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Vol.41 No.6
JournalJournal of the Magnetics Society of Japan
Electronic Journal URL: https://www.jstage.jst.go.jp/browse/msjmag

2017

Review Article

Development of Media Nanostructure for Perpendicular Magnetic Recording
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
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M. Futamoto and M. Ohtake…108

Hard and Soft Magnetic Materials
Synthesis and Magnetic Properties of (Zn2+Ti4+) Substituted W-type and Y-type Ferrites
K. Watanabe, K. Kakizaki, and K. Kamishima…127

Power Magnetics

Low Permeability Composite Magnetic Core Transformer with High Coupling Coeﬃcient andits Application to PFM Controlled Quasi-Resonant Mode Flyback-Type DC-DC Converter

広告◇東英工業株式会社

K. Sato, T. Sato, M. Sonehara, and H. Takeuchi…132◇計測エンジニアリングシステム株式会社◇ネオアーク株式会社◇TDK株式会社

41

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

6

2017

INDEX

Journal of the Magnetics Society of Japan
Vol. 41, No. 6
Electronic Journal URL: https://www.jstage.jst.go.jp/browse/msjmag

CONTENTS
Review Article
Development of Media Nanostructure for Perpendicular Magnetic Recording…………………………………………………………………………………… M. Futamoto and M. Ohtake

108

Hard and Soft Magnetic Materials
Synthesis and Magnetic Properties of (Zn2+Ti4+) Substituted W-type and Y-type Ferrites……………………………………………………………… K. Watanabe, K. Kakizaki, and K. Kamishima

127

Power Magnetics
Low Permeability Composite Magnetic Core Transformer with High Coupling Coefficient andits Application to PFM Controlled Quasi-Resonant Mode Flyback-Type DC-DC Converter…………………………………………………………… K. Sato, T. Sato, M. Sonehara, and H. Takeuchi

Board of Directors of The Magnetics Society of Japan
President:

K. Takanashi

Vice President:

K. Nakagawa, S. Nakamura

Director, General Affairs:

Y. Miyamoto, K. Niiduma

Director, Treasurer:

K. Aoshima, K. Ishiyama

Director, Planning:

Y. Saito, S. Nakagawa

Director, Editorial:

K. Kobayashi, T. Ono

Director, Public Relations:

H. Itoh, S. Greaves

Director, International Affairs:

Y. Takemura, M. Nakano

Auditor:Y. Suzuki, R. Nakatani

132

INDEX
J. Magn. Soc. Jpn., 41, 108-126 (2017)<Review>

Development of Media Nanostructure for Perpendicular Magnetic
Recording
Masaaki Futamoto and Mitsuru Ohtake

Faculty of Science and Engineering, Chuo University, 1-13-27 Kasuga, Bunkyo-ku, Tokyo 112-8551, JapanThe review covers the research and development of perpendicular media nanostructure focusing on the recording
layer. The recording layer material includes the Co-alloys with hcp structure, the magnetic multilayers, and the alloys
with ordered structures that have high magneto-crystalline anisotropies (K u ). The technologies of underlayer orinterlayer for aligning the easy magnetization axis of magnetic crystal grain perpendicular to the film plane are briefly
reviewed including the high K u magnetic materials for future recording media.
Observations of compositional and magnetization structures in sub-µm scale have played important roles in
improving the recording layer. Typical data on these characteristics are explained in relation to the media structure
improvements. Considering that high K u magnetic materials will be employed in future recording media, basic
experimental results related in controlling the easy magnetization axis and in keeping the surface flatness of L1 0 ordered magnetic materials are explained. Future possibilities for increasing the areal density beyond 1 Tb/in2 byimproving the magnetic material and the nanostructure of recording layer are also discussed.
Key words: perpendicular magnetic recording, thin film media, nanostructure, compositional structure, magnetization
structure, Co-alloy, ordered alloy, multilayer, high K u magnetic material, epitaxial growth
1. Introduction
The areal density of hard disk drive (HDD) is now
around 1 Tb/in2 which is 5 × 108 of the areal density
employed in the world-first HDD developed in 1956 (2Kb/in2, 100 bits/in, 20 tracks/in). Now all the HDDs are
based on the perpendicular magnetic recording (PMR)
technology that was invented by Prof. Iwasaki in 1976 1).
PMR involves nearly 30 years of research anddevelopment before the technology started to be used incommercial HDDs in 2005 by replacing the conventional
longitudinal magnetic recording (LMR) technology. Table
1 shows the brief history of research and development
until the shipment of PMR-HDD, listed by keywords froma viewpoint of perpendicular magnetic recording media
technology 2). Although the PMR research anddevelopment attracted much attention in the magnetic
recording community from late 1970’s to early 1980’s,
there was a dip period of 10 - 15 years, from middle 1980’s
to late 1990’s which is shown as the background shadow
in the table, with low activities in the academic and the
HDD-related communities. This was partially becausethat there were a couple of engineering issues 3-6) which
needed to be solved before applying PMR to thecommercial HDDs, and mainly due to that it wasconsidered at that time possible to increase the areal
density with the conventional LMR. The high cost for
production facilities in changing the recording scheme
could be an additional reason for the industries.
However, as the areal density increased beyond 10 Gb/in2,
the signal decay problem associated with thermal
instability with LMR 7-9) gradually narrowed theallowance in HDD product design. Triggered by a highdensity magnetic recording demonstration using PMR10)

108

and by the developments of PMR media technology whichsolved the engineering issues11-18), activities towardrealization of HDD products based on PMR emerged
drastically in the magnetic recording community fromthe year around 2000 and the first PMR-HDD productwas shipped in 2005 with an areal density of 133 Gb/in2
19, 20). The areal density of PMR-HDD increased sincethen about 10 times and the technologies for achieving 5- 10 Tb/in2 areal densities are investigated in the cuttingedge research laboratories21, 22).
In the PMR technologies, the recording media and thewriting head are different from those for conventional
LMR whereas other technologies like the reader head
and the mechanical positioning have much in common.
The present paper reviews the development of PMR
media technology focusing on the nanostructure of Coalloy recording layer. To further increase the areal
density well beyond 1 Tb/in2, magnetic materials with themagneto-crystalline anisotropy (K u ) higher than those of
Co-alloys need to be employed. Some of the basic key
points in tailoring the high K u magnetic thin films arealso briefly reviewed.
2. Technologies in controlling the structure of Co-alloy
recording layer
When a CoCr-alloy magnetic material with hcp
structure is deposited on a substrate, a texture growth
proceeds with the c-axis direction perpendicular to thesubstrate surface23), where the c-axis is the easymagnetization of Co-alloy crystal with an hcp structure.
However, crystal grains with different growthorientations coexist in the early stage of film growth near
the substrate as shown in the cross-sectionaltransmission electron microscope image of Fig.1 (a) 24).

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

Table 1 Research and development history of perpendicular magnetic recording media technologies 2)
Year
Areal density

5 Mb/in2

Recording layer

CoCr

1990

1995

2000

100 Mb/in2

1 Gb/in2

10 Gb/in 2

1985

1980

1975

10 Mb/in2

CoCrTaCoCrPt2005
100 Gb/in2

2010
500 Gb/in2

CoPtCr + Oxides

SUL noise reduction (JVC, 1996)
PMR invention by Prof. Iwasaki (Tohoku Univ., 1976)
Technologies
Improved underlayers (Hitachi, 1996-1997)
Media noise reduction (Hitachi, 1997-1998)
CoCr-alloy PMR media (Tohoku Univ., 1978)
High Hc, high Mr/Ms media (Hitachi, 1998)
・Thin film media
SUL noise observation (Hitachi, 1984)
CoPtCrO high Hc, high Mr/Ms media (Toshiba, 2000)
technology
c-axis control by underlayer (Hitachi, etc. 1985)
Improved SUL (Hitachi & many companies, 2000)
CGC media (IBM & Tohoku Univ., 2002)
Improved Ru-underlayer (Fujitsu, etc., 2005 - )PMR media, L10-FePt (AIT, 1997)
PMR media, SmCo5 (Shinshu Univ. Waseda Univ. etc., 2004)
Cross-sectional TEM (Hitachi, etc., from around 1984)
Segregation structure observation by TEM using chemically etched sample (NTT, 1984-1994)
Magnetization structure observation by E-holography (Hitachi, 1986-1987)
Magnetization structure observation by MFM (Hitachi, etc., from around 1992)
Compositional segregation observation by AP-FIM (Tohoku Univ., 1992-1994)
・Structure analysis
Compositional segregation observation by EELS and EDX TEM (Hitachi, from 1995)
technology
Basic magnetic properties determined using Co-alloy single-crystals, Ku, α, etc. (Hitachi, Tohoku Univ. from 1996)
Tribology-related problems, PMR-floppy-disk (Many companies, middle ‘80s)
Pulling out from Ba-ferrite floppy PMR media business (Toshiba, late ’80s)
EventsUS HDD companies shifted to LMR, Japanese companies followed (middle-late ‘80s)
Drastic growth of PMR R&D
Giving-up PMR-HDD business (Censtor, Fujitsu,1991-1993)(Japan, US, Europe, 1976 - 1985)
Negative report on PMR (IBM, 1998)
Thermal stability problem prediction (CMU, 1994)
52.5 Gb PMR demonstration (Hitachi, 2000)
Commercialization of PMR-HDDs (Toshiba, 2005)
Areal density: the areal density of commercialized HDD and therefore the values before the year 2005 are those of LMR-HDDs.
Recoding layer: the typical recording layer material studied for PMR media. There were many variations, for example, CoCrPtX (X=Ta, Nb, Si, B, etc.).
Background shadow: it shows the R&D activities of PMR in the world. There was a dip period of 10 - 15 years from middle 80s to late 90s.

(a)

(b)

CoCr-alloy

CoCr-alloy

Amorphous Ge

Polyimide substrate(No underlayer)

50 nm

50 nm

(d)

(c)

CoCrPt + SiO2

Ru

CoCr-alloy

SUL

Ti
Amorphous Ge

50 nm

10 nm

Fig. 1 TEM micrographs of PMR media 2). The numbers inparentheses are the years of observation or publication.
(a) CoCr-alloy thin film deposited on polyimide substrate withno underlayer, (1985). (b) CoCr-alloy thin film deposited on
amorphous Ge underlayer, (1985). (c) CoCr-alloy thin film
deposited on Ti/Ge composite underlayer, (1988). (d) CoCrPt+ SiO2 thin film deposited on Ru layer, (2008).

The presence of initial growth region is undesirable for
PMR media, since it deteriorates the perpendicularmagnetic anisotropy25, 26). To enhance the crystal
nucleation with the c-axis perpendicular to the substratesurface, underlayer materials were investigated aroundmiddle 1980’s, and the materials like hcp-Ta 27), hcp-Ti 28,29), amorphous-Ge 30) were found suitable for this purpose.
The underlayer has to offer a surface condition for Coalloy crystal to nucleate with the basal (0001) parallel tothe surface through hetero-epitaxial growth. Suchnucleation of Co-alloy crystal is also possible on aneutral-type surface that promotes free nucleation of Cocrystal, where the (0001) basal plane with the lowest
surface energy tends to be favored parallel to thenonmagnetic underlayer surface. Underlayer materials
were then extended to nonmagnetic hcp-CoCr 31, 32), hcpRu 30, 33), fcc-Au, fcc-Al 34), fcc-Pt 35), etc. and dualunderlayer structures; Ti/Ge 36), CoCr/TiCr 37, 38),
Co 3 O 4 /Pt 35), Pt/Ti 39), Pd/Ti 40), CoCrRu/TiCr 41), Ru/Ta 42,43), Ru-oxide/Ru 44), etc. for the preparation of PMR media.
Figures1 (a)-(d) show the variation in cross-sectional
structure of CoCr-alloy film by employing amorphous Ge,
hcp-Ti/Ge, and hcp-Ru underlayers 2). The c-axis of CoCralloy crystal is well controlled perpendicular to thesubstrate surface, and the crystal grain diameter
distribution is narrowed for the films with hcp-Ti/Ge andRu underlayers. The distributions of crystal grain
diameter for Ti and Ru crystals are relatively narrow
when the underlayer materials are deposited on neutraltype amorphous layers. The perpendicular magnetic
anisotropy was enhanced greatly by introducing such
underlayers.

