagnetization in Bi,Ga:NIG could respond at higher

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INDEX

frequencies.
MO signals measured at high frequencies using
Bi2.5Ga:NIG films on glass substrates are shown in Fig.
12. The AC magnetic field induced by the AC current
with frequencies as high as 6 GHz was measured at theedge of a microstrip line by the stroboscopic method using
short laser pulses without a bias magnetic field48),49).
Figure 12 shows signal voltages corresponding to MO
responses measured using Bi2.5Ga:NIG films with a Ga
content 0 ≤ y ≤ 1 on glass substrates. For comparison, theresult for a Bi1:YIG thin film on a GGG(111) substrateprepared by MOD is also shown. It should be noted thatthese samples were measured with the same condition.
MO signals were detected at frequencies as high as 6 GHz
for all of the Bi2.5Ga:NIG thin films, and they had broad
4.2 磁気転写特性の評価
resonance peaks between 1.5 and 2 GHz. By contrast, the4.2.1 磁気転写膜の性能指数
MO signal measured using a Bi:LIG film grown by the
BYIG と BIG について，磁気転写膜としての特性を比較
LPE method showed a sharp resonance peak49).
するために，
それぞれ最も磁気光学効果の大きかった本焼成温度が 700℃で作製された本焼成温度 BYIG と 490℃で作製された

わかる．

4. Mの磁気光学スペクトルと透過率をO imaging using Bi-substitFig.7
utedに示す．膜厚はそれ
iron garnet films
BIG

ンプルの AFM

ぞれ 200nm である．
ファラデー回転のピーク位置での透過率

prepared by MOD

4.は，それぞれ約
1 MO imaging40%，30%であった．
using Bi:YIG on GGG(001)
Figure13shows
MO images of magnetic recording
た薄膜について
磁気転写膜としての性能を比較するために，性能指数
marks
invideotapes
measured， using a Bi2.5:YIG film
いが，その理34). The wavelengths of thewithathickness
of
200
nm
13)
を用いて評価した ．ここで T は透過率である．この式からグレイン程度のmeasuredrecording marks
Fig. 13(a) and 13(b) were 4
性能指数を計算した結果をFig.8inに示す．この結果から性能指数
非常に小さくand 1.16 µm, respectively. Zigzag-shaped magnetic
はファラデー回転角が最大値をとる波長ではなく，BYIG ではと考えている．
domains recorded in the videotapes were clearly observed.
530 nm, BIG では 550 nm 付近で最も大きな値を示すことがわ
界に不純物層This result indicates that the spatial resolution in MO
かった．また，530〜650 nm の波長において BIG の方が大と予想していimaging using the Bi2.5:YIG film is better than 0.6 µm,
きな性能指数を示した．
which is sufficient for measurements using an optical
4.2.2 空間分解One
能の評価
は，磁気転写膜 microscope.
of the advantages of highly Bi空間分解能を評価するためにサブミクロンサイズの磁気記録テ磁化が得やすsubstituted iron garnet is its high submicron-scale
ープの観察を行った結果をFig.
Fig. 9の(a)，れた特徴を有spatialresolution, which 9に示す．
is achievedby a(b)はBYIG
submicronで，(c)，BIGを用いて観察を行った結果である．
thickfilm(d)はwith
a large
Faraday rotation angle.(a)，(b)，イン構造になっ14. A camera and an analyzer were used to acquire theMO signals50). In addition, a quarter-wave plate wasplaced in front of the analyzer for quantitative
measurement. Figure 15 shows the magnetic field
distribution for a smart phone. The magnetic field
distribution from permanent magnets used in the camera,
speakers, and the taptic device were clearly observed.
The magnetic field distributions were quantitatively
obtained by measuring the MO effect, rotation angle 𝜃𝜃, or
ellipticity 𝜂𝜂 using equations50),51)
2𝐼𝐼gh − (𝐼𝐼ijh + 𝐼𝐼gjh )2(𝐼𝐼ijh + 𝐼𝐼gjh )𝐼𝐼ijh − 𝐼𝐼gjh𝜂𝜂 ≈2(𝐼𝐼ijh + 𝐼𝐼gjh )

𝜃𝜃 ≈

(5)

(6)

where 𝐼𝐼mn , 𝐼𝐼opn , and 𝐼𝐼mpn are the light intensities forlinear and right- and left-circular polarization,
respectively, measured at each pixel of the imaging
sensor in the imaging device. Quantitative values of themagnetic field were obtained from the magnetic-fielddependences of 𝜃𝜃 or 𝜂𝜂. An MO image of a stack of Ndmagnets and the corresponding line profile are shown in
Fig. 16. The validity of the measured values has been
discussed47).

(c)，(d)の記録波長はそれぞれ 4 m，1.16 m，
1.54 m，0.9 m，4.である．記録波長は図中の明るい線と暗い線の一組の幅であるの2 MO imaging plate combined with backlight
An optical setup with extraordinarily large lenses andで，空間分解能は記録波長の約半分となる．したがって今回の測
a 定から
half-mirror
be required to achievesufficiently
BYIG とwill
BIG の空間分解能はそれぞれ0.58 m，0.45
uniform polarization over a large area. To solve thism よりも高いことがわかる．このサブミクロンサイズの空間
problem, Nagakubo et al. have proposed combining an
分解能は，今回観察に使用した光学系の分解能とほぼ等しい空間
MO imaging plate with a thin electroluminescent (EL)
分解能である．しかし，問題点として得られた磁気イメージにざ
backlight and a linear polarizing film, as shown in Fig.

Fig. 14 An MO imaging setup with an MO imaging
plate with a backlight.

らつきが見えるという点が挙げられる．このざらつきは AFM 測定で見られたグレイン構造あるいはガーネット薄膜自体の磁区構

Speaker

CameraSpeaker
Taptic device

1 cm

(a)
Fig. 13 MO images of video tapes magneticallyrecorded with wavelengths of (a) 4 µm and (b) 1.16 µm,
measured using a B2.5:YIG film on GGG(001) 34).

(b)

Fig. 15 (a) A Photograph of a smartphone and (b) its
MO image.

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INDEX

Fig. 17 Kerr spectrum of the Bi2:NIG film on a glass
substrate.