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INDEX(a)(b)
CoCrPt (+SiO2)

CoCrTaInterface

Disordered region

Position (nm)

18

(c)

(d) 330 C

214 C

12

Ru (0001)

Ti (0001)
3 nm

3 nm

Fig. 2 Cross-sectional high-magnification TEM micrographs
showing the interface between recording layer and underlayer 2).
(a) CoCrTa film deposited on Ti (0001) underlayer, (1991) 47).
(b) CoCrPt-SiOx thin film deposited on Ru (0001) underlayer,(2008).
With increasing linear recording density, theperpendicular coercivity must be increased while therecording layer thickness needs to be reduced to enhance
the writing efficiency of a recording head. In the nearly
30 years since the PMR invention in 1976, the recording
layer material varied from a binary CoCr-alloy to ternaryor more complex alloys of CoCrTa, CoCrPt, and CoPtCr +X (X = Ta, Si, Nb, B, oxides). The shift was necessary
for increasing the K u value of recording layer, which was
directly related with increasing the perpendicularcoercivity. When the layer thickness was decreased for
Co-alloy recording layer materials to be less than 50 nm,
a notable decrease in perpendicular coercivity wasobserved and the decrease depended on the underlayermaterial 12, 45, 46). Careful examinations of the interface
between the Co-alloy recording layer and the underlayer
revealed that there was an atomically disordered regionof a few nanometers around the interface 38, 46, 47).
Figure 2(a) shows the TEM image of initial growth region
observed for a CoCr-alloy film deposited on a Ti
underlayer 2, 47). The formation of disordered region waspartially due to a large lattice misfit of about 15%
between the CoCr-alloy magnetic material and the Ti
underlayer, and also due to an inter-diffusion of elements
which was caused during a deposition process at anelevated substrate temperature of around 250 °C. An
introduction of nonmagnetic layer with hcp-crystal
structure such as CoCr 35 , CoCr 25 Ru 25 , or Ru, which
possessed a lattice constant close to that of the Co-alloy
recording layer material, was proved effective in forminga sharp elemental interface while keeping an epitaxial
relationship between the two layers 12, 46-50).
Compositional variations across the interface were
investigated by using a TEM equipped with an EDX, andit was confirmed that the distance of elements diffusionfrom the underlayer toward the magnetic layer was verysmall, less than the electron-beam diameter of 2 nm

110

6
0
0

20
10
Number of SFs

30 0

20
10
Number of SFs

30

Fig. 3 Distribution of stacking faults (SFs) in CoCrPt recording
layer (2002).
(a) Cross-sectional TEM image, (b) high-resolution image
showing an example of stacking fault, (c) SF distribution along a
CoCrPt layer grown at a substrate temperature of 214 °C, (d) SF
distribution along an another CoCrPt layer grown at 330 °C.
The structure is CoCrPt(18 nm)/NiTaZr/Glass substrate 52).
employed in the EDX analysis 49). An example of CoCrPt(+SiO 2 ) magnetic layer deposited on a Ru layer is shownin Fig. 2(b). Formation of atomically sharp interface wasproved effective in maintaining high perpendicularcoercivities of PMR media for reduced magnetic layer
thicknesses.
Another origin that deteriorates the perpendicularmagnetic anisotropy is stacking faults (SFs) in the hcpCo-alloy crystal. When a SF exists in an hcp-Co-alloy
crystal, the local atomic stacking sequence becomessimilar to that of fcc-structure. The magnetic anisotropyof fcc-Co with cubic symmetry is one magnitude lower
than that of hcp-Co 51). Perpendicular coercivity
decreases when SFs exist in the crystals of hcp-Co-alloy
recording layer. Distribution of SF was investigated for
the Co-alloy recording layers using a high-resolution
TEM and it was revealed that it depended on the process
condition like substrate temperature, deposition rates ofrecording layer, underlayer, and intermediate layermaterials. An example of such analysis on the
distributions of SF along CoCrPt crystal columns isshown in Fig. 3 52). The recent investigations on CoPtalloy materials have shown that the atomic attacking
sequence is also influenced by the average number ofvalence electrons of alloy material, and the SF density
decreases when the number approaches to a certain
value through adjustment of compositional alloy
elements 53). In PMR media productions, the media
structure and the process conditions are carefully
controlled to lower the SF density in the recording layer.
3. Compositional structure of recording layer
Compositional fluctuations of Co and Cr in vacuum

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX(a)

Plan-view TEM image

(c)

(b)

Cr image(a)(b)

Co image

Fig. 4 Cr and Co distributions in a CoCrTa perpendicular medium
deposited at 230 °C investigated by EELS-TEM (1996), (a) planview TEM image, (b) Cr distribution map, and (c) Co distribution
map 66, 67).
evaporated or sputter deposited CoCr alloy thin films
were studied initially from a viewpoint to explain themagnetic property differences between thin film and bulk
samples, and they were investigated by using NMR 54),
FMR 55), chemical etching 56-58), and TEM equipped withan analytical facility in middle to late 1980’s 59, 60). It
was in late 1980’s when a relation between the nonmagnetic Cr concentration in LMR-CoCrTa recording
layer and the medium noise property was investigatedand it was demonstrated that a Cr-rich CoCrTa thin film
medium showed a lower noise property 61-65). Magnetic
crystal grain isolation while controlling the diameter inthe recording layer was realized necessary for highdensity perpendicular magnetic recording to increase thecoercivity, the recording resolution and to decrease themedium noise. Magnetic decoupling and/or reduction ofinteraction between the neighboring magnetic crystal
grains were achieved by enhancing segregation ofnonmagnetic metallic elements (Cr, Ta, etc.) toward
crystal grain boundaries. The technology has been
improved, from around the year 2000, by employing atechnique of nonmagnetic material (SiO x etc.)
precipitation along the magnetic crystalline grainboundaries, which is based on a physical segregation of
two different phases, metallic magnetic and nonmagnetic
amorphous materials, during a film formation process.
Figure 4 shows the Cr and the Co distributions visualized
for a Co-17at%Cr-3at%Ta PMR medium, where
nonmagnetic Cr atoms are strongly segregated near thegrain boundaries and are depleted inside the crystal
grains 66, 67). The distribution of Co atom is showing an
opposite tendency. When the Cr concentration in a CoCralloy material exceeds 30 at. %, the material becomesnonmagnetic at room temperature.
Employment of a sputter deposition system withultra-high-vacuum background pressure 37, 68) and with ahigh purity Ar sputtering gas were shown effective inenhancing nonmagnetic elements segregation toward themagnetic crystalline grain boundaries 69). The residual
gasses in a sputter deposition chamber such as watervapor deteriorate the perpendicular magnetic properties
through selective reaction with active Cr atoms forming
oxides and/or hindering segregation within the magnetic
crystal grains 70). Strong Cr segregation around the

10 nm

2 nm

Fig. 5 Plan-view TEM images of CoCrPt-SiOx perpendicular
medium, (2008). (a) TEM image showing the magnetic crystal
grains isolated by thin oxide layer and (b) high magnification
TEM image 88).
magnetic crystal grains weakens or decouples the direct
magnetic exchange between the neighboring magnetic
crystal grains, thus contributing to a reduction of
medium noise, an increase of coercivity, and an increaseof recording resolution.
The high-density magnetic recording demonstration
using PMR in the year 2000 10) was performed by
employing CoCr-alloy perpendicular media withenhanced Cr segregation around the magnetic crystal
grain boundaries 48, 71). Although the low noise and thehigh recording resolution properties were achieved withthe media, time dependent decays of recorded signals
were observed particularly at low linear densities 72. 73).
This was because that the remanent magnetization M r
was lower than the saturation magnetization M s ,M r /M s <1, where a demagnetization field was working inan opposite way to the recorded bit direction and caused
a time-dependent thermal decay in the recorded
information. Although the necessity of media preparation
with squareness, M r /M s = 1, was realized for suppressing
such time-dependent thermal decay of signal, the M r /M svalues of CoCr-alloy media were less than 1 up to the late1990’s 71-73). Magnetic multilayers composed of Co andnoble metals (Co/Pt, Co/Pd) were, on the contrary, known
to be tuned easily to possess the value, M r /M s , to be 1,
and they were investigated as practical magnetic
recording media 74, 75-79). However, the multilayer
structure which needed complicated film deposition
processes was one of the drawbacks of media production.
A novel technology in increasing the M r /M s to be nearly
unity while achieving a magnetic grain isolation wasdeveloped from the year around 2000. The technology
introduced oxygen into the Co-alloy thin films, which had
a similarity with that of the magnetic tape technologythat employed low noise Co-CoO recording layers 80-82).
Addition of oxygen or oxide to a CoPt or a CoCrPt
magnetic material during the sputter deposition process

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INDEX

Oxide region(nonmagnetic)

Segregation
Co-alloy crystal(magnetic)
c-axis(Co-alloy)

Hetero -epitaxy
c-axis(Ru)

Ru underlayer

Fig. 6Structure model of CoCrPt-SiO x perpendicular
medium. Co-alloy magnetic crystals are epitaxially grown on caxis oriented Ru crystals. Thin oxides are separating magnetic
interaction between Co-alloy crystal grains126).

CoCr recording layer

Scale downin 22 years
350nm

CoPtCr-SiOx recording layer

18 nm
Ru

Ti/Ge underlayer
Hc=850 Oe, d=40nm, h=300nm
h/d=7.5

SUL

Hc=4600 Oe, d=9 nm, h=18 nm
h/d=2.0[Single-layer perpendicular medium, (1988)] [Double-layer perpendicular medium, (2010)]

Fig. 7 Comparison of perpendicular recording media. Recording
layer height (h) and grain diameter (d) have been reduced greatly
while perpendicular coercivity has been increased in the 22 years
from 1988 to 2010 2).

showed a drastic effect in enhancing the magnetic
crystalline grain isolation. The technology was first
applied to CoPt media in 1994 83) and then to CoPtCr
media in 2000 84). Perpendicular magnetic properties
including the coercivity and the squareness (M r /M s ) weregreatly improved by addition of oxides like SiO x , whichwere carried out in the early years of this century,
accelerated the realization of commercial PMR HDDs 85,
86).
A plan-view TEM image of CoCrPt-SiO x
perpendicular medium is shown in Fig. 5.
Thecompositions of CoCrPt-SiO x perpendicular media werestudied by employing a TEM equipped with an EDX
facility, where the local compositions of 1 - 2 nm indiameter area could be determined. With this type media,
the magnetic crystal grains of sub-10 nm in diameter are
separated by oxide-based grain boundaries of around 1
nm width and this type granular media have been usedas the recording media of commercial HDDs. The highmagnification TEM image clearly indicates that most of
the grains are single crystals and the grain boundary112

structure is amorphous. The grain boundary
compositions were investigated, and it was found that
not a small amount of metallic elements other than Si
were included in the grain boundaries 87-89).
The
metallic elements (Co, Cr, Pt, etc.) are considered to bedissolved in the matrix SiO x through forming oxides like
CoO and Cr 2 O 3 . The composition indicates that the grainboundary is non-magnetic and the thin oxide layer isdrastically reducing the magnetic interaction between
the neighboring magnetic crystal grains. A model of
PMR medium126), based on the structural and thecompositional investigations, is depicted in Fig. 6.
Various nonmagnetic materials such as C 90), SiO 2 91),
TiO 2 92), Ta 2 O 5 93), etc. have been investigated formagnetic crystal grain isolation of hcp-Co-alloy materials.
Although any nonmagnetic material that segregatesalong the grain boundary during a sputter deposition
process is possible for magnetic crystal grain isolation,practical conditions in a mass-production of recording
media including easy handling, homogeneity ofnanostructure, reproducibility in the fabrication, etc. areconsidered. As a result, SiO x -based oxide materials are
now widely employed in the commercial PMR media
fabrication.
4. Tuning the nanostructure of recording layer
Figure 7 compares the cross-sectional TEM
micrographs of Co-alloy PMR media developed in middle
1980’s and in 2010 showing the development of
nanostructure in over a quarter century. While achievinga drastic scale down in the length and the width of c-axis
oriented Co-alloy crystal grain, the medium coercivity
was increased from around 0.8 kOe to be larger than 4
kOe by controlling the alloy composition, the underlayermaterial, and the processing condition, all of which were
inevitable in increasing the areal density. The PMR
media used in commercial HDDs employ the double
magnetic layer structure consisting of a semi-hardmagnetic recording layer stacked on a soft magnetic
underlayer, which is the basic PMR structure proposedby Prof. Iwasaki 1). Until late 1990’s, the Co-alloy
recording layer was deposited directly on a thick softmagnetic underlayer of crystalline Fe-Ni 94, 95),
amorphous Co-Zr-Nb 11, 96), etc. However, a series ofresearch has shown that the medium noise can be
reduced by introducing a thin nonmagnetic layer
between the soft magnetic underlayer and the semi-hardrecording layer, where the nonmagnetic layer is now
defined as an intermediate layer since the soft magnetic
layer is called as “underlayer” in the double layer
structure 97). Such magnetic decoupling between thetwo magnetic layers was effective in preventing atransmission of short wave length noise caused in the soft
magnetic underlayer 98-100).
The introduction of nonmagnetic interlayer offeredanother positive effect in controlling the structure ofrecording layer independently from that of soft magnetic

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

underlayer (SUL), where a smaller thickness SUL wasdesirable in achieving an efficient magnetic writing witha single-pole writing head. As explained in the previous
section, the nanostructure of recording layer must be
controlled simultaneously in the crystal orientation, thegrain diameter, as well as the separation between
magnetic grains with a presence of thin nonmagnetic
layer. Dual layer structures of hcp-Ru/Ta 42, 43), hcpRu/Ni-W, or even more complicated stacked structures 101,102) were employed as the intermediate layer, where thefirst layer of Ta or Ni-W provided a surface condition that
promoted free nucleation of hcp-Ru crystal with the(0001) basal plane, energetically most stable plane,parallel to the surface. A sputter deposition consistingof low Ar pressure deposition followed by high Ar
pressure deposition for the Ru layer preparation was
employed to produce an intermediate layer with the caxis oriented dome-like Ru crystal grains which wereeffective in enhancing nonmagnetic SiO x segregation
around the Co-alloy magnetic crystal grains 101, 102). The
Ar pressure for sputtering is known to give an influenceon the topological nanostructure of deposited film 103).
Since a PMR medium is used in combination with amagnetic head in an HDD system, a careful adjustmentin the magnetic property of recording layer is carried out
through an optimization not only from the intermediate
layer but also for the Co-alloy composition, thenonmagnetic material for magnetic crystal separation,
the sputter deposition condition, etc. Magnetic property
control along the thickness of Co-alloy recording layer by
employing a compositional gradient technique or adeposition of layers with different properties was used fortuning the medium property suitable for high areal
density PMR 101, 104, 105).
5. Recorded magnetization structure
Magnetization structure observation gives valuable
information for the investigations on the interrelationship between medium structure and recording
characteristics of recording resolution and medium noise.
Bitter method is the easiest technique and it was appliedin middle 1980’s to investigate the effect of c-axis
orientation of CoCr-crystal grains on perpendicularmagnetization. An example is shown in Fig. 8. The
Bitter observation from the top and the bottom sides ofrecording medium showed that a stronger perpendicularmagnetization was realized for a medium with higher caxis orientation 106). More quantitative estimation of
perpendicular magnetization was later carried out in1986 - 1987 by using the electron holography
technique106-108). Figure 9 shows the electron holography
images of CoCr-alloy media with different degrees of caxis orientation observed for the media cross-sections.
The data showed that a perpendicular magnetization at
300 kFCI (bit length = 80 nm) recording was achieved
through the whole film thickness of 200 nm for the highly
c-axis oriented CoCr medium sample, indicating that the(a-1)

Front surface(a-2)

0.20 µm thickness

Rear surface

(b-2)

(b-1)

0.35 µm thickness

10 µm

Fig. 8 Optical microscope images of recorded magnetization
observed by Bitter method. CoCr perpendicular thin film media
with thicknesses of 0.20 and 0.35 µm were observed from thefront and the rear surfaces, (1987).
Bitter images of 20 kFCI recording on a 0.20 µm thick CoCr/Ge
perpendicular thin film observed from (a-1) top and (a-2) rearsurfaces.
Bitter images of 20 kFCI recording on a 0.35 µm thick CoCr/Ge
perpendicular thin film observed from (b-1) top and (b-2) rearsurfaces 106).