Fig. 16 (a) Digital photograph of Nd magnets; (b)MO image and (c) line profile corresponding to thewhite line.
4.3 MO color imaging
Spectral shapes of the Kerr rotation of the MO imaging
plates measured in reflection geometry are influenced bylight interference. Consequently, they differ from their
original spectral structures, resulting in MO images that
could become colorful. For example, the Kerr rotation
spectrum of an MO imaging plate using Bi2:NIG with athickness of 450 nm is shown in Fig. 17. Peaks areobserved at wavelengths of 530 and 595 nm,
corresponding to green and yellow, respectively. By
varying the chemical composition and/or the thickness ofthe garnet film, or the light source, researchers could
vary the color depending on the magnetic field as
previously reported42). Figure 18 shows three results
obtained experimentally using various light sources (i.e.,a white LED, green and yellow LEDs, and blue andyellow LEDs) plotted in the chromaticity diagram(CIE1931), in which each curve indicates that color
varies with the magnetic field strength. The magnetic
field was quantitatively evaluated with colors expressed
as RGB values. More details of the MO color imaging
technique will be found in Ref. 42.
Figure 19 shows a demonstration of quantitative MO
color imaging using a white LED and a commercial
ferrite magnet. The color scale indicates the strength ofthe magnetic field. In the MO image, south (S) and north(N) magnetic poles are indicated by yellow and blue,
respectively; the zero point is indicated by violet.
We expect the resolution of the displayed magnetic field
to be improved through the use of a color instead of a
monochrome display. Specifically, the resolution of themagnetic field in color can potentially be the cube of thatin monochrome. For example, in the case of an 8-bit
camera, magnetic field values could be expressed with

114

Fig. 18 Color table showing the location of dataexperimentally obtained using white LEDs, green
and yellow LEDs, and blue and yellow LEDs as lightsources.

Fig. 19 (a) Digital photograph of a spherical magnet
and (b) the corresponding MO color image

28×3 = 16,777,216 colors, compared with 28 = 256 inmonochrome.
5. Summary

An MO imaging technique using an MO imaging plate
has been developed as a powerful method that visualizes
the magnetic field distribution at a wide range oftemperatures from cryogenic to greater than 100ºC, at
length scales from micrometers to several tens ofcentimeters, and at frequencies from DC to GHz. The
spatial and time resolutions have been improved throughthe use of Bi-substituted iron garnet films as MO imaging

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INDEX

plates. In this paper, Bi-substituted iron garnet films
fabricated by the MOD method were reviewed, where
highly Bi-substituted iron garnet films on glass
substrates as well as on GGG substrates were discussed.
In addition, an MO imaging plate with an EL backlightfor wide-area MO imaging and for MO color imaging withhigh magnetic resolution were reviewed. Further
advances in MO materials and measurement techniques
are expected to broaden the application range of MO
imaging.
Acknowledgements
I would like thank K. Sato for supporting my research on
magneto-optics and magnetic materials, and T. H.
Johansen for the opportunity to study MO imaging at
Oslo University. I am also grateful to numerous
coworkers—T. Kato, S. Iwata, T. Nishi, N. Adachi, H.
Asada, T. Nomura, M. Kawahara, K. Ishiyama, S. Hashi,
M. Veis, M. Sasaki, and M. Nishikawa—for support. I
also thank the many students of the Sato and Ishibashi
laboratory who contributed to studies on Bi-substituted
iron garnet films and MO imaging.
This work was supported by JST A-STEP Grant
Number (08003108), JSPS KAKENHI Grant Number(JP14655151, JP16360003, JP18H03776), the National
InstituteofInformation andCommunications
Technology (NICT) in Japan (program period: 2009–2016), the Nanotechnology Platform Program (Moleculeand Material Synthesis) of the Ministry of Education,
Culture, Sports, Science and Technology (MEXT), Japan,
and a joint research program with the Institute ofMaterials and Systems for Sustainability, Nagoya
University.
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Received June 2, 2020; Revised July 9, 2020; Accepted July 29,
2020

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J. Magn. Soc. Jpn., 44, 117-121 (2020)<Paper>MnGa Film with (001) Texture Fabricated on Thermally Oxidized
Si Substrate Using CoGa Buffer Layer
Y. Miwa*, D. Oshima**, T. Kato*, and S. Iwata***Department of Electronics, Nagoya University, Furo-cho, Chikusa-ku, Nagoya, 464-8603, Japan**

Institute of Materials and Systems for Sustainability, Furo-cho, Chikusa-ku, Nagoya, 464-8603, Japan(001)-oriented MnGa films were grown on thermally oxidized Si substrates using CoGa buffer layers by themagnetron sputtering method. The deposition and annealing temperatures of the CoGa buffer layer were optimized to
improve (001)-orientation of the MnGa layer. As a result, a saturation magnetization of 290 emu/cc was obtained
when CoGa was sputter-deposited at 400 ˚C and post-annealed at 700 ˚C. This saturation magnetization was almost
equal to that of MnGa grown on an MgO(001) substrate. The (001)-oriented MnGa film grown on the Si substrate had
surface roughness compared to that grown on MgO(001) due to the high temperature annealing of the CoGa layer.
Improvement of surface flatness is necessary for applications to magnetic memory and storage such as magnetic
tunnel junctions and bit patterned media.
K e y w o r d s : MnGa, CoGa, magnetic recording, recording media, bit patterned media,
1. Introduction
The areal density of hard disk drive (HDD) is
increasing every year, and further increase is desired.
Energy-assisted magnetic recordings, such as heat- or
microwave-assisted recordings, are under development
to boost the areal density of HDD. These technologies
will be followed by heated dot magnetic recording(HDMR), which combines heat- (or energy) assisted
magnetic recording and bit patterned media (BPM) to
extend the density more than 10 Tb/in2 1). In BPM, bitboundaries are defined by a nano-fabrication technique,and low transition noise is obtained without sacrificing
thermal stability compared to a conventional granular
media2). In general, ion milling followed by trench fillingwith nonmagnetic materials between bits is used todefine the bit boundary in BPM. However, there areseveral problems with this fabrication method of BPM,
including rough surface of the media, large distributionof the switching field, and high fabrication cost. An
alternative way to fabricate BPM is an ion irradiation
technique, in which accelerated ions are irradiated tomodify the magnetic properties locally without etching
the media physically3), 4). Ion irradiation is a simple andcost-effective fabrication method, and no flattening
process, such as trench filling, is required after the bitboundary has been defined.
In the early stages of the study of the ion irradiation
method, Co/Pt3)-6) and Co/Pd7)-8) multilayer films withperpendicular magnetic anisotropies were used as media.
The ion irradiation caused interdiffusion at the Co/Pt or
Co/Pd interface, which lowers the perpendicularanisotropy and makes them in-plane magnetized films
due to the demagnetizing energy. Thus, when these
multilayers are used as ion irradiation BPM,
perpendicularly magnetized bits are separated by thein-plane magnetized region, which causes exchange
coupling between the bits, and limits the ultimate

density of BPM. In order to avoid the exchange coupling,
we focused on perpendicularly magnetized Cr- andMn-based ordered alloy films, because their magnetism
was varied from ferromagnetic to paramagnetic by ion
irradiation through the change in their crystal structurefrom the ordered to the disordered phase9)-12). In ourprevious work, we demonstrated dot-patterned MnGa
films with a flat surface and low bit edge damage by 30
keV Kr+ ion irradiation12)-13). However, these results
were obtained using MnGa grown on MgO(001) single
crystal substrates, which are costly for practicalapplication. Thus, we have studied the growth of(001)-textured L10-MnGa films on thermally oxidized Si
substrates14).
For the mass production of ion beam patterned BPM,
resist masks are expected to be imprinted on themagnetic layer, and then accelerated ions are uniformly
irradiated through these masks15). Thin (low-height)
resist masks are generally required to reduce the pitch
size of the patterning. However, when a thin resist mask
is used, the energy of incident ions must be lowered tostop the incident ions by the resist masks / to protect the
MnGa bit by the irradiation. On the other hand, thelow-energy of ion irradiation requires a reduction in the
thickness of MnGa films, because the penetration depthof incident ions depends on the ion energy.
The thickness of MnGa should be less than thepenetration depth to kill the ferromagnetism of the“space” regions. Therefore, thin and uniform MnGa withlarge perpendicular anisotropy should be prepared on
low-cost substrates for practical application. However,reducing the MnGa thickness is known to degrade themagnetic properties, such as squareness and increased
coercivity16). Recently, Suzuki et al. reported that a
CoGa buffer layer was effective for obtaining ultrathin
MnGa films with a magnetic squareness close to unity17).
In this study, we report the fabrication of ultrathin
MnGa films on thermally oxidized Si substrates using

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the CoGa buffer layer, and discuss the optimized growth
conditions.
2. Experimental methods
The sample stack was Cr (2 nm) / MnGa (5 nm) / CoGa
(0 or 30 nm) / Cr (20 nm) / MgO (20 nm) / CrB (5 nm) /NiTa (25 nm) / SiO2 / Si substrate. An Ar+ ion beam wasused to clean the surface of the substrate. Most layers
were deposited by RF magnetron sputtering, and only
the MgO layer was deposited by e-beam evaporation.
NiTa and CrB were inserted to obtain the (001)-oriented
MgO buffer layer17). The sample was transferred
between sputtering and e-beam evaporation chambers
without breaking the vacuum to avoid contamination
during the deposition of the Cr/MgO/CrB/NiTa buffer
layer. The buffer layer was deposited at room
temperature, and the sample was then annealed at 800
˚C for 60 min. The CoGa layer was deposited at TsCoGa =200–500 ˚C, and then annealed at TaCoGa = 500–800 ˚Cfor 30 min. The MnGa layer was deposited at 200 ˚C, andthen annealed at 400 ˚C for 60 min. The Cr layer was
sputter-deposited as a protective layer at a temperature
lower than 150 ˚C. An epitaxial MnGa film on the
MgO(001) single crystal substrate was also prepared as

F i g . 1 (a) M-H curves, (b) XRD profile, (c) AFM,
and (d) MFM images of ultrathin MnGa film on MgOsubstrate using CoGa buffer layer. Full scales (FS) of
the images are (b) 20 nm and (c) 10 Hz.

118

a reference. The MgO substrate was heated at 600 ˚C for
10 min before deposition. The structure was Cr (2 nm) /MnGa (5 nm) / CoGa (30 nm) / Cr (20 nm) / MgO(001)substrate. The growth temperatures were 400 ˚C, 400 ˚C,and 300 ˚C for the Cr, CoGa, and MnGa layers,respectively. The post-annealing was performed at 600
˚C for 60 min, 400 ˚C for 30 min, and 400 ˚C for 60 min
after the deposition of the Cr, CoGa, and MnGa layers,respectively. Finally, the protective Cr layer was
deposited on the sample.
Magnetization curves were measured by analternating gradient field magnetometer. Crystal
structures were analyzed by X-ray diffraction (XRD)
using Cu Kα radiation. Surface morphologies andmagnetic domain structures were characterized byatomic force microscopy (AFM) and magnetic force
microscopy (MFM) with a frequency mode, respectively.
3. Results and discussions
3.1 Ultrathin MnGa film grown on MgOsubstrate using CoGa buffer layer
Figure 1(a) shows the out-of-plane and in-plane M-Hcurves of the ultrathin MnGa film grown on the MgOsubstrate using a CoGa buffer layer. The out-of-plane
curve showed a hysteresis with a squareness of unity; incontrast, the in-plane curve was almost linear with little
hysteresis, showing a strong perpendicular anisotropy.
The saturation magnetization Ms was ~ 300 emu/cc,
which was smaller than thick MnGa films12). Figure 1(b)shows the XRD profile of the ultrathin MnGa film. CoGa
001, MnGa 002, Cr 002 and CoGa 002 peaks wereobserved, implying all the layers were epitaxially grownwith (001)-orientation. The MnGa film had theL10-ordered phase from the fact that the ferromagnetism
was observed as shown in Fig. 1 (a), however, MnGa 001
peak was not visible due to the thin MnGa layer. The
peak position of MnGa 002 shifted to high angle sidecompared to the peak position of the bulk MnGa,
implying lattice constant shrinks in the film normal andexpands in the film. In addition, the peak shift for theultrathin MnGa was larger than that for the thick MnGa

F i g . 2 Saturation magnetization Ms and averageroughness Ra of ultrathin MnGa films on Si
substrates using CoGa buffer layers as a function of
growth temperature TsCoGa of CoGa layer.

Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

INDEXF i g . 4 Saturation magnetization Ms and averageroughness Ra of ultrathin MnGa films on Si substrates
using CoGa buffer layers as a function of annealing
temperature TaCoGa for the CoGa layer.

F i g . 3 (a) XRD patterns of ultrathin MnGa films
on Si substrates using CoGa buffer layers. The
growth temperature of CoGa TsCoGa was varied from
200–500 ˚C. (b) Lattice constant c of CoGa as afunction of TsCoGa. The dashed line shows the value ofbulk CoGa.
with Ms = 450 emu/cc12) (not shown here). Accordingly,
possible reasons for the small Ms of the thin MnGa films
are low degree of L10 ordering and lattice distortion.
Figures 1(c) and (d) show AFM and MFM images of theultrathin MnGa film, respectively, grown on the MgOsubstrate using the CoGa buffer layer. Averageroughness Ra estimated from the AFM image was 0.2 nm.
In the MFM image, bright and dark regions wereobserved, which reflects the maze-like domain structurecommonly seen in perpendicularly magnetized film.
These results show that an ultrathin MnGa film withperpendicular anisotropy and flat surface was grown on
the MgO substrate similarly to the results in ourprevious study14).
3.2
Ultrathin
MnGa(001)-oriented
polycrystalline films grown on Si substrates
using CoGa buffer layer
Details regarding the growth and annealing
conditions of the seed layers underneath the CoGa layer
are provided in our previous study13). In this paper,
growth conditions for the CoGa layer were investigatedto obtain high-quality ultrathin MnGa films. Figure 2
shows TsCoGa dependence of the Ms and Ra of ultrathin