(b)

(a)

Front
Surface

CoCr
Ge
100 nmRear
SurfaceFig. 9 Electron holography images of magnetization recorded at
300 kFCI observed for cross-sectional CoCr perpendicular
media samples without and with Ge underlayer, (1987).
Magnetic leakage flux intensity is different between the front
and the rear surfaces for a CoCr film without underlayer (a).
Magnetic leakage flux intensity is similar at the front and therear surfaces, indicating a strong perpendicular magnetization isrealized for the film with Ge underlayer (b) 106-108).
medium structure, particularly the alignment of easymagnetization orientation, was giving a very strong
influence in achieving a high linear density magnetic
recording.
Magnetization structure observations were also
carried out by employing high-resolution magnetization
observation techniques such as Lorentz TEM 109), spin
SEM 110), and magnetic force microscopy (MFM) 37, 75, 111113). MFM is the technique developed as an extension of
scanning force microscopy invented in 1986 114) and thetechnique started to be employed in the magnetics
community from late 1980’s. Since then, MFM has been
most widely used to visualize the recorded information
because of its easy operation and no special requirementof pre-processing for the observation samples. Figures
10 (a) and (b) compare the magnetization structure of
PMR- and LMR-CoCrTa media with similar alloy

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

113

INDEX(a)

20 kFCI-PMR(b)

10 kFCI -LMR
10

1

Track pitch (nm)10

FCI/TPI=1

1000

100

2 µm

(c)

100
100
1000
10

500 nm
70 Gb/in2 – LMR

500 nm185 Gb/in2 – PMR

Limit of granular media?

10000

Super paramagnetic limit ?

Atom
100000
10000
100000

PMR
1000

Track density (kTPI)

LMR
100

10

1

500 nm528 Gb/in2 – PMR

500 nmFig. 11 Decreasing trend of bit size in commercial HDDs 2).

2 µm

Fig. 10 MFM images of magnetization structure recorded on Coalloy recording media, (a), (b): (1994, and (c) (2007) 2,113). (a)MFM image of Co-17at%Cr-5at%Ta perpendicular medium
recorded at 20 kFCI. (b) MFM image of Co16at%Cr-4at%Ta
longitudinal medium recorded at 10 kFCI. (c) MFM image ofcommercial PMR medium shipped in 2007 with an areal densityof 185 Gb/in2.
compositions observed in 1994 113). The LMR medium
was observed with a relatively uniform MFM contrast
within sharply recorded bit transitions, whereas the
PMR medium was with an irregular MFM contrast andwavy bit transitions which correspond to higher medium
noise and poorer recording resolution. The difference inmagnetization structure was correlated with thenonmagnetic elements (Cr, Ta) segregation around themagnetic crystal grain boundaries, which was later
confirmed to be depending on the crystal grain boundary
structure. When deposited under a similar condition,enhanced segregation was achieved more easily with an
LMR medium that included higher-angle crystal grainboundaries, whereas segregation of nonmagnetic
elements was poor for a PMR medium consisting of c-axis
oriented hcp-crystal grains which tended to form lowerangle crystal grain boundaries. The crystal lattice of caxis oriented hcp-grain in PMR medium matches withthat of neighboring crystal every 60 degrees of rotation
along the perpendicular direction, whereas that of LMR
medium with the c-axis lying in the film plane matches
every 180 degrees and therefore crystal grains of LMR
medium tend to form large angle crystal grainboundaries. Because the local stress in crystal grains of
LMR medium presumably larger than that of PMR
medium, Cr segregation was accelerated for the LMR
medium.
The amorphous phase segregation along magnetic
crystalline grain boundaries by employing SiO x oxide as
the nonmagnetic material sputter deposited together
with a CoPt-alloy magnetic material, which was based on

114

Bit length (nm)

2 µmLinear density (kFCI)

1000
FCI/TPI=10

10 Gb/in2 – LMR

a two-phase separation mechanism, showed a drastic
effect in forming very sharp compositional transitions
from magnetic to nonmagnetic regions. As a result,smaller-size recorded bits surrounded by sharp
magnetization transitions became possible as shown in
the MFM image of CoCrPt-SiO x perpendicular media
depicted in Fig. 10 (c). The signal to noise ratio was
improved and thus the continuous decrease in bit size
became possible while keeping the thermal stability ofrecorded information which had caused problems with
LMR media 7, 8). The decreasing trend of bit size incommercial HDDs is shown in Fig. 11 together withrelated MFM images 2). The bit shape is changing froman elongated rectangular to a square with increasing
areal density and by shifting the recording technology
from LMR to PMR.
With increasing areal density toward 1 Tb/in2, thebit length is decreasing to be shorter than 25 - 30 nm
where a spatial resolution of 10 nm or better is necessary
to clearly observe the recorded bit patterns. Since thespatial resolution of commercially available MFM tip was
limited at around 20 nm, the structure and the magnetic
material for MFM tip fabrication were systematicallyinvestigated to improve the spatial resolution and themagnetic switching field 115-119). High magnetic switching
field is required for an MFM tip in the magnetization
structure observation of high coercivity medium, where
the tip is exposed to a strong magnetic flux emanatingfrom the sample that may change the tip magnetization.
When the tip magnetization changes, the MFM contrast
varies during an observation process. MFM tips withhigh magnetic switching fields were prepared by coatinghigh K u magnetic materials on nonmagnetic base tips 120122). Through optimization of tip fabrication condition,
MFM tips with spatial resolutions of 6 - 7 nm could be
developed 123, 124). The switching field was increased to behigher than 2 kOe with maintaining high spatial
resolutions of 10 nm or better 117, 125).
Figure 12 shows the magnetization structure of 500 -Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

1000[ 10 6 erg/cm 3 ]500 kFCI (50.8 nm)

750 kFCI (33.9 nm)

SmCo 5PrCo5FePt
Sm 2 Fe17N3
YCo5
Nd 2Fe14B
Sm 2 Co17
MnAl MnBi
Co 3Pt
FePd

100

CoPt

Ku

10

Co

CoPtCr-alloy

Ku = 2πMs2
11000 kFCI (25.4 nm)

0

5

Ms2

10

15

20

[ 10 5 (erg/cm 3 ) 2 ]25

Fig. 14 Ku values of magnetic materials plotted as a functionof Ms2 ,127-129).

1100 kFCI (23.1 nm)
Fig. 12 Magnetization structure of 500 – 1100 kFCI (bit length:50.8 – 23.1 nm) recording on a CoCrPt-SiOx PMR medium
observed by using a high-resolution MFM tip, (2013) 118, 119).

(a)and TEM images) were not observed for a same area but
from different areas of a same medium sample, a detailed
information how the individual bit is recorded on
granular magnetic crystal grains can be revealed, which
is useful in finding the points for further improving the
medium structure to be more suitable for higher areal
density magnetic recording 126).
6. Nanostructure tailoring of magnetic materials

(b)

50 nm

50 nm
Fig. 13 Relationship between magnetization structure (1000kFCI recording) and magnetic crystal grain distribution, (2013).
The schematic MFM image, (a), is obtained by overlapping an
MFM image with a plan-view TEM image, (b), observed for asame PMR medium 118,119).
1100 kFCI (bit length: 51 - 23 nm) recorded on a CoCrPtSiO x PMR medium observed by employing a highresolution tip 119). When a plan-view medium structure
observed by TEM is overlapped with an MFM image, a
relationship between the recorded magnetization and themagnetic crystal grain structures can be demonstrated
as shown in Fig. 13 118).
Here, the magnetic crystal
grains along one track are colored for easy recognition of
the inter-relationship. Although the two pictures (MFM

6.1 High K u magnetic materials
The K u value of CoCrPt-alloy material used as therecording layer in the current HDDs is limited to bearound 107 erg/cm3 even the alloy composition and the
media production process conditions are modified.
Magnetic materials with higher K u values are required
for further increasing the recording density. The volume(V) of magnetic crystal grain included in a magnetic bitof recording layer needs to be continuously decreased
while keeping the thermal stability condition of recorded
information, K u V/k B T > 35-70 127, 128), where k B is
Boltzmann’s constant and T is absolute temperature.
There are candidates for such magnetic materials with
K u values greater than 107 erg/cm3; (1) magnetic
multilayers of Co and noble metals, Co/Pd, Co/Pt, (2)ordered alloys of rare-earth and 3d-transition metals,
SmCo 5 , PrCo 5 , GdCo 5 , Sm 2 Co 17 , etc., (3) alloys of 3dtransition and noble metals with ordered structures, L1 0 FePt, L1 0 -FePd, L1 0 -CoPt, D0 3 -Co 3 Pt, τ-MnAl, (4) Co-Pt
alloys with metastable ordered structures, L1 1 -CoPt, B h CoPt, and (5) other compounds, Nd 2 Fe 14 B, Sm 2 Fe 17 N 3 ,α-MnBi, etc. Figure 14 shows the K u values of magnetic
materials plotted as a function of Ms2, which are citedfrom the references127-129). For application to PMR media,
the K u should be larger than the magneto-static energy(2πM s 2), thus a relation, K u > 2πM s 2, must be satisfied.
The materials above the K u = 2πM s 2 line shown in Fig.

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017115

INDEX

14 are basically applicable to PMR media. However, forpractical applications, several conditions such as foraligning the easy magnetization axis perpendicular tothe substrate, for achieving higher order degrees inordered alloys to get high K u values, for controlling the
crystal grain size suitable for magnetic recording,chemical stability, etc. must be considered. These
conditions are usually different depending on thematerial and thus must be investigated independently
for the respective magnetic material. Some of the basic
conditions in tailoring the crystal structure, the crystal
orientation, the chemical ordering, etc., have been
investigated by employing epitaxial magnetic thin films
prepared on nonmagnetic single-crystal substrates. For
example, it has been shown that the easy magnetization
axis of SmCo 5 ordered alloy and related compounds of
RT 5 -type ordered structure (R: rare-earth element, T: 3dtransition element), YCo 5 , GdCo 5 , SmNi 5 , SmFe 5 , etc.,
which possess K u values in the order 106 – 108 erg/cm3,
can be aligned perpendicular to the substrate byemploying a Cu(111) 130-137) or a Ru(0001) underlayer 138).
The effects of deposition temperature and film
composition on the formation of ordered alloys and thedegree of ordering are also investigated by using an
UHV-MBE with co-evaporation of R and T elemental
materials. Amorphous phases tend to be mixed with
RT 5 -ordered crystals when the composition shifts fromthe optimum ranges which depend delicately on the R-Tmaterial combination 130, 134). Substrate temperature
higher than 500 °C is necessary for the preparation of
RT 5 -type epitaxial thin films on Cu(111) and Ru(0001)
underlayers. Higher temperature, however, promotes
diffusion of metallic Cu or Ru to the magnetic layer 131,139), which generally lowers the K value. Further studiesufor forming a sharp interface while achieving a high orderdegree in the RT 5 structure in very thin films of less than
10 nm in thickness are apparently necessary for theapplication to future perpendicular recording media.
Chemical stability of RT 5 material is an issue to be solved
before practical applications, since rare-earth elements
react easily with oxygen or water vapor.
6.2 Magnetic multilayer films
Magnetic multilayers of Co and noble-metal element(Pd or Pt) can be prepared at relatively low temperatures.
From the investigations of Co/Pd multilayers depositedon single-crystal Pd underlayers of (001), (011), and (111)orientations at RT, it was shown that higher
perpendicular magnetic anisotropies were observed forthe epitaxial multilayer films prepared on fcc-Pd(111)underlayer 140, 141). The multilayer was composed of
repeated number of stacked fcc-Co(111)/fcc-Pd(111) bilayer with very thin Co-Pd alloyed regions at the Co/Pd
interface 142-144). Perpendicular magnetic anisotropy
increased with increasing the repetition number, and theorigin was interpreted to be coming from an interface
anisotropy and a magneto-crystalline anisotropy of CoPd(111) alloy-crystal whose lattice was slightly deformed

116(a) A1(b) L10
FePt, CoPt, etc.

(d) A3(e) D019
Co3Pt(c) L11,
CoPt(f) D2d
SmCo5, PrCo5, etc.

Fig. 15 Crystal structures of high Ku magnetic materials, basedon fcc-cubic (A1) and hexagonal (A3) lattices.
along the perpendicular direction in accommodation of alarge lattice misfit of 9.2% between the fcc-Co and the fccPd layers 145). As no ordered phase exists in the Co-Pd
binary phase diagram 146), the origin of perpendicularanisotropy is possibly due to a presence of very thinmetastable structure similar to L1 1 -phase which has ahigh K u and is recognized in a CoPt alloy material 147).
The magnetic anisotropy and the saturation
magnetization can be controlled in a certain range to beusable for magnetic recording media by varying the
thickness ratio of the Co and Pd layers and the repetition
number of Co/Pt bi-layer. Although a low substratetemperature for film formation and a high perpendicularmagnetic anisotropy are the advantages with multilayer
magnetic films, they are not easy to be tuned for
application to granular-type perpendicular media. The
multilayer films are considered to be more suitable forthe bit patterned media which involve nanometer scale
fabrications using chemical etching, ion-beam, or othersophisticated techniques 148-150).
6.3 Magnetic materials with ordered structures
The high K u magnetic materials of (2) - (5) groups
are with ordered structures consisting of more than twoelements with atomic arrangements specific to therespective ordered structures. Figure 15 shows the
typical crystal structures of high K u magnetic materials.
Table 2 lists the structural and magnetic properties ofhigh K u magnetic materials 127-129, 151-156). The degree of
atomic ordering in a magnetic material, where a higher
order degree corresponds to a higher K u value, depends
on the processing condition, in particular, thetemperature. Thus higher substrate temperatures are
generally necessary for film preparations. In these highK u magnetic materials, the alloys of 3d-transition andnoble metals with L1 0 -type ordered structure, L1 0 -FePt,
L1 0 -FePd, and L1 0 -CoPt alloys, have been mostintensively investigated for the granular-type recording
layers 127, 128, 157-161). Thus, high K u magnetic materials
with L1 0 ordered structure are focused in the present

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

Table 2 Structure and magnetic properties of high Ku magnetic materials 127-129, 151-156).
Crystal structure

Material
L10 phases
FePt
FePd
CoPt
τ-MnAl
B81 phase
α-MnBi

Tetragonal (fcc-based)a = 384 pm, c = 370 pm, c/a = 0.96
a = 381 pm, c = 372 pm, c/a = 0.98
a = 377 pm, c = 370 pm, c/a = 0.98
a = 392 pm, c = 357 pm, c/a = 0.91
Hexagonal
a = 428 pm, c = 611 pm, c/a = 1.43

L11 phase
CuPt-type rhombohedral (fcc-based)
CoPt
a = 534 pm, α = 61.5°
D019 phase
Rhombohedral (hcp-based)
Co3Pt
a = 367 pm
Compounds of rare-earth
Hexagonal
and transition elements
a = 499 pm, c = 398 pm, c/a = 0.80
SmCo5
a = 494 pm, c = 398 pm, c/a = 0.81
YCo5
a = 501 pm, c = 399 pm, c/a = 0.80
PrCo5
a = 498 pm, c = 397 pm, c/a = 0.80
GdCo5
Sm2Co17
Sm2Fe17N3
Nd2Fe14B

Disordered
A1 structure

a = 838 pm, c = 1221 pm, c/a = 1.45
a = 873 pm, c = 1264 pm, c/a = 1.45
Tetragonal
a = 879 pm, c = 1218 pm, c/a = 1.39(a) L10-(001) variant with (b) L10-(100), (010) variants
the c-axis perpendicular with the c-axis in-plane
Fe or Co
Pd or Pt
Substrate
Oxide [001]
[100]
[010]

Lattice mismatch : (aL10 – a)/a

(cL10 – a)/a

(aL10 – a)/aFig. 16 Possibilities of phase transformation from disordered
A1 structure to L10 ordered structure. The c-axis of L10(001)variant crystal is perpendicular to the substrate surface,whereas the c-axis is in-plane for the L10(100) and (010)
variant crystals. The lattice mismatch is supposed to give an
influence which type of variant crystal is likely favored.
review. Concerning the preparation and thecharacterization of other-type ordered alloy films of RT 5 ,L1 1 -CoPt, B h -CoPt, and DO 19 -Co 3 Pt, they are reported
in the references 130-139, 162-172).
There are technical issues associated with thematerials of L1 0 ordered structure to be solved for
application to PMR recording media. They are basicallysimilar to the case of conventional Co-alloy material withan hcp structure, but are more complicated. The
technologies for aligning the easy magnetization axis of
L1 0 -magnetic-crystal grain perpendicular to thesubstrate surface and for preparation of a flat medium
surface that can allow a very narrow spacing of less thana few-nm between the medium and a magnetic head. The
latter is a newly added issue associated with themagnetic material which will be used for HDDs with theK1, (10 7 erg/cm3)

Ms, (emu/cm3)

Tc, (K)

6.6
1.8
4.9
1.7

1140
1100
770
560

750
749
840
650

1.2

700

633

3.7 (S = 0.54)

800

-

2.0

1100

1190

11 - 205 - 6.5
8.3
4.6

910
850
930
498

1020
987
885
1014

4.2
8.6

838
1230

1190
749

4.6

1270

588

areal densities greater than 2 - 3 Tb/in2, where a newtype media, bit patterned media (BPM), is included.
Preparation of magnetic crystal with L1 0 -orderedstructure involves atomic ordering from a disordered fcc(A1 structure), which is different from the case of
traditional Co-alloys.
The possibilities of easymagnetization distribution associated with the phasetransformation from A1 to L1 0 ordered structure areshown schematically in Fig. 16. Here, the underlayer ofoxide with NaCl-type structure is assumed, consideringthat magnesium oxide (MgO: NaCl-type structure) has
been widely used as an underlayer for aligning the easymagnetization axis (c-axis) perpendicular to the film
plane. A1 structure based crystallographic notations forthe direction and the plane are employed in this paper for
the L1 0 -ordered structure to make easy comparison withthose of A1 structure, though the unit-cell of L1 0 -orderedstructure is tetragonal as shown in Table 2.
The easy magnetization axis of L1 0 -orederdmagnetic crystal is the c-axis. When an atomic ordering
takes place in a bulk crystal of A1 structure, there arethree possibilities of c-axis orientation with the ordered
L1 0 structure, L1 0 [100], [010], and [001] parallel withrespect to the original A1 [001], as shown in Fig. 16.
Three kinds of L1 0 crystal domain variant with the c-axis
perpendicular each other were observed in the highresolution TEM pictures of FePt film samples preparedon NaCl(001) substrates 173, 174). The presence of crystal
domains with the c-axis lying in the film plane
deteriorates the perpendicular magnetic anisotropy andwidens the magnetic switching field distribution of media,
both of which are undesirable for high density magnetic
recording applications.

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

117

INDEXThe other issue is the surface roughness of granulartype magnetic layer consisting of L1 0 ordered magnetic
crystal grains. Possibly due to a large difference insurface energy between the magnetic material (FePt,
FePd, or CoPt) and the oxide underlayers (MgO, etc) andto a high substrate temperature for enhancing L1 0
ordering, the magnetic material tends to form crystal
grains of spherical or facetted shape with different
diameters forming notable surface undulations, which
may hinder a closer approach of magnetic head in
performing a high density magnetic recording. Basictechnologies in controlling the crystallographic
orientation and in forming thin films with reduced
surface undulations are briefly explained in the following
sections.
(a) c-axis alignment of L1 0 -ordered magnetic crystal
In order to apply L1 0 -oredered alloy magnetic films to
PMR media, the c-axis of crystal grain must be aligned
perpendicular to the film surface. The alignment involves
a more complicated technology compared with the caseused for the conventional hcp-Co-alloy crystal. This isbecause that L1 0 -crystal structure is based on cubic (A1,fcc) structure with three crystallographic symmetries
compared with less crystallographic symmetries in the
hcp-structure (A3). When a magnetic material isdeposited on a substrate where an L1 0 -ordered alloy
crystal with the basal L1 0 -(001) plane parallel to thesurface is expected through an epitaxial growthmechanism, there are other possibilities of crystal growthwith the L1 0 -(100) and the L1 0 -(010) planes parallel tothe surface as well. This is due to that the atomic
arrangement of L1 0 -(001) is similar to that of otherplanes and the unit size of L1 0- (001) basal plane (a × a) isvery similar to that of other L1 0- (100) and L1 0- (010)
planes (a × c), where the c/a ratios of L1 0 crystals are very
close to unity ranging between 0.96 and 0.98 for theordered FePd, FePt, and CoPt materials as listed in Table
2. When the order degree decreases, the c/a ratio
approaches to 1.0 175) thus reducing the difference inlattice dimension between the L1 0 -(001) basal plane andthe other L1 0 -(010), L1 0 -(100) planes. As a result, threepossibilities of ordering from disordered A1 structure to
L1 0 -ordered structure exist as shown schematically in
Fig. 16.
Table 3 shows the distributions of L1 0 -crystal
variants which were investigated for FePd, FePt, and
CoPt magnetic materials 153, 176). 40 nm thick films were
deposited on the single-crystal substrates of MgO,
SrTiO 3 , and LaAlO 3 with (001) orientation at a substratetemperature of 600 °C. The lattice constants of substratematerial and the misfits with the magnetic materials arealso shown in Table 3. All the magnetic films grew
epitaxially on the (001) single-crystal substrates. The
L1 0 -variant configurations and order degrees were
carefully investigated by reflection high-energy reflection
diffraction (RHEED) and by X-ray diffraction (XRD)
employing out-of-plane and in-plane measurements.

118

Similar results on the variant configurations and theorder degrees are confirmed for the magnetic materialsof 40 nm thickness prepared by employing a two-stepprocess which consists of low temperature deposition at
200 °C followed by high temperature annealing at 600 °C
153, 177). When the magnetic materials were deposited at
200 °C, the films grew epitaxially with the disordered A1
structure. In the case of two-step process, an orderingfrom disordered A1 structure to L1 0 -ordered structure istaking place during the annealing process at 600 °C. On
the LaAlO 3 (001) substrate, two-types of variant crystal
were coexisting for the three magnetic materials. On thecontrary, FePd film grew with the L1 0 -(001) plane
parallel to the substrate surface on the MgO(001) and the
SrTiO 3 (001) substrates and the L1 0 -(100), (010) variants
were not recognized. However, two variant crystals were
coexisting on the LaAlO 3 (001) substrate. FePt and CoPt
films were consisting of two-types of variant crystal onthese oxide substrates. The data indicate that L1 0 crystal variant structure changes depending on the
combination of magnetic and substrate materials. It
also shows that the order degree, S, is different between
the L1 0 -variant crystals with the c-axis perpendicularand in-plane for a magnetic film. Higher order degrees
tend to be observed for the L1 0 -(001) variants than for
the L1 0 -(100) + (010) variants as shown in the cases of
CoPt on MgO, SrTiO 3 , and LaAlO 3 substrates, FePt on
SrTiO 3 and LaAlO 3 substrates, and FePd on LaAlO 3
substrate. Such difference seems to be caused by adelicate difference of atomic movement within amagnetic crystal during the ordering process, which ispossibly related with the internal stress of magnetic
crystal caused by a lattice misfit with the substratelattice.
Higher S value of L1 0 -(001) variant is observed in
the order FePd > FePt > CoPt on the oxide substrates,
indicating that the activation energy for L1 0 -ordering is
the lowest for the FePd material. When the order degree
is compared for the FePd films deposited on MgO, SrTiO 3 ,and LaAlO 3 substrates, higher S value is observed in theorder MgO > SrTiO 3 > LaAlO 3 . This tendency is similar
to the order of lattice misfit, (a L10 – a)/a, where a L10 is thelattice length of L1 0 -(001) basal plane and a is the lattice
constant of oxide crystal. A negative larger lattice misfit
along L1 0 [100], [010] can be interpreted to yield a higher
S value in the L1 0 -(001) variant. The rule is also
applicable for the S value of L1 0 -(001) variant in CoPt
film. The data for FePt film show almost similar tendency
but with a reversed order between the LaAlO 3 substrate(S = 0.30, misfit = - 0.8%) and the SrTiO 3 substrate (S =0.34, misfit = - 0.2%). The discrepancy may be within anexperimental error since the misfit difference is only
0.6%. The stress in L1 0 -(001) variant crystal caused by alattice misfit with the substrate is considered to havegiven a positive effect in enhancing ordering.
Table 3 shows that the ratio of L1 0 -variants with thec-axis perpendicular and in-plane differs depending on

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

Table 3 L10 crystal variants in magnetic thin films deposited on different oxide substrates153, 176).
Substrate