F i g . 5 XRD patterns of ultrathin MnGa films on Si
substrates using CoGa buffer layers. The annealing
temperature of CoGa TaCoGa was varied from 500–800
˚C. The dashed line shows the peak positions of thebulk MnGa.
MnGa films on Si substrates, where the CoGa buffer
layer was post-annealed at TaCoGa = 500 ˚C. Ms wasestimated from the hysteresis loop along the film normal
direction, because the MnGa films were hard to saturate
along the in-plane direction. When TsCoGa = 300 ˚C and400 ˚C, the MnGa films showed perpendicularanisotropy. As shown in Fig. 2, large Ms and small Ra
were obtained at TsCoGa = 400 ˚C. The Ms shown in Fig. 2
is considered to be sensitive to chemical ordering, lattice
distortion, and (001)-orientation of the MnGa film.
However, X-ray diffraction does not give us a clear
answer for the variation of Ms, because MnGa 001 and
002 peaks were not visible, as shown in Fig. 3(a), due tothe thin MnGa layer. Figure 3(a) shows the XRD
patterns of ultrathin MnGa films on Si substrates using
CoGa buffer layers, where the CoGa layer was deposited
at TsCoGa = 200–500 ˚C and post-annealed at TaCoGa = 500
˚C. CoGa 001 and 002, Cr 002, and MgO 002 and 004
peaks were observed for all the samples. This meansthat the CoGa, Cr, and MgO layers were all grown with(001)-orientation. Figure 3(b) shows the TsCoGa

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INDEXF i g . 6 (a), (c) M-H curves and (b), (d) AFM images
with a full-scale (FS) 20 nm of ultrathin MnGa films
on a Si substrate. The films were grown on (a), (c) a
CoGa and (b), (d) Cr layer. The CoGa layer was
sputter-deposited at TsCoGa = 400 ˚C and annealed at
TaCoGa = 700 ˚C.
dependence of the lattice constant c of the CoGa layer,
which was calculated from the peak position of CoGa 001,
since CoGa 002 and Cr 002 peaks were overlapped. The
lattice constant of the CoGa layer became close to that ofa bulk CoGa19) (a = 0.2877 nm) with increasing TsCoGa up
to 400 ˚C, as shown in Fig. 3(b). This result may be oneof the reasons for the increase of Ms with increasing
TsCoGa up to 400 ˚C. Further increase of TsCoGa caused
significant increase in the roughness of the CoGa surface,
as shown in Fig. 2, which might inhibit growth of the(001)-oriented L10-MnGa layer.
Hereafter, TsCoGa was fixed to be 400 ˚C, and thepost-annealing temperature of the CoGa layer TaCoGa
was varied from 500–800 ˚C. Figure 4 shows TaCoGa
dependence of Ms and Ra of ultrathin MnGa films on Si
substrates. Ms reached maximum value at TaCoGa = 700
˚C, while Ra increased with increasing TaCoGa. The
increase in Ra is considered to be due to the increase of
the surface roughness of the CoGa layer with increasing
TaCoGa, because CoGa films without MnGa layers
annealed at 700 ˚C exhibited large Ra (not shown here).
Figure 5 shows XRD patterns of ultrathin MnGa films
on Si substrates using CoGa buffer layers. The CoGa
layer was deposited at TsCoGa = 400 ˚C, and thepost-annealing temperature TaCoGa was varied from
500–800 ˚C. Similar to Fig. 3(a), the samples prepared at
up to TaCoGa = 700 ˚C show (001)-oriented growth of the
CoGa, Cr, and MgO layers. At TaCoGa = 800 ˚C, CoGa 001
and 002 and Cr 002 peaks disappeared, which may havebeen due to the alloying of CoGa and Cr. The variationin the crystal structure of the CoGa layer may haveresulted in increased CoGa surface roughness. For the
samples prepared at TaCoGa = 600 ˚C and 700 ˚C, the
MnGa 002 peak was visible around 2θ = 53˚, which is120higher than the angle estimated in a bulk MnGa (dashedline). This suggests that the MnGa lattice wascompressed along the [001] direction and expanded inplane to decrease the lattice mismatch between MnGa
and CoGa. The appearance of the MnGa 002 peak is
considered to reflect the formation of (001)-oriented
L10-MnGa, and to be related to the increase of Ms, asseen in Fig. 4.
As discussed above, the Ms of the MnGa layer
estimated from the hysteresis loop along the film normal
direction was increased with increasing TsCoGa up to acertain temperature, and then abruptly disappeared. In
the present film stack, the rather high annealing
temperature of TaCoGa = 700 ˚C was effective inpromoting the crystal growth of the MnGa layer. This
post-annealing temperature was much higher thanthose for the CoGa buffer layer on the MgO single
crystal substrate (see Fig. 1). Such a high annealing
temperature was also necessary in the MnGa film grownon the Si substrate with Cr buffer layer13).
Figure 6 shows (a) out-of-plane and in-plane M-Hcurves, and (b) an AFM image of the ultrathin MnGa
film grown on the Si substrate using the CoGa buffer
layer with TsCoGa = 400 ˚C and TaCoGa = 700 ˚C (optimum
condition in the present study). As a reference, M-H
curves and an AFM image of the MnGa film grown on a
Cr layer are shown in Figs. 6 (c) and (d), respectively. Msof the MnGa films on the CoGa and Cr layers were 290
and 280 emu/cc, respectively, which were almost thesame as that on the MgO substrate shown in Fig. 1(a).
The squareness of the out-of-plane M-H curve increased
and the coercivity decreased by inserting the CoGa layeron the Cr layer, which suggests the (001) orientation andcontinuity of the MnGa layer were improved comparedto our previous study. However, Ra of the film with the
CoGa layer was 3.9 nm, greater than that of the MnGa
film grown on the Cr layer. This may have been due tothe rough surface of the CoGa buffer layer, andimprovement of the surface roughness is necessary forthe application to HDD. Further study of thecomposition and annealing conditions of the CoGa layer,such as rapid thermal annealing, should be investigated.
4. Conclusion
Ultrathin MnGa films on thermally oxidized Si
substrates were fabricated by inserting a CoGa buffer
layer, and the growth conditions of the CoGa buffer layer
were investigated to obtain MnGa film with a large
saturation magnetization. In the MnGa film on the Si
substrate, saturation magnetization of 290 emu/cc was
obtained when the CoGa buffer layer was
sputter-deposited at 400 ˚C and post-annealed at 700 ˚C.
The saturation magnetization of the MnGa film on the Si
substrate was almost the same with that on the MgOsubstrate. The hysteresis loop of the MnGa film
exhibited almost unity squareness and low coercivity by
the insertion of the CoGa buffer layer, suggesting the

Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

INDEX

improvement of (001) orientation and continuity of the
MnGa layer. However, the post-annealing of the CoGa
layer at high temperature resulted in increased surface
roughness. Improvement of the surface roughness isnecessary for practical application to BPM.