Property

V L10 (001)
S L10 (001)(aL10 – a)/aa = 421 pm V L10 (100), (010)
S L10 (100), (010)(cL10 – a)/a
V L10 (001)
SrTiO3 (001) S L10 (001)(aL10 – a)/aa = 390 pm
V L10 (100), (010)
S L10 (100), (010)(cL10 – a)/a
V L10 (001)
LaAlO3 (001) S L10 (001)(aL10 – a)/aa = 381 pm
V L10 (100), (010)
S L10 (100), (010)(cL10 – a)/a
MgO (001)

FePd

Magnetic material
FePt

100 %
0.69
- 9.7 %
0%-11.8 %
100 %0.61
- 2.6 %
0%
- 4.9 %
42 %0.51
- 0.2 %
58 %0.19
- 2.5 %

35 %0.34
- 8.7 %
65 %0.39
-12.1 %
40 %0.38
- 1.6 %
60 %0.33
- 5.2 %
42 %0.30
+ 0.8 %
58 %0.22
- 2.9 %

aL10 = 381 pm, cL10 = 372 pm aL10 = 384 pm, cL10 = 370 pm

CoPt

aL10 = 377 pm, cL10 = 370 pm

41 %
0.23
- 10.0 %59 %0.03-12.1 %
48 %0.20
- 3.3 %
52 %0.06
- 5.1 %
40 %0.15
- 1.0 %
60 %0.11
- 2.9 %V L10 (001): volume ratio of variant with c-axis perpendicular to the substrate. S L10 (001): order degree of L10 (001) variants.
(aL10 – a)/a : lattice mismatch between L10-crystal and the substrate.
V L10 (100) + (010) : volume ratio of variants with c-axes in-plane. S L10 (100) + (010) : order degree of L10 (100) + (010) variants.
(cL10 – a)/a : lattice mismatch between L10 (100) or (010) crystal and the substrate.

the combination of magnetic material and substrate.
Considering the volume ratio (33%) of L1 0 -(001) variant
expected when the phase transformation from A1 to L1 0
structure is taking place randomly, the nucleation of L1 0 (001) variant is apparently favored since the volume ratio
exceeds 33% for all the combinations. However, it is noteasy to discern which factor is playing a dominant role infixing the variant orientation. Only the L1 0 -(001) variant
is observed for the FePd films formed on MgO and SrTiO 3
substrates, but on LaAlO 3 substrate two kinds of variant
are coexisting. For FePd film, a larger negative lattice
misfit with the substrate seems to be working to promote
the nucleation of L1 0 -(001) variant. However, this rule isnot straightforwardly applicable to the FePt and the
CoPt cases. Some other material factors such asactivation energy of atomic diffusion in magnetic crystal,
L1 0 crystal nucleation density, and mechanical property
like Young’s modulus are considered to be working in
addition to the lattice misfit effect.
The variant configurations were investigated for the
40 nm-thick Fe 50 (Pt 1-x Pd x ) alloy films deposited on
MgO(001) substrate to study the influence of Pd to Pt
ratio on L1 0 variant structure 178, 179). The result showedthat the films with x > 0.5 consisted of L1 0 (001) crystals,
whereas the films with x < 0.25 included L1 0 (100)crystals with the c-axis lying in the film plane. Suchdata is apparently indicating that material related
factors are giving influences in fixing the L1 0 -variant
crystal orientation. The effect of strain caused by lattice
misfit on L1 0 variant configuration can be investigatedby employing film samples with different thicknesses,
since the tension force caused by a misfit between thesubstrate and the deposited magnetic material decreases
with increasing the magnetic layer thickness along thefilm growth direction. The effects of magnetic layer

thickness on the variant structure and the order degree
were investigated for FePd films deposited on MgO(001)substrates 180). Although the films thinner than 40 nm
consisted of L1 0 (001) crystal, a notable volume of
L1 0 (100) and L1 0 (010) variants started to be included
when the film thickness was increased to be larger than
100 nm. The easy magnetization axis varied fromperpendicular to in-plane, as the thickness was increased
from 100 to 500 nm where the volume percentage of
L1 0 (001) variant decreased down to about 20% which islower than that expected for random nucleation (33%).
The order degree of L1 0 structure also decreased withincreasing the magnetic layer thickness. Similar data
were obtained for FePt thin films 181). These results showthat the lattice strain caused by misfit with the substratematerial is playing an important role in fixing the L1 0
variant configuration and in enhancing the ordering
degree.
The effects of variant configuration and the orderdegree are also investigated for magnetic materials
formed on fcc (001) metallic underlayers such as Ag, Au,
Pt, Pd metals epitaxially grown on MgO(001) substrates
182). Although 10 nm thick FePt films deposited at a
substrate temperature of 600 °C grew with the c-axis of
L1 0 structure perpendicular to the substrate surface they
included L1 0 (100) variants. The order degree varied
widely in a range between 0.2 – 0.78 depending on the fcc
underlayer material, which is different from the cases ofoxide substrate. The difference may be attributed to thedifference of lattice strain in the magnetic material
around the interface between the metallic underlayer
and the oxide substrate. In the cases of metallic
underlayer, the strain will be more easily relaxed due tosimilar Young’s modulus and possibly due to element’s

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INDEX

Domain B

Domain C
Domain A
Fe(Pt,Pd)
Interface

┴

┴

┴

┴

MgO

1 nm

Fig. 17 High-resolution TEM image of interface between
MgO(001) substrate and L10-ordered Fe(Pt, Pd) epitaxial thinfilm. The thick and thin arrows respectively show the lines
enriched with high atomic number atoms (Pt, Pd) and with low
atomic number atoms (Fe) corresponding to the atomic
stacking along [001] of L10-ordered structure. The dotted lines
show the anti-phase boundaries and the mark ⊥ shows themisfit dislocation 181).

(b)

(a)
Epitaxial film
A1 (001)(c)

L10(100)

L10-crystal
nucleation

D-boundary

L10(100)
Lattice strain
APboundary

Substrate (001)

Substrate (001)

L10(001)

Substrate (001)Fig. 18 Schematic model to explain the nucleation and growthof L10-crystal variants in disordered A1 phase matrix. (a)Lattice of A1 crystal epitaxially grown on MgO(001) substrateis strained to accommodate the lattice mismatch, where the A1
lattice is expanded in lateral direction whereas it is contracted
in vertical direction. (b) Nucleation of L10-crystal is influenced
by the A1 lattice strain. L10(001) crystal with c/a < 1 will beformed preferentially in a region close to the substrate interface
while a more isotropic nucleation will take place in a regionaway from the interface. (c) Coalescence of growing crystal
variants forms domain boundaries (D-boundary). Neighboring
L10(001) crystals with a half-pitch difference (d002) in theatomic stacking forms an anti-phase boundary (AP-boundary)181).
diffusion across the interface of the two layers. Theseexperimental results indicate that the lattice misfit ofmagnetic material with respect to the substrate, theunderlayer, or the interlayer plays important roles indetermining the variant orientation as well as the degreeof L1 0 ordering.
The structure variation from A1 to L1 0 phase wasinvestigated by XRD and HRTEM observations for FePt
and FePd films on MgO(001) substrates as a function offilm thickness to study the L1 0 variant crystal growthmechanism 181). The results indicate that the resulting
L1 0 -variant structures are related with the degree of120

Fig. 19 AFM images of (a) FePt, (b) FePd, (c) CoPt, and (d)CoPd films grown at different substrate temperatures.
(e)–(h) Cross-sectional profiles along the red lines in (a)–(d),respectively 185).
lateral lattice strain existing in the A1 phase film andthat the preferential formation of L1 0 (001) variant is due
to larger lattice strains in the FePd film compared withthe case of FePt film.
The effects of lateral lattice strain on L1 0 -ordering
were also studied for 10-nm-thick FePt films depositedon (001) oriented oxide single crystal substrates of MgO,
MgAl 2 O 4 , SrTiO 3 , and LaAlO 3 at a low substratetemperature of 200 °C and then annealed at 600 °C 183),
where the lattice mismatch of A1-FePt crystal withrespect to the substrate ranged from -9.8% to +1.0%.
Higher L1 0 -ordering degree, S, is observed as MgO >MgAl 2 O 4 > SrTiO 3 > LaAlO 3 . The XRD and HRTEM
study has shown that the lattice strain existing in thedisordered (A1) film decreases during the hightemperature L1 0 -ordering process through introductionof misfit dislocations. The net misfit between L1 0 -crystal
and MgO(001) substrate is estimated to be less than 0.5%
for the L1 0 -ordered film. Figure 17 is an example ofinterface structure of L1 0 -ordered Fe(Pt,Pd) epitaxial
film formed on MgO(001) substrate. From the
distributions of variant domain boundaries observed by
HRTEM, the variant size is estimated to be 2 - 5 nm indiameter. A nucleation and growth model for explaining
the L1 0 -variant structure formation mechanism shownin Fig. 18 suggests a possibility in tailoring the variant
structure in L1 0 -ordered magnetic thin films bycontrolling the alloy composition, the layer structure, andthe substrate material.
(b) Reduction of surface roughness for L1 0 -orderedmagnetic thin films
When a magnetic material is deposited on a

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INDEX

substrate, atom migration, adhesion to the substratesurface, and formation of three-dimensional clusters takeplace. The film growth mode depends basically on adifference between the binding energy of deposited atoms,
Ε f , and the energy to bond to the substrate, E fs . In a caseE fs > E f , the film growth takes place with the Frank-van
der Merwe growth mode or the layer by layer mode,
whereas the film growth follows the Volmer-Webergrowth mode or the island-type growth mode for E fs < E f
184).
Most of the combinations of magnetic materials(FePt, FePd, CoPt) and oxide substrates belong to thelatter case. However, as the migration and diffusion
behaviors of deposited atoms are influenced by the film
formation process such as the deposition method and thesubstrate temperature, the resulting film nanostructure
varies depending on these conditions, primarily on thesubstrate temperature. Figure 19 shows the variationof surface profiles observed for 40 nm-thick FePt, FePd,
CoPt, and CoPd films deposited on MgO(001) substrates
at RT, 400 °C, and 600 °C. Epitaxial growth is realized
for all the films 185).
The nm-scale film surface
undulations are caused by the atomic movements
reflecting the surface energy variations of crystal facet as
schematically illustrated in Fig. 19 (e)-(h). When thesubstrate temperature is lower than 200 °C, the films arecontinuous with relatively smooth surfaces. The
arithmetical mean surface roughness, R a , values forthese films are less than 2 nm. However, with increasing
the substrate temperature, the surface roughness
increases to be larger than 10 nm and deep crevices areobserved for the L1 0 ordered films prepared at 600 °C,
which are caused by clustering of deposited magnetic
material and consisting of facetted surfaces. Surface
roughness caused by clustering at elevated substratetemperatures is noted for the L1 0 -(001) oriented
polycrystalline thin film media 186-188).
Considering that the film morphology is influenced
during the film deposition process, a two-step process
technique consisting of a low temperature film deposition
followed by a high temperature annealing for L1 0
ordering has been investigated to prepare an L1 0 orderedfilm with improved surface flatness 153, 174, 175, 177). Figure
20 compares the surface structures of FePd thin films of
2 - 40 nm thicknesses observed by atomic force
microscopy (AFM). The films were deposited on
MgO(001) substrates at 200 °C and then annealed at
600 °C. Smooth and continuous surface structures are
realized for the films with thickness larger than 5 nm,
showing an effectiveness of the two-step technique for the
preparation of flat thin films. The surface roughness R aof 40 nm-thick film prepared by employing the technique
is as small as 0.15 - 0.3 nm. On the contrary, isolatedisland-like cluster structures are observed for the films
thinner than 2 nm, where a clustering force of Fe and Pd
atoms has apparently overcome the lateral binding force
in the film. The degrees of L1 0 ordering in the films
prepared by using the two-step method are similar to
those of films directly deposited at 600 °C 180).

(a) 2 nm

30.5
nm
0

Ra: 4.4 nm(b) 5 nm3.0
nm0

Ra: 0.2 nm(c) 10 nm

2.5
nm
0

Ra: 0.2 nm(d) 40 nm

Ra: 0.2 nm
MgO[010]2.0
nm
0200 nmFig. 20 AFM images of L10 ordered FePd films with (a) 2, (b) 5,(c) 10, and (d) 40 nm thicknesses prepared on MgO (001)substrates by deposition at 200 °C followed by annealing at
600 °C 175).

(e) VN-TEM(d) MgO-TEM

L10-FePt (001) layer

L10-FePt (001) crystal

VN (001)

MgO (001)
5 nm

SrTiO3 (001)

5 nm

SrTiO3 (001)Fig. 21 AFM images of L10-FePt thin films of 2 nm in average
thickness formed on (a) MgO, (b) VC, and (c) VN epitaxial
underlayers 189). (d) Cross-sectional TEM image of L10-FePt
film of 2-nm average thickness formed on MgO layer and (e)cross-sectional TEM of 2-nm thick L10-FePt film formed on
VN epitaxial layer. (a) (d) and (c) (e) images are for the same
samples, respectively.
Although the two-step method is useful for the
preparation of L1 0 -ordered thin film with a flat andcontinuous surface down to a thickness around several
nanometers, island-like isolated crystals are formed
when the thickness is further decreased. In order toprepare a continuous thin film, the energy balanceworking on an isolated island needs to be considered. Asa deposited material behaves like a liquid droplet on thesubstrate surface at high temperatures, the contact angle
of island to the substrate surface is determined by abalance working between the surface energies ofsubstrate (γ SG ), deposited material (γ LG ), and interfacial
energy between these materials (γ SL ), which is known as
Young relationship 189). For achieving a lower contact
angle, the condition, γ SG > γ SL + γ LG , must be satisfied.
Thus it seems useful to employ a substrate material ofwhich surface energy is larger than that of magnetic
material (FePt, FePd, CoPt). Figure 21 compares themorphology of L1 0 -FePt(001) films of 2 nm average
thickness formed on MgO (γ SG = 1.4 J/m2), VC (2.8 J/m2),
and VN (2-3 J/m2) single-crystal underlayers of (001)orientation by employing the two-step method 190). The

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121

INDEX

10 9

MgO underlayer(substrate)(c)Stressdistribution

Magnetic film, (d)disordered phase (A1) L10-ordered film
MgO overcoat

High temperature(600 °C)

Fig. 22 Schematic model of L10 ordering from disorderedstructure, A1, magnetic film without overcoat (a), (b) and withovercoat (c), (d) epitaxially grown on MgO(001) substrate at
low temperature (200 °C) followed by high temperatureannealing (600 °C). The overcoat prevents clustering ofmagnetic material during the annealing process, thus providing
a smooth surface of L10-ordered film. The order degree, S, ishigher for the film with overcoat than that without overcoat 190,
191)

underlayer materials have a common NaCl-type crystal
structure with similar lattice constants of 0.41 - 0.42 nm
and L1 0 -FePt crystal grows epitaxially with the c-axis
perpendicular to the substrate surface. Very flat andcontinuous films with the surface roughness, R a < 0.4 nm,
are formed on the VC and VN underlayers, whereas
isolated L1 0 -FePt islands are formed on the MgOunderlayer (R a = 4.2 nm). An employment of underlayermaterial which has a surface energy larger than that ofmagnetic material seems to be crucial for the preparationof L1 0 -ordered continuous magnetic thin films with very
flat surfaces.
An another possibility to prepare L1 0 (001) oriented
magnetic thin films with flat surfaces is an employmentof a thin overcoat on top of the magnetic thin film
prepared by using the two-step method 191, 192). A thinovercoat, 2 nm thick MgO layer was formed at 200 °C on
FePt/MgO(001) samples with different FePt film
thicknesses before annealing at 600 °C. Very flat andcontinuous surfaces with the R a values less than 0.5 nm
were realized for the 1-10 nm thick L1 0 (001) FePt
epitaxial thin films. Here the overcoat is interpreted tohave worked in suppressing the atomic movements in thesurface region of magnetic film during the hightemperature ordering process at 600 °C preventing aformation of clusters. Furthermore, higher order degrees
are also observed for the magnetic films with the MgO
overcoats compared with the values of films of same
thickness without overcoats. Formation of MgO overcoat
layer is effective also in enhancing L1 0 ordering
providing a lateral tension force to the magnetic layerfrom the surface side. Figure 22 shows a schematic model
explaining the effect of overcoat. The basic findings
obtained from the experimental studies employing L1 0
epitaxial thin films are believed to be useful in tailoring
the high K u L1 0 magnetic thin films to be applied forpractical recording media.

122

KuV/kT=40 (at 300 K)[ Thermally stable ]

10 8

10

5

KuV/kT=80 (at 300 K)

Orderedalloys (SmCo5 )10 7Orderedalloys (L10, L11)

FePt, CoPt, FePd

10 6

MgO underlayer(substrate)

Low temperature(200 °C)

Ku (erg/cm 3 )(a)Magnetic film,(b) L10-ordered isolated crystal
disordered phase (A1)

[ Thermally unstable ]1

10

Co alloys

CoPt, CoCrPt, etc.

10 2

10 3

10 4

Volume (nm 3 )1

Cube diameter (nm)10Fig. 23 Relationship between Ku and magnetic crystal grainvolume for keeping the thermal stability condition, KuV/kBT> 40 – 80126).

7. Future possibilities of high K u magnetic
materials
In order to apply the high K u magnetic materials
with L1 0 ordered structure for PMR media, practicaltechnologies in tuning the nanostructure such as
magnetic crystal grain isolation, high L1 0 ordering
within a small crystal grain, appropriate control of grainsize, narrow distribution of grain size, etc. while
adjusting the magnetic properties, need to be developed.
In addition, surface flatness and mechanical durability
conditions that are related with very narrow physicalspacing of less than a few nanometers between amagnetic head and a medium must be simultaneously
satisfied. The development of such a highly sophisticatedtechnology is a big challenge. However, the authors
believe it will be possible likewise the past developmentsof media technology that enabled a dramatic areal
density growth of HDD through intensive experiments
carried out based on the basic science of materials.
Thermal stability of recorded information depends
on K u V/k B T. When T is room temperature (300 K), therelation between magnetic crystal grain volume V (nm3)and K u (erg/cm3) for maintaining the thermal stability isplotted as shown in Fig. 23 2, 126). Assuming the magnetic
grain to be a cube shape, the cube edge length is also
indicated as a measure in the figure. To decrease thelength while keeping the thermal stability K u V/k B T > 60,
K u must be increased from the order of 106 erg/cm3 which
is with the Co-alloy materials currently employed as
PMR media, to the order of 107 erg/cm3 or higher.
Although there are more than several magnetic
materials with the K u values higher than 107 erg/cm3 as
explained before, L1 0 -ordered magnetic alloy materials
seem to be the practical candidate for future PMR media
because of the excellent chemical stability and the higher
potential in tuning the nanometer level structure. The

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

minimum edge length of cube crystal for keeping thethermal stability is calculated to be 3.4 nm for the L1 0 ordered FePt material which possesses the K u value 6.6
× 107 erg/cm3 129). Assuming a physical separation of 0.5
nm between magnetic crystals, the areal density will be
increased up to around 40 Tb/in2. If a SmCo 5 crystal with
the K u value of 2 × 108 erg/cm3 is applied, an areal
density of more than 80 Tb/in2 seems feasible. Fromthermal stability and available magnetic material pointsof view, an areal density of 100 Tb/in2 which is more than
100 times larger than that of current areal density seems
to be achievable by carefully tuning the sub-nanometerlevel structure of high K u magnetic materials.
8. Summary
The developments of perpendicular magnetic
recording media technology are briefly reviewed from aview point in tailoring the nanostructure of recording
layer. Before establishing the PMR media structure that
employs the granular-type Co-alloy recording layer withan hcp structure, there was a long history of research anddevelopment (R&D) since the invention of PMR in 1976.
Including the early stage of R&D of Co-alloy recording
layer, the technological developments in aligning theeasy magnetization axis by hetero-epitaxial thin film
growth and in achieving magnetic separation between
the magnetic crystal grains by nonmagnetic elements
segregation around the grain boundaries are explained.
It is also shown that observations of composition andmagnetization structures have played key roles for the
improvement of medium nanostructure to be tuned
suitable for high density magnetic recording.
Although Co-alloy recording layer materials havebeen used since the invention of PMR and will be
continuously used up to an areal density beyond 1 Tb/in2
by adding further tunings to the medium structure andpossibly through a technological modification in therecording scheme 21, 22, 161), an employment of high K umagnetic material is necessary for the recording layer inkeeping the thermal stability of recorded information forfurther increasing the areal density well beyond 2 - 3
Tb/in2. However, an employment of high K u magnetic
material brings in new technological issues which are
different from those associated with conventional Coalloy materials. From such viewpoints, basic research
results related with technologies in controlling the easymagnetization and in keeping the surface flatness for
L1 0 -ordered magnetic thin films are also briefly
explained. Finally, future possibilities of PMR media
materials are discussed from the thermal stability ofrecorded information and the available magnetic
material points of view. The authors believe that an areal
density of 100 Tb/in2 which is more than 100 times larger
than that of current areal density seems to be achievableby carefully tuning the sub-nanometer level structure ofmagnetic materials.

Acknowledgements
The author (M. F.) would like to thank the colleagues
at Hitachi CRL who worked with the author during “thedip period” lasted for 10 - 15 years, and the members of
the committee 144 of the Japan Society for the Promotionof Science who worked toward the realization of PMRHDDs, of which names are listed as the reference paper
authors. The authors also thank the students of Chuo
University who patiently collected valuable scientific
data.
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Received May 05, 2017; Revised June 28, 2017; Accepted August17, 2017

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J. Magn. Soc. Jpn., 41, 127-131 (2017)<Paper>ofSynthesis and Magnetic Properties
Substituted W-type and Y-type Ferrites(Zn 2+ Ti 4+ )K. Watanabe, K. Kakizaki, and K. Kamishima
Graduate School of Science and Engineering, Saitama University, 255 Shimo-Okubo, Sakura-ku, Saitama 338-8570, Japan

We report the synthesis of (Zn2+Ti4+) substituted hexagonal ferrites, BaZn2(Zn0.5Ti0.5)xFe16-xO27 (W-type) and
Ba2Zn2(Zn0.5Ti0.5)xFe12-xO22 (Y-type), and the relationship between their composition and magnetic properties.
Substituted W-type ferrite was synthesized in the region of 0.0 ≤ x ≤ 1.5. A decrease in the Curie temperature of theW-type phase indicated that the exchange interactions between the Fe3+ ions weakened due to the substitution of Zn2+
and Ti4+ ions. Spontaneous magnetization (σS) at x = 0.5 was similar to that at x = 0.0. This reflects the equal
distribution of nonmagnetic Zn2+ and Ti4+ ions to spin down sites and spin up sites, respectively. At x > 0.5, σS
decreased clearly, which was possibly caused by a deviation from the collinear arrangement. Substituted Y-typeferrite was also obtained up to x = 0.5. The Y-type phase is more unstable than the W-type phase with respect to thissubstitution rate.
Key words: W-type ferrite, Y-type ferrite, (ZnTi) substitution, solid state reaction, X-ray diffraction,
Curie temperature, spontaneous magnetization, cation distribution
1. Introduction
The hexagonal ferrites have been investigated forelectronic devices at microwave frequencies since theirdiscovery in 1950s 1). Their availability at highfrequencies is caused by the large magnetic anisotropy
constants and the preference to several magnetized
directions, parallel or perpendicular to the hexagonal
c-axis. Most of ferrites also have low eddy current loss athigh frequency and corrosive resistance because they areoxides.
The crystal structure of hexagonal ferrites is based
upon the close-packed structure of large oxygen anions
and barium cations. The small transition metal cations(Fe3+ and Me2+) are located between the larger ions. The
structures of hexagonal ferrites can be represented by
the stacking combinations along their hexagonal c-axisof R, S, and T blocks 2). The chemical formulae for theseblocks are as follows: R = (BaFe6O11)2-, S = Me2Fe4O8 /(Fe6O8)2+, T = Ba2Fe8O14. The composition of the S block
corresponds with the two formula units of the spinel
ferrite. The major hexagonal ferrites are described by
the following symbols: RSR*S* for M-type (BaFe12O19),(RSR*S*S*)3 for X-type (Ba2Me2Fe28O46), RSSR*S*S* forW-type(BaMe2Fe16O27),
(TS)3
for
Y-type(Ba2Me2Fe12O22),
RSTSR*S*T*S*for
Z-type(Ba3Me2Fe24O41),
and(RSTSR*S*)3for
U-type(Ba4Me2Fe36O60) 3), 4). The superscript * denotes that thelayer is turned for 180° around the c-axis.
M-type barium ferrite (BaFe12O19) is widely used forpermanent magnets due to its large magneto-crystalline
anisotropy 5). In addition, many research groups haveinvestigated barium ferrite as microwave absorbers 6)-8)
and magnetic recording media 9) by modifying its
anisotropy. The control of its magnetic property isachieved by the replacement of Fe3+ ions by Al3+ 10)-13),

Cr3+ 13)-15) ions or pairs of (Me2+Mt4+) ions such as(Co2+Ti4+) 16, 17), (Zn2+Ti4+) 18), 19), (Mn2+Ti4+) 6), 20),
(Ni2+Sn4+) 21), (Co2+Sn4+) 22), (Co2+Zr4+) 23), 24), where the
Me and Mt represent divalent and tetravalent cations,
respectively. The substitution effect on anisotropy is sosignificant that even the hard magnetic property of
BaFe12O19 was changed to soft magnetic by substitutionof (Co2+Zn2+)Ti4+ for Fe3+ 25).
On the other hand, W-type and Y-type ferrites havedifferent magnetic properties from those of M-typeferrite due to their different stacking structures. W-typeferrite can possess higher saturation magnetization thanthat of M-type ferrite 26). Y-type ferrite has planar
anisotropy 27), which gives hope for high-frequency
application. Therefore, the control of the magnetic
properties of W-type and Y-type ferrites has been an
important issue.
To the best of our knowledge, however, there are few
investigations on the substituted W-type and Y-typeferrite. Maeda et al. reported the preparation by aconventional ceramic technique, magnetic properties,
and microwave absorption properties of the (Mn2+Sn4+)
and (Mn2+Zr4+) substituted W-type ferrite 28). Ul-ain et al.
reported the synthesis by sol-gel method and catalytic
activities of the (Mn2+Ti4+) substituted Y-type ferrite 29).
Also, substitution limits of Me2+Mt4+ for Fe3+ in W-type
and Y-type structures are unclear in contrast with thatin the M-type structure where complete substitution of
Co2+Ti4+ for Fe3+ is possible to form BaCo6Ti6O19 30).
We report the synthesis of the (Zn2+Ti4+) substituted
hexagonal ferrites, BaZn2(Zn0.5Ti0.5)xFe16-xO27 (W-type)
and Ba2Zn2(Zn0.5Ti0.5)xFe12-xO22 (Y-type), and the relation
between their composition and magnetic properties inthis paper. Here, we employed Zn2+Ti4+ cations for
substitution because an increase of the magnetization
was expected if nonmagnetic cations replaced Fe cations