5)

Acknowledgments
The authors thank Mr.
Kumazawa for his assistance with the experiments. This
work was supported in part by JSPS KAKENHI Grant
Number
17H03249,
19K15044,
and
by
the
TATEMATSU foundation. This work was also supported
in part by the Project of Creation of Life Innovation
Materials for Interdisciplinary and International
Researcher Development of the Ministry of Education,
Culture, Sports, Science and Technology (MEXT), Japan.
A part of this work was performed under the Research
Program of “Dynamic Alliance for Open Innovation
Bridging Human, Environment and Materials” in“Network Joint Research Center for Materials andDevices.”
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Received Mar. 09, 2020;
Accepted Jul. 22, 2020

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J. Magn. Soc. Jpn., 44, 122-128 (2020)<Paper>Media Designfor Three-Di mensio nal Heat-Ass isted Ma gne tic Recording
T. Kobayashi, Y. Nakatani* , and Y. Fujiwara

Graduate School of Engineering, Mie Univ., 1577 Kurimamachiya-cho, Tsu 514-8507, Japan*

Graduate School of Informatics and Engineering, Univ. of Electro-Communications, 1-5-1 Chofugaoka, Chofu 182-8585, JapanWe design media for three-dimensional heat-assisted magnetic recording at 2 Tbpsi (total density of 4 Tbpsi),taking account of heat transfer simulation, field calculation, and information stability in high Curie temperature(HC) and low Curie temperature (LC) layers. The medium surface temperature while writing in the HC layer, thethermal gradients while writing in the HC and LC layers, and the grain temperature difference between the HC andLC layers while writing in the LC layer are discussed by using a heat transfer simulation. The magnetic fieldstrength as a function of distance from the layer surface for a single layer is discussed by employing a fieldcalculation. The Curie temperatures of T LC and T HC , and the grain heights of hLC and hHC in the LC and HC
layers, respectively, are also discussed in terms of information stability during 10 years of archiving and writing in
the LC layer. As a result, a media structure is decided whereby the upper, namely, the surface layer is an LC layer
with T LC = 625 K and hLC = 4.5 nm, and the lower layer is an HC layer with T HC = 750 K and hHC = 6.0 nm with!
!
!
!
an anisotropy constant ratio of 0.8.

!

K e y w o r d s : 3D HAMR, heat transfer simulation, field calculation, information stability

!

!

!

2 . Ca lc ula tio n Me tho d

1 . I ntr o duc tio nHeat-assisted magnetic recording (HAMR) is a
2 .1 He a t tr a n sfe r simula tio npromising candidate as a next generation magnetic
We assumed the medium to be granular. The
recording method capable of operating beyond thetemperature profile for 3D HAMR depends on thetrilemma limit1). HAMR is a recording technique wherematerials between the grains and between the HC andLC layers. Suitable intergrain and interlayer materials
the medium is heated to reduce coercivity during theare currently unknown. However, it is important towriting period.
obtain the characteristics of the temperature profile inThree-dimensional magnetic recording is useful for
3D HAMR. Therefore, we analytically estimated theincreasing the recording capacity. Three-dimensionaltemperature profile for the cross-track direction y
HAMR (3D HAMR) has been proposed2) where themedium consists of a high Curie temperature T HC
using data published in previous papers 4) and 5)
employing Eq. (1)(HC) layer and a low Curie temperature T LC (LC)layer with no exchange coupling between the HC andT 2 (y) " T amb Pw2!
LC layers. With 3D HAMR, once data have been
=
,
(1)
!
T1 (y) " T amb Pw1
written in the HC layer, other data can be written in!
the LC layer by employing lower temperature
heating.
where T i (y) is the temperature for a heating power
We have previously discussed the stability of the
Pwi . The ambient temperature T amb is the maximum
information in the HC layer while writing is under way
working
temperature of the hard drive, and is assumed!
in the LC layer, and the information stability in the HC
to be 330 K. The temperature increase above theand LC layers for 10 years of archiving3). This result ! ambient temperature T (y) " Ti
amb is proportional towas obtained under the assumptions that !the!
Pwi .
temperatures in the HC and LC layers are the same,
Figure 1 shows a schematic illustration of theand that the thermal gradient for the cross-track
structure of a medium that consists of four layers,
direction is constant. However, the two layers have!
namely a recording
layer (RL) (FePt base, HC 4 nm +different temperatures and thermal gradients.
!
LC 4 nm = 8 nm), interlayer 1 (IL1) (MgO base, 5 nm),
Furthermore, the magnetic field strength from theinterlayer 2 (IL2) (Cr base, 10 nm), and a heat-sink
HC and LC layers is also important as regards readout.
layer (Cu base, 30 nm). IL1 is a layer for c-axis
In this paper, we design the 3D HAMR media at 2orientation control of RL, and IL2 is a seed layer for IL1.
Tbpsi (total density of 4 Tbpsi), while taking account of
There is no intergrain material and no interlayera heat transfer simulation, a field calculation, and thematerial between the HC and LC layers. The x , y ,information stability in the HC and LC layers.
and z axes are the down-track, the cross-track, and!
122

the film normal direction, respectively, where y = 0 at
the track center and z = 0 at the RL surface. We!
focused on the medium surface temperature!T surf while!

! No.5, 2020
Journal of the Magnetics Society of Japan Vol.44,!

INDEX

writing in the HC layer, the thermal gradients "T /"yfor the cross-track direction while writing in the HC
and LC layers, and the grain temperature difference"T HL at the track edges between the HC and LC layers!
while writing in the LC layer.
The thermal constants are summarized in Table 1.
Two thermal conductivities of 50 and 4 W/(mK) wereused for IL1.

!

F i g . 2 Grain arrangement for field calculation. (a) Top
and (b) side views.
F i g . 1 Film structure for heat transfer simulation.

T a b l e 1 Thermal
simulation.

constants

for

heat

transfer

!

!

!

!

2 .2 Fie l d c a lc ula tio n
The magnetic field strength from the upper and lower
layers, as shown in Fig. 2, was calculated using ananalytical equation. One bit has 3 grains for the
down-track direction and 3 grains for the cross-track
direction, namely, there are 9 grains/bit. The bit aspectratio is one. The grain size Dm and the intergrain
spacing " G are 5.0 nm and 1 nm, respectively, with nograin size distribution. The flying height hfly is the!
distance between the magnetic head reader and theupper layer surface.! The grain heights and the!magnetizations for the upper and lower layers are h ,!UhL , M sU , and M sL , respectively. !
!
The interlayer spacing " L for suppressing the!
exchange coupling between the upper and lower layers! And "
strongly depends on the interlayer material.
L!
should be thinner due to the higher field strength at the! layer. On the assumption that!wereader from the lower
had selected an appropriate material, we chose a " L!
value of 1 nm.