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127

INDEXin spin down sublattices as observed in substituted
spinels 31).
2. Experimental procedure
The samples were prepared by solid-state reaction. We
used BaCO3, ZnO, TiO2, and α-Fe2O3 as starting
materials. They were mixed in the desired composition of
Ba : Zn : (Zn0.5Ti0.5) : Fe = 1 : 2 : x : (16-x) for the W-typeferrite and Ba : Zn : (Zn0.5Ti0.5) : Fe = 2 : 2 : x : (12-x) for
the Y-type ferrite with 0.0 ≤ x ≤ 2.0 in a ball-milling pot
for 24 h. The mixed powders were pressed into disks andthen pre-sintered at 900°C in air for 5 h. The products
were ground in a mortar and then pulverized for 10 min.
at the rate of 1100 r.p.m. in a vario-planetary mill. The
fine powders were pressed into disks again and thensintered at 1100-1350°C in air for 5 h. Finally, theproducts were ground into powders in a mortar.
The crystal structures were examined by powder X-ray
diffraction (XRD) with Cu-Kα radiation. The
magnetizations were measured with a vibrating sample
magnetometer (VSM), and a superconducting quantum
interference device magnetometer (SQUID).

sintered at 1250°C, provided that the W-type phase was
the minority phase for the sample with x = 1.5. The
samples with 0.0 ≤ x ≤ 1.5 sintered at 1300°C and 1350°C
consisted of the main W-type phase. On the other hand,at sintering temperature of 1350°C, the sample
contained the M-type and spinel minority phases. Since
this temperature is near to the melting point of the otherbarium-iron oxides (e. g. Ba2Fe2O5; Tm = 1370°C 32)), theW-type phase is likely to decompose. These results showthat the sintering temperature of 1300°C is the suitablecondition for the synthesis of the substituted W-typeferrites.
Fig. 2 shows XRD patterns of W-type composition
samples with Ba : Zn : (Zn0.5Ti0.5) : Fe = 1 : 2 : x : (16-x)
sintered at 1300°C. The samples had the W-type
structure at 0.0 ≤ x ≤ 1.5. At x = 1.5, there were some

3. Results and discussion
We identified the crystalline phases of the W-type
composition samples with 0.0 ≤ x ≤ 2.0 sintered at
1200-1350°C by XRD. Fig. 1 shows phase relationshipbetween sintering temperature and substitution rate x.
At 0.0 ≤ x ≤ 1.5, the W-type phase formed in the samples
Fig. 2 XRD patterns for products sintered at
1300°C whose raw materials were mixed incomposition of Ba : Zn : (Zn0.5Ti0.5) : Fe = 1 : 2 : x :(16-x) with 0.0 ≤ x ≤ 2.0.

Fig. 1 Relationship between obtained crystalline
phase and sintering temperature and substitution
rate x for samples whose raw materials were mixed
in composition of Ba : Zn : (Zn0.5Ti0.5) : Fe = 1 : 2 : x :(16-x) with 0.0 ≤ x ≤ 2.0. Phases in parentheses areminority phases.

128

Fig. 3 XRD patterns for products sintered at
1100°C whose raw materials were mixed incomposition of Ba : Zn : (Zn0.5Ti0.5) : Fe = 2 : 2 : x :(12-x) with 0.0 ≤ x ≤ 2.0.

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weak diffraction peaks of the M-type and spinel ferrites.
On the contrary, the peaks of the W-type phase
disappeared in the pattern at x = 2.0. It should be notedthat the first and second peaks of the W-type phase do
not overlap those of the spinel and M-type phases ascomparing the patterns between x ≤ 1.5 and x = 2.0. Fig.
3 shows XRD patterns of Y-type composition samples
with Ba : Zn : (Zn0.5Ti0.5) : Fe = 2 : 2 : x : (12-x) sintered at
1100°C. We obtained the Y-type phase in the samples
with x = 0.0 and 0.5. On the other hand, the diffraction
peaks of the Y-type phase weakened at x ≥ 1.0 andfinally disappeared at x = 2.0. The phase identification
was based on the diagrams of the BaO-ZnO-Fe2O3 33),
BaO-ZnO-TiO2 34) and BaO-Fe2O3-TiO2 35) systems,
which showed no corresponding compounds except forthe spinel and Y-type phases. The maximum
substitution rate of (ZnTi)/(Fe+(ZnTi)) to form the single
Y-type phase (0.5/12 ~ 4.2%) is smaller than that to form
the single W-type phase (1.0/16 ~ 6.3%). Therefore, theY-type phase is less stable than the W-type phase for(Zn2+Ti4+) substitution.

between Fe3+ ions weakened. This fact implies that Zn2+
and Ti4+ cations successfully replaced Fe3+ cations in theW-type structure at x ≤ 1.5.
Fig. 5 shows magnetization curves at T = 5 K (= –268°C) for W-type samples of BaZn2(Zn0.5Ti0.5)xFe16-xO27(0.0 ≤ x ≤ 1.5) sintered at 1300°C. The magnetization
curve of the substituted samples at x = 0.5 was almost
the same as that of the no-substitution sample (x = 0.0).
Then, the magnetization decreased with increasing thesubstitution rate from x = 0.5 up to 1.5. This result is
consistent with the fact that M-type ferrite showedsimilar Zn-Ti substitution dependence of spontaneous
magnetization 37), where the magnetization began toFig. 5 Magnetization curves at T = 5 K (= –268°C)
for samples sintered at 1300°C whose raw materials
were
mixed
in
compositionof
BaZn2(Zn0.5Ti0.5)xFe16-xO27 with 0.0 ≤ x ≤ 1.5.

Fig. 4 Temperature dependences of magnetization
at µ0H = 0.1 T for samples sintered at 1300°C whose
raw materials were mixed in composition of
BaZn2(Zn0.5Ti0.5)xFe16-xO27 with 0.0 ≤ x ≤ 1.5. Inset
shows Curie temperatures TC of those samples.
Fig.
4
shows
temperaturedependencesofmagnetization at µ0H = 0.1 T for W-type composition
samples of BaZn2(Zn0.5Ti0.5)xFe16-xO27 (0.0 ≤ x ≤ 1.5)
sintered at 1300°C. Inset of Fig. 4 shows the Curie
temperature TC of the W-type phase as a function of thecomposition x, where TC was estimated by linear
extrapolation to zero-magnetization from the inflection
point of the M-T curve. The Curie temperature TC of the
no-substitution sample (x = 0.0) was 305°C, which is inagreement with the previous report about BaZn2Fe16O27
single crystals 36). TC linearly decreased with increasingx at x ≤ 1.5, suggesting the exchange interactions

Fig. 6 Spontaneous magnetization and high field
differential susceptibility for samples sintered at
1300°C whose raw materials were mixed incomposition of BaZn2(Zn 0.5Ti0.5)xFe16-xO27 with 0.0 ≤ x≤ 1.5.

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

129

INDEX

decrease above a certain Zn-Ti substitution rate.
Fig. 6 shows spontaneous magnetization σS and highfield differential susceptibility χd as a function of thecomposition x. They were estimated by linear
extrapolation of the magnetization curves from a highfield region. The spontaneous magnetization σS slightly
decreased from x = 0.0 to 0.5. In the range of 0.5 < x ≤ 1.5,
the decrease of σS became more obvious. The
substitution effect on σS seems to reflect Zn2+-Ti4+ cation
distribution in the W-type structure.
In order to discuss the relationship between the change
in σS and the cation distribution in the W-type structure,
we consider the spin configuration of the W-type
structure. Table 1 shows the distribution of thetransition metal sites and the expected directions of themagnetic moments in the R and S blocks of the W-type
structure 38). Fig. 7 also shows magnetic structure of
W-type ferrite, corresponding to Table 1.
Concerning the cation distribution in the W-type
structure, we assume that the Zn2+ ions prefer the 4e
and 4fIV sites and that Ti4+ ions prefer the 12k sites. In
the case of the (ZnTi) substituted M-type ferrite 37), Zn2+
ions tend to enter the tetrahedral 4fIV sites in the S block.
Also, Ti4+ ions tend to enter the octahedral 4fVI (the R
block) sites and the 12k (the S and R blocks border) sites.
In our W-type ferrite case, even with no substitution
sample of x = 0.0, however, the positively charged part of
two S blocks of (Zn2Fe10O16)2+ and the negatively charged
R block of (BaFe6O11)2- are alternately stacked up as
SSRS*S*R*. The additional Zn2+ ions seem to occupy 4eand 4fIV sites in the S blocks like the ZnTi-doped M-typeferrite. In order to keep the local electronic balance inthe additional Zn2+ doped S blocks, Ti4+ ions may prefer
the 12k sites to the 4fVI sites because of the proximity to
the Zn2+ sites in the S block.
We discuss the effect of the nonmagnetic ions upon the
spontaneous magnetization σS depending on x. Between
x = 0.0 and 0.5, the occupation of spin down (4e and 4fIV)  sites by Zn2+ ions can be cancelled out by the occupationof spin up (12k) sites by Ti4+ ions. Therefore, thedecrease of σS is slight in this range. At x > 0.5, theobvious decrease of σS is possibly caused by the deviationfrom the collinear arrangement which can be dominant
at high substitution rates of Zn2+ and Ti4+ ions 37). To
ascertain our estimation, further experimental
investigations are needed to estimate the spin
configuration and the cation distribution by Rietveld
analysis on powder neutron diffraction 39) and X-ray
absorption fine structure analysis 40).
The high field differential susceptibility χd for thesubstituted W-type ferrite increases with increasing x as
shown in Fig. 6. It suggests that the magnetic structureof the substituted W-type ferrite becomes non-collinear
especially at x > 0.5. It is consistent with theinterpretation of the variation in σS discussed above.

130

Table 1 Coordination, number of ions, and spinorientation for various cation sublattices of W-typeferrite.
Block

Sublattice

Coordination

Number
per
block

Spin

S-S

6g

Octahedral

3

Up

S

4e

Tetrahedral

2

Down

S

4f

Octahedral

2

Up

S

4fIV

Tetrahedral

2

Down

S-R

12k

Octahedral

6

Up

R

4fVI

Octahedral

2

Down

2d

Trigonal-bipyramidal

1

Up

R

Fig. 7 Magnetic structure in half of unit cell (informula unit ) of W-type ferrite.
4. Conclusions
We report the synthesis condition and the magnetic
properties of the (Zn2+Ti4+)-substituted W-type andY-type hexagonal ferrites that have hardly been
investigated so far. The synthesis condition shows thesubstitution limits of Zn2+Ti4+ for Fe3+ in the W-type andY-type ferrites, suggesting the stability of the W-type
and Y-type structures.
The W-type composition samples sintered at 1300°C
had the W-type structure in the region of 0.0 ≤ x ≤ 1.5.
The Curie temperature of the W-type phase decreased
linearly with increasing x at x ≤ 1.5. It suggests that theexchange interactions between the Fe3+ ions weakened
due to the substitution of nonmagnetic Zn2+ and Ti4+ ionsin the W-type structure. Spontaneous magnetization σS

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEXat x = 0.5 was similar to that at x = 0.0. It reflects theequal distribution of nonmagnetic Zn2+ and Ti4+ ions tospin down sites and spin up sites, respectively. At x > 0.5,
σS decreased clearly, which is possibly caused by thedeviation from the collinear arrangement. It isconsistent with the increase in the high field differential
susceptibility χd with x.
The Y-type single phase was also obtained for the
Y-type composition samples up to x = 0.5. The Y-typephase is more unstable than the W-type phase withrespect to this substitution rate.
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Sonntag: J. Magn. Magn. Mat., 192, 83-99 (1999).
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Received Jul. 20, 2017; Revised Sept. 3, 2017; Accepted
Sept. 19, 2017

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INDEX
J. Magn. Soc. Jpn., 41, 132-139 (2017)<Paper>Low permeability composite magnetic core transformer
with high coupling coefficient and its application
to PFM controlled quasi-resonant mode flyback-type DC-DC converter
K. Sato*,**, T. Sato**, M. Sonehara**, and H. Takeuchi**

Nagano Prefecture General Industrial Technology Center, 1-3-1 Osachi-katamacho, Okaya, Nagano 394-0084, Japan**