2 .3 In fo r ma tio n s ta b ility
The bit error rate (bER) for estimating information
stability was calculated using the grain error
probability3), taking account of the shape anisotropy
M s H d / 2 using a self-demagnetizing field H d as
shown in Fig. 3. The temperature dependence of themagnetic properties was determined by performing amean field analysis3).
The conditions used when calculating!the information
stability in the HC and LC layers are summarized inTable 2.
We assumed that the grain size distribution was
log-normal with a mean grain size Dm of 5.0 nm and astandard deviation " D / Dm of 15 %.
The temperature in the center of the track whilewriting is higher than that at the track boundary by("T /"y) # (DT / 2) where !DT is the track width.
Therefore, !
the temperature of the writing head and thesurface lubricant while writing in the HC layer
decreases as DT , namely the bit aspect ratio DT / DB ,! peak temperature while writing indecreases5). And thethe LC layer also decreases as DT / DB decreases asshown in Fig. 4. Furthermore, increasing the grain! number per bit, namely, decreasing!
columnDT / DB , iseffective in reducing the bER caused by the Curie! 6). Therefore, we chose atemperature T c variation
DT / DB value of 1.
We have introduced an HAMR!design parameter,
namely, the medium anisotropy constant ratio
K bulk 4), in place of the medium anisotropy constant
Ku / !
K u since the K u value is a function of the medium T c .
K u / K bulk is the intrinsic ratio of the medium K u tobulk FePt K u regardless of T c where FePt is acandidate HAMR medium material thanks to its highand relatively low T c . Although a !low T c isKu !
!
advantageous for writing in HAMR, a higher
K u / K bulk!
!
is needed for a low T c medium for 10 years ofarchiving7), and a medium with a high K u / K bulk , e.g.
!
!
to manufacture regardlessof T c 8).
0.6 " 1.0, is difficult!
And then, a high K u / K bulk must be achieved in 3D! information stability in the HC layer
HAMR to realize

!

!
!

!
!

Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

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INDEX

while writing in the LC layer, and to realize
information stability in the LC layer during 10 years ofarchiving3). The K u / K bulk value was revised to 0.8
from 0.85 or 0.86 in a previous paper3) since the value of0.85 appears to be too high for manufacturing a
medium.
! assumed that the T distribution was
We alsocnormal with a mean Curie temperature T cm and astandard deviation " Tc /T cm of 2 %. No intrinsic
distribution of K u was assumed. However, there was a!by " .
fluctuation in K u caused
Tc
The temperature profile for the !
cross-track direction! in the LC layer3) was also revised, andwhile writing! in this paper is shown in Fig. 4. The LC
that used!
writing temperature!of the grains at the track edges
was assumed to be T LC + 2" TLC , taking account of thestandard deviation " TLC of T LC . The LC writing
temperature of the grain in the center of the track was
assumed to be

!
#T !
!
(Dm + $ G ) ,T LC + 2" TLC +#y(2)taking account of "T /"y = 10 K/nm.

!

!

F i g . 4 Temperature profile for information stability in
HC layer while writing in LC layer.
3 . Ca lc ula tio n Re sult s
3 .1 He a t tr a n sfe r simula tio n
The writing temperature of the grains at the trackedges was assumed to be T cm + 2" Tc for the HC andLC layers where T cm and " Tc are the mean Curie
temperature and the standard deviation of T cm ,respectively, taking account of the Curie temperature!
variation.
!
Figure 5 shows !the temperature profiles for the! interlayer 1
cross-track direction where we used an(IL1) thermal conductivity of 50 W/(mK) and an HC(upper, namely, surface, T HC = 750 K) / LC (lower, T LC
= 550 K) layer structure. The solid lines indicate thetemperatures at the layer boundaries, and the dottedlines indicate those at the layer centers.

!

!

F i g . 3 Self-demagnetizing field H d for grains withmagnetization M s and volume D " D " h where D
and h are grain size and height, respectively.

!

!
!
!
T a b l e 2 Calculation conditions.

!

(a) "T H /"y = 25.2 K/nm!

(b) "T L /"y = 9.5 K/nm
F i g . 5 Temperature profile for the cross-track direction
for an IL1 thermal conductivity of 50 W/(mK) and an
HC (upper) / LC (lower) layer structure while writing inthe (a) HC!
and (b) LC layers.

124

Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

INDEX!

!

!
!

(a) When writing in the HC layer (HC writing), thetemperature at y = ± 6 nm and z = " 2 nm is
T HC + 2" THC = 780 K at which "T H /"y is 25.2 K/nm.
The T surf value is 908 K.
(b) When writing in the LC layer (LC writing), the!
!
temperature
at y = ± 6 nm and z = " 6 nm! is! which "T /"y is 9.5 K/nm.
T LC + 2" TLC = 572 K at
L
! The "T value is 40 K.
HL
On the other hand, Fig. 6 shows the temperature! the cross-track !
profiles for
direction where the LC!
! structure is used.
(upper) / HC (lower) layer!
(a) When HC writing, the temperature at y = ± 6 nm
and z = " 6 nm is T HC + 2" THC = 780 K at which!
"T H /"y is 17.6 K/nm. The T surf value is 1004 K.
(b) When LC writing, the temperature at y = ± 6 nm!
and z = " 2 nm is T LC + 2" TLC =! 572 K at which!
! K/nm. The "T value is "35 K.
"T L /"y is 13.5
HL
The T surf value !
of 1004 K for LC / HC in Fig. 6 (a) is! in Fig. 5 (a) sincehigher than that of 908 K for HC / LC!
!
HC is the lower layer.
!
! of 25.2 K/nm
The "T /"y value
for the HC / LC
! structure and the upper HC writing in Fig. 5 (a) ishigher than that of 13.5 K/nm for LC / HC and theupper LC writing in Fig. 6 (b) due to their Curie! temperatures.

Employing Eq. (1), we find that the thermal gradient"T H /"y for the HC layer at the track edges is
intrinsically larger than "T L /"y for the LC layer asshown in Eq. (3),T H (y) " T amb T H (±6 nm) " T amb T HC + 2# THC " T amb==,
T L (y) " T amb T L!(±6 nm) " T amb T LC + 2# TLC " T amb
T H (y) " T amb ="T HC + 2# THC " T amb(T L (y) " T amb ) ,T LC + 2# TLC " T amb#T H (y) T HC + 2$ THC % T amb #T L (y)=&.
#y
T LC + 2$ TLC % T amb#y(3)

And then, the "T H /"y of 25.2 K/nm for HC / LC
structure and the upper HC writing in Fig. 5 (a) ishigher than that of 17.6 K/nm for LC / HC and thelower HC writing in Fig. 6 (a) due to a heat flow in thein-plane !
direction at the depths of the layer.
Therefore, the thermal gradient has the following
characteristics:(1) The thermal gradient for the HC layer isintrinsically larger than that for the LC layer due totheir respective Curie temperatures.
(2) The thermal gradient for the upper layer isintrinsically larger than that for the lower layer due toa heat flow.
These two characteristics are very important in 3D
HAMR.
The "T H /"y of 25.2 K/nm for the HC / LC structureand HC writing in Fig. 5 (a) is relatively large due tothe combination of the HC layer and the upper layer.
And the "T L /"y of 9.5 K/nm for the HC / LC structure! and LC writing in Fig. 5 (b) is relatively small due tothe combination of the LC layer and the lower layer.
On the other hand, for the LC / HC structure, the!
difference between the "T /"y values of 17.6 K/nm for
HC writing and 13.5 K/nm for LC writing in Fig. 6 issmall due to the opposite combinations.
If "T /"y is small, the information in tracks!
(a) "T H /"y = 17.6 K/nm
adjacent to the writing track becomes unstable.
Furthermore, "T /"y and "T /"x are almost the samefor the cross-track and down-track directions,! respectively4). If "T /"x is too large, the writing!
temperature decreases before writing, and write-error!
!
increases. If "T /"x is too small, it is difficult to reduce
the writing temperature, and the magnetization is!
reversed by the reversed writing field when writing thenext bit. Then, erasure-after-write increases. Therefore,!
the small difference in the "T /"x values for HC and
LC writing is preferable in terms of the writing
property8).
When the "T HL value is negative, the temperature!
in the HC layer while writing in the LC layer is lower(b) "T L /"y = 13.5 K/nm
than that in the LC layer, and this is advantageous in
F i g . 6 Temperature profile for the cross-track direction
relation to the information stability in the HC layer.
for an IL1 thermal conductivity of 50 W/(mK) and an
And! the "T HL value of "35 K for the LC / HC
LC / HC layer structure while writing in the (a) HC and!
structure in Fig. 6 (b) is negative since HC is the lower(b) LC layers.
layer.

!

!

Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

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INDEX

layers for 4 W/(mK) in Fig. 7 is lower than that for 50
W/(mK) in Fig. 6 since the heat flow perpendicular to
the film plane was reduced by the adiabatic effect of
IL1. Therefore, the T surf values for 4 W/(mK) in Table 4
are lower than the corresponding values for 50 W/(mK)in Table 3. Although the T surf values can be reduced,
T a b l e 3 Results of heat transfer simulation (thermal
the "T /"y values become small.
!
!
! be noted that the temperature of the LC
conductivity of interlayer 1 (IL1): 50 W/(mK)).
It should
layer ( z = " 2 nm) in the center of adjacent tracks! writing in the HC layer is 635 K as
( y = ± 18 nm) while! shown in Fig. 7 (a), and HC writing will erase data in
the adjacent tracks. Therefore, it is necessary to devise!
a writing method, e.g. combination with shingled!
magnetic recording.
In short, an LC (upper, namely, surface) / HC (lower)
layer structure is preferable as regards "T /"y and"T HL . An adiabatic medium structure with a low
thermal conductivity of IL1 is also preferable as
regards T surf .
The results obtained for the IL1 thermal conductivityof 50 W/(mK) are summarized in Table 3. Although the
T surf value is high, an LC (upper) / HC (lower) layer
structure is preferable as regards "T /"y and "T HL .

!

!

!
T a b l e 4 Results of heat transfer simulation (thermal!conductivity of interlayer 1 (IL1): 4 W/(mK)).

(a) "T H /"y = 13.0 K/nm
3 .2 Fie l d c a lc ula tio n
Figure 8 shows the peak z component H zpeak of themagnetic field as a function of the distance from thelayer surface for a single layer where M s is 1000
emu/cm3. For the upper layer,!

!

Distance = hfly ,and for the lower layer,!
(b) "T L /"y = 9.4 K/nm
F i g . 7 Temperature profile for the cross-track direction!
for an IL1 thermal conductivity of 4 W/(mK) and an LC/ HC layer structure while writing in the (a) HC and (b)LC layers. !
The results for the IL1 thermal conductivity of 4
W/(mK) are shown in Fig. 7 and summarized in Table 4.
The temperature difference between the HC and LC

!

(4)

!

Distance = hfly + hU + " L .

(5)

The dotted lines indicate fitting by an exponentialfunction.
The H zpeak value from the medium decreases
exponentially with a distance of more than 3 nm.
Therefore, a low hU (thin upper layer) is preferable interms of ensuring a high H zpeak from the lower layer.
!On the other hand, increasing the grain height (layer
thickness) is an effective way of increasing H zpeak asshown in!Fig. 9.

!

!
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Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

INDEX

The temperature dependence of M s was roughly
estimatedundertheassumptionofM s (T c = 770 K, T = 300 K) = 1000 emu/cm3 as shownin Fig. 10. M s (T c , T ) increases as T c increases. The! lower layer since it has a
HC layer is preferable as thelarger M s .
!
In short, a thin LC (upper) / HC (lower) layer!
!
structure is preferable as regards H zpeak .

!
!

!

bit error rate (bER) as a function of the grain height hunder the conditions of a T cm of 550 K, no writing field
H w , a storage temperature T sto of 350 K, and astorage time t sto of 10 years. It can be seen that the!
grain height of 6.5 nm is necessary for bER = 1E " 3 .
! value for 10 years of archiving was
The required T cm
! of h with a bER of 1E " 3 asobtained as a function!
shown in Fig. 12.
!
In short, the hLC value can be reduced by increasing
T LC !since K u increases as T c increases under a!
!
constant K u / K bulk 8).

!
!

!

!

!

F i g . 8 Peak field strength H zpeak as a function of
distance from the layer surface for a single layer withgrain height h .

!

!

F i g . 1 1 Bit error rate necessary for 10 years ofarchiving as a function of grain height (layer thickness)h.

!

F i g . 9 Peak field strength H zpeak as a function ofgrain height h for various distances.
F i g . 1 2 Mean Curie temperature T cm necessary for 10
years of archiving as a function of grain height (layer
thickness) h .

!

!

F i g . 1 0 Temperature dependence of magnetization
M s for various Curie temperatures T c .
3 .3 In fo r ma tio n s ta b ility fo r 1 0 ye a r s o f!
archiving!
Next, we discuss a way of reducing the LC grainheight (layer thickness) hLC .
The hLC value is limited by the information stability
needed for 10 years of archiving. Figure 11 shows the!

!