Faculty of Engineering, Shinshu Univ., 4-17-1 Wakasato, Nagano, Nagano 380-0928, JapanWe describe a low permeability composite magnetic core transformer with a high coupling coefficient and its
application to a PFM controlled quasi-resonant mode flyback-type DC-DC converter. To increase the coupling
coefficient between primary and secondary windings embedded in a carbonyl-iron powder/epoxy composite magnetic
core with a relative permeability of 6, three types of winding layouts were investigated by using a 2D-FEA numericalsimulation and the characteristics were confirmed experimentally. The primary and secondary windings that werelaid out adjacent to each other were the best choice as it had the highest coupling coefficient at over 0.99. By
applying the composite magnetic core transformer with this coupling coefficient to a 48 V input/12 V output PFM
controlled quasi-resonant mode flyback-type DC-DC converter, the maximum efficiency was 92.5% at an output
power of around 28 W, and an efficiency of 90.0% was obtained at an output power of 60 W.
Key words: low permeability magnetic core, carbonyl-iron powder/epoxy composite magnetic core transformer,
winding layout, coupling coefficient, DC-DC converter, efficiency
1. Introduction
Recently, SiC/GaN power devices have been applied
to switching power converters. It is because they makepossible to have good advantages such as lower ON
resistance, higher speed switching operation and higher
temperature operation than the Si-based power devices1), 2). By using the novel power devices, it is possible not
only to increase the power conversion efficiency of theconverter but also to downsize the magnetic devices
used in the switching power converters since theswitching frequency can be increased to MHz band ortens of MHz. However, it is considered difficult to use
the magnetic devices consisting of dust core or Mn-Zn
ferrite core for such the very high frequency switching
power converter. M. Harrison3) stated that the possible
magnetic core for MHz band or beyond 10 MHz power
conversion would be Ni-Zn ferrite core only.
The authors are developing a novel magnetic core
as a candidate substitute for the Ni-Zn ferrite core, forexample, the carbonyl-iron powder/epoxy composite
magnetic core was developed4) and was applied toswitching power converters5),6). In order to increase thevolume resistivity of the composite magnetic core body,the surface-oxidation for the carbonyl-iron powder wasvery
effective4).
Thoughthesurface-oxidized
carbonyl-iron powder/epoxy composite magnetic core
had a low relative permeability of about 6, the magnetic
loss was very small up to 10 MHz in frequency. By
applying the composite magnetic core inductor to a 1
MHz switching buck-type DC-DC converter using GaN
power device, a maximum power conversion efficiency of
93% was obtained5). It was also reported that acomposite magnetic core transformer with copper-cladpolyimide tape windings embedded in the composite

132

magnetic core could realize a coupling coefficient k >0.99 even when using such the low permeability
composite magnetic core6). Furthermore, by applying
the composite magnetic core transformer to a PWM(Pulse Width Modulation) controlled flyback-type
DC-DC converter, the maximum power conversion
efficiency was 91% at 400 kHz switching and 85% at 1
MHz switching6).
In general, the low permeability magnetic material
is not suitable for transformer core because of large
leakage inductance owing to the low coupling coefficient
between the windings. This study was done to develop
the composite magnetic core transformer with highcoupling coefficient and to apply it to a PFM (Pulse
Frequency Modulation) controlled quasi-resonant mode
flyback-type DC-DC converter7).
This paper describes an investigation on the designof low permeability composite magnetic coretransformer with high coupling coefficient on the basesof the two-dimensional finite element analysis (2D-FEA)
numerical simulation and experiments. In addition, aPFM controlled quasi-resonant mode flyback-type
DC-DC converter using the fabricated composite
magnetic core transformer is also shown.
2. Surface-oxidized carbonyl-iron powder/epoxy
composite magnetic core transformer
2.1 Properties of surface-oxidized carbonyl-iron
powder/epoxy composite magnetic coreAs the soft magnetic metal powder for thecomposite magnetic core, non-reduction carbonyl-iron
powder with a mean diameter of 1.6 m was used. Inorder to prevent the electrical contact between theadjacent powders in the composite magnetic core and to

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

suppress the over-lapped eddy current loss, highresistive oxidized layer of several tens of nm wasintroduced on the surface of the powders by thermaloxidation in air at 200 degrees-C for 6h4). The composite
magnetic core was fabricated by the casting andpost-curing using the composite slurry with thesurface-oxidized carbonyl-iron powder and epoxy
precursor solution. Since the nonmagnetic epoxy resin
between the magnetic particles behaves as a magnetic
gap, the composite magnetic core is hard to bemagnetically saturated due to its low permeability andhas superior superimposed DC current characteristics
when using the magnetic devices, so that thetransformer core can be gapless.
Fig. 1 shows the static magnetization curve of thesurface-oxidized carbonyl-iron powder/epoxy composite
magnetic core measured by using a VSM (Riken Denshi
Co., Ltd.; BHV-55). The saturation magnetization was
about 1 T. Though the composite magnetic core hadnonmagnetic epoxy resin, it exhibited a high saturation
magnetization around 1 T, so it was considered that theenergy density could be increased and the transformer
core could be downsized.
Fig. 2 shows the frequency dependence of therelative complex permeability of the surface-oxidized
carbonyl-iron powder/epoxy composite magnetic coremeasured by using an impedance/material analyzer

Magnetization M [T]

1.5
1
0.5
0
-0.5
-1
-1.5
-800
-400
0
400
800
Applied magnetic field H [kA/m](Agilent; 4291B) and a coaxial waveguide. The real partof the complex relative permeability r’ was low as about
6, but the imaginary part r” was as small as 10-2 orderat 10 MHz or less. Although not measured below 1 MHz
in frequency, the relative permeability r’ of about 6 wasconsidered to be constant at hundreds kHz band and themagnetic loss r” was considered to be small enough at
same hundreds kHz band.
The volume resistivity of the surface-oxidized
carbonyl-iron powder/epoxy composite magnetic core
was about 100 m4), which was 100 times higher thanthat of Mn-Zn ferrite core. Therefore, it was consideredthat the eddy current loss of the transformer core could
be significantly suppressed at least up to 10 MHz.
2.2 Transformer structure
The output voltage of the quasi-resonant mode
flyback-type DC-DC converter is controlled by changing
both the on pulse width and the switching frequency
according to the output current7). As the output power
becomes larger, the on pulse width expands and theswitching frequency decreases. Also, the soft switching
operation can be realized by using the resonance of thedrain-source voltage of the MOSFET due to the
inductance of the transformer primary winding and thecapacitor connected in parallel to the MOSFET.
Therefore, in order to increase the switching frequency,
it is necessary to reduce the inductance of the primarywinding. Since the composite magnetic core underinvestigation has a high saturation magnetization andthe operating range ΔB of the transformer can beenlarged, even when the switching frequency decreases
under heavy load, the magnetic saturation occurs
hardly, which is due to the low inductance (lowpermeability composite magnetic core).
Fig . 3 sh o w s th e c o mposite magnetic coretransformer structure under investigation in this paper.
The transformer winding was embedded in thecomposite magnetic core with a 30 mm-square size andFig. 1 Static magnetization curve of surface-oxidized
carbonyl-iron powder/epoxy composite magnetic core.

Permeability r', r"

10
5

r'

1
0.5

r"
0.1
0.05
0.01
1M

10M
100M
Frequency f [Hz]1G

Fig. 2 Complex relative permeability of surface- oxidized
carbonyl-iron powder/epoxy composite magnetic core.

Fig. 3 Composite magnetic core transformer withembedded windings.

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017133

INDEX

type-A
Primary winding
Secondarywinding

type-BPrimary
winding-1

Primary
winding-2

Secondarywinding

type-CPrimary winding
Secondarywinding-1

0

Secondarywinding-2

Cross-sectional
structure of winding

2

2

2

1

1

1

Fig. 4 Three types of winding layouts for composite
magnetic core transformers with embedded windings.
10 mm in thickness. A copper-clad polyimide tapehaving a 70 m thick, 5 mm wide copper conductor line
formed on a 100 m thick, 5.2 mm wide polyimide tape
was used as the primary and secondary winding.
In this paper, inductance of the primary winding
was designed so that the switching frequency was about
500 kHz when the converter output power of 60 W8). Fora reason of the fabrication process of the tape wound
coil, the number of turns was determined to be an evennumber, the primary winding turn was 8, and thesecondary winding turn was 4. Since the composite
magnetic core had a low relative permeability, winding
layout becomes more important than when using othermagnetic core materials having a high relative
permeability. In general, when the primary andsecondary winding were wound in an interleaved
structure, the coupling coefficient between the windings
could be increased9)-11). It was easy to wind in aninterleaved structure when the number of turns ofprimary winding and that of secondary windings wereequal. However, since the number of turns of primaryand secondary winding of the transformer discussed inthis paper was different, three types of transformers
having different winding layout as shown in Fig. 4 werestudied. In the cross-sectional view of windings in thelower figure of Fig. 4, “P-1” to “P-8” means from firstturn to eighth turn of the primary winding, respectively.
Similarly, “S-1” to “S-4” means from first turn to fourth
turn of the secondary winding, respectively.
Type-A transformer consisted of one primarywinding and one secondary winding. The secondary
winding was concentratedly laid out near the center ofthe primary winding when looking at the cross-sectional
structure of the windings. Therefore, the innermost
conductors of the primary winding and the secondary
winding were not adjacent to each other.
Type-B transformer consisted of two primarywindings and one secondary winding. The secondary
winding was sandwiched between two primary windings

134

and wound four turns. In order to make the primarywinding eight turns, the winding end of one of the two
primary windings and the winding start of the otherwere connected in series outside the magnetic core.
Hence, the innermost conductors of the primarywinding and the secondary winding were adjacent toeach other, and the secondary winding were distributedand evenly laid out between the primary winding whenlooking at the cross-sectional structure of the windings.
Type-C transformer consisted of one primarywinding and two secondary windings. After the firstsecondary winding was wound four turns with theprimary winding, the second secondary winding waswound continuously four turns with the primarywinding. In order to make the secondary winding four
turns, two secondary windings were connected inparallel outside the magnetic core. As shown in Fig. 4,
the innermost conductors of the primary winding andthe secondary winding were adjacent to each other andthe secondary winding of type-C were alternately andevenly laid out between the primary winding whenlooking at the cross-sectional structure of the windings.
The inductance and the resistance of the two secondary
windings were different because the lengths and innerdiameters were different. Hence, it was considered thatan imbalance in the current flowing of the twosecondary windings may occurs. The following 2D-FEA
numerical simulation and experiments took thisbehavior into account.
2.3 Fabrication procedure of composite magnetic coretransformer
The fabrication procedure of the composite
magnetic core transformer is as follows. An air-corewinding having an inner diameter of 17.5 mm was
fabricated. Air-core windings were covered withpolyimide tape to prevent contact between the winding
and the composite magnetic core material. The
composite
slurrywas
preparedbymixingsurface-oxidized carbonyl-iron powder and two
component epoxy precursor solution. The slurry wasfilled in a mold in which the winding was arranged andthen cured at 120 degrees-C for 5 h. Thereafter, the sizeof the composite magnetic core was adjusted bypolishing. As the lead wire from the winding inside thecomposite magnetic core, the same copper-cladpolyimide tape was used.
3. Operation analysis for composite magnetic coretransformer
3.1 Analysis model and conditions
In order to clarify the operation characteristics ofthe three types of transformers, 2D-FEA numericalsimulation (Ansys; Maxwell) was introduced. As shownin Fig. 3, the analytical region was set to be the righthalf of the transformer cross section because of thesymmetry of the device structure. The analytical modelfor deriving various parameters was a model rotated

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

Table 1 Parameters for 2D-FEA numerical simulation
for three types of composite magnetic core transformers.

one time around c-c’ by post processing of the simulator.
For this reason, the transformer core of the analyticalmodel had a cylindrical shape with a diameter of 30 mm
and a thickness of 10 mm, and the outer shape wasdifferent from the fabricated transformer. However, it
was considered that there was little influence on theanalysis result since the magnetic circuit of thetransformer core had almost the same structure. The
2D-FEA
could
notaccurately
simulatethethree-dimensional structure of the transformer
including the influence of the lead wire etc. Therefore,
quantitative consideration was difficult. However, it
was considered sufficient to consider the relationshipbetween the winding layout of the transformer and its
operating characteristics under the same magnetic
circuit structural condition.
Table 1 shows the properties used for analysis. The
copper loss of the winding was calculated with taking
the losses due to the skin effect and the proximity effectinto account. The magnetic properties of the composite
magnetic core were regarded as linear, and themagnetic hysteresis was not considered. The core loss
was calculated from the imaginary part of the complex
permeability shown in Fig. 2 and set as the loss tangent
tan (=r”/r’) in the 2D-FEA numerical simulation.
Since the measurement variation of the imaginary partof complex permeability in the low frequency below 10
MHz was large, the value of tanat 10 MHz was also
used for the analysis below 10 MHz. Analysis was
carried out under the constant voltage source excitationby using “Maxwell Circuit Editor” function capable ofcoupled analysis of FEA and the circuit simulation, thecircuit diagram as shown in Fig. 5 was introduced, and

5
2.5
0

Meas. Calc.
: type-A: type-B: type-C1

0.1

Meas. Calc.
: type-A: type-B: type-C

0.01
1

Coupling coefficient k

Fig. 5 Relationship between 2D-FEA numericalsimulation for composite magnetic core transformer andequivalent circuit.

7.5

-2.5
10
Resistance Rs []Conductor line
(Copper)Insulator layer(Polyimide)

Constant
Permeability: r’ = 6 (const.)
Loss tangent: tan= 0.002 (const.)
Conductivity: 0.01 S/m (100 m)
Conductivity: 5.8 × 107 S/mThickness: 70 m, Width: 5 mm
Conductivity: 0 S/m
Thickness: 100 m, Width: 5.2 mm

Inductance Ls [H]

Material
Composite
magnetic core

10

0.995
0.99
0.985
Meas. Calc.
: type-A: type-B: type-C

0.98
0.975
100k

1M
Frequency f [Hz]10M

Fig. 6 Frequency characteristics of three types oftransformers.
the operation characteristics of the transformers were
simulated. In Fig. 5, a sinusoidal voltage source wasconnected to the primary side of the transformer, and a
resistive load was connected to the secondary side. The