3 . 4 I n f o r m a t i o n s t a b!
ili ty in HC la ye r while
wr iting in LC la ye r!
The T LC value is limited by the information stability
in the HC layer while writing in the LC layer. The T HC
value required for the information stability was
obtained as a function of T LC for various HC grain! heights (layer thicknesses) h for a writing field HHCwof "10 kOe, a time t of 1 nm, and a bER!of 1E " 4 asshown in Fig. 13. The dotted line indicates the limit for! taking account of the result in
10 years of archiving,!
!
Fig. 12.
!
!
In short, the T LC value can be raised by increasing
T HC and/or hHC . However, increasing hHC to morethan 6 nm does not seem to be effective for increasing
T LC as shown in Fig. 14.

!

!

!

!
!

!
Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

127

INDEXF i g . 1 3 HC Curie temperature T HC required for
information stability in the HC layer as a function of
the LC Curie temperature T LC for various HC grainheights (layer thicknesses) hHC .

!

!
!

F i g . 1 4 T LC required for information stability in the
HC layer as a function of hHC at T HC = 750 K.
3.5 Layer de sign

! Finally, we decided the layer structure for the 3D! shown! in Fig. 15 using the above
HAMR media as!
!

results.
We assumed the upper layer to be the LC layer, andthe lower layer to be the HC layer with T HC = 750 K
and hHC = 6.0 nm, taking account of the results in 3 . 1 ,3 . 2 , and 3 . 4 . Then the LC layer has T LC = 625 K and
hLC = 4.5 nm as a result, taking account of "T HL in
3 . 1 and h in 3 . 3 , respectively. !
Figure 16 shows the z component H z of the! and HC layers for themagnetic field from the LC! are 1098 anddown-track direction. The H zpeak values!
570 Oe from the LC and HC layers, respectively, where!
!
hfly is 4.0 nm.

!
!

F i g . 1 5 Layer structure for 3D HAMR.

F i g . 1 6 Field strength H z from LC and HC layers forthe down-track direction.
4. Conclusion s
We designed! 3D HAMR media at 2 Tbpsi (total

density of 4 Tbpsi), taking account of a heat transfer
simulation, field calculation, and information stability
in high Curie temperature (HC) and low Curie
temperature (LC) layers.
(1) An LC (upper, namely, surface) / HC (lower) layer
structure is preferable as regards the thermal gradients
while writing in the HC and LC layers and information
stability in the HC layer while writing in the LC layer.
An adiabatic medium structure is also preferable as
regards the surface temperature.
(2) A thin upper layer and a lower HC layer are
preferable in terms of ensuring a high magnetic fieldfrom the lower layer.
(3) The LC grain height (layer thickness) hLC can be
reduced by increasing the LC Curie temperature T LC .
(4) The T LC value can be raised by increasing the HC
Curie temperature T HC and/or the HC grain height!
(layer thickness) hHC .
We decided the layer structure for !
the 3D HAMR
!media using the above results. The upper layer is an! T = 625 K and h = 4.5 nm, and the
LC layer withLC
LC
! is an HC layer with T = 750 K and hlower layer
HCHC= 6.0 nm with an anisotropy constant ratio of 0.8.

! acknowledge the support ofA c k n!
o w l e d g e m e n t We!
! Japan.
the Advanced Storage ResearchConsortium,
Re fe r e nc e s1) S. H. Charap, P. -L. Lu, and Y. He: IEEE Trans. Magn., 3 3 ,
978 (1997).
2) Western Digital: private communication.
3) T. Kobayashi, Y. Nakatani, and Y. Fujiwara: J. Magn. Soc.
Jpn., 4 4 , 34 (2020).
4) T. Kobayashi, Y. Isowaki, and Y. Fujiwara: J. Magn. Soc.
Jpn., 3 9 , 8 (2015).
5) T. Kobayashi, Y. Isowaki, and Y. Fujiwara: J. Magn. Soc.
Jpn., 3 9 , 139 (2015).
6) T. Kobayashi, Y. Nakatani, and Y. Fujiwara: J. Magn. Soc.
Jpn., 4 3 , 70 (2019).
7) T. Kobayashi, Y. Nakatani, and Y. Fujiwara: J. Magn. Soc.
Jpn., 4 3 , 114 (2019).
8) T. Kobayashi, Y. Nakatani, and Y. Fujiwara: J. Magn. Soc.
Jpn., 4 2 , 110 (2018).
R ece iv ed Jun. 30, 2 020; A cc ept ed Ju l. 27, 202 0

128

Journal of the Magnetics Society of Japan Vol.44, No.5, 2020

INDEX

Editorial Committee Members・Paper Committee Members
T. Ono and T. Kato (Chairperson), K. Koike, T. Taniyama and K. Kobayashi (Secretary)
H. GotoT. Hasegawa
S. Isogami
K. KamataH. Kikuchi
T. Kouda
S. Kokado
Y. Kota
T. Kubota
T. MakiS. Muroga
T. Nagahama
H. Nakayama
M. NaoeT. Narita
J. OzekiN. PhamT. Sasayama
T. SatoK. Sekiguchi
Y. Shiratsuchi
T. Takura
S. YamadaT. Yamamoto
K. Yamazaki

T. Kimura
T. Morita
D. Oyama
T. Shima

N. Adachi
N. Inaba
Y. Kamihara
M. Ohtake
T. Tanaka

S. Honda
K. KatoT. Nozaki
I. Tagawa
S. Yoshimura

K. Bessho
S. Inui
A. Kuwahata
T. SatoM. Takezawa

M. DoiK. ItoK. Masuda
S. Seino
M. Tsunoda

T. DoiY. Kanai
Y. Nakamura
T. Suetsuna
N. Wakiya

K. Hioki
H. KatoK. Nishijima
K. Tajima
S. Yabukami

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編 集 委 員・論 文 委 員
小 野 輝 男 （理事） 加 藤 剛 志 （理事）
磯上慎二
小瀬木淳一
小山大介
小田洋平
後藤博樹
笹山瑛由
中山英俊
長浜太郎
成田正敬
山崎慶太
山田晋也
山本崇史
安達信泰
桑波田晃弘
土井達也
薮上
信

伊藤啓太
佐藤
拓
土井正晶
吉村
哲

乾
成里
末綱倫浩
仲村泰明
脇谷尚樹

小 池 邦 博 （幹事） 谷 山 智 康 （幹事） 小 林 宏 一 郎 （幹事）
鎌田清孝
菊池弘昭
木村
崇
窪田崇秀
神田哲典
佐藤
岳
嶋
敏之
白土
優
関口康爾
田倉哲也
長谷川 崇
PHAM NAMHAI
槙
智仁
室賀
翔

稲葉信幸
清野智史
西島健一

大竹
充
田河育也
野崎友大

加藤宏朗
竹澤昌晃
日置恵子

加藤和夫
田島克文
別所和宏

金井
靖
田中哲郎
本多周太

古門聡士
直江正幸
森田
孝

神原陽一
角田匡清
増田啓介

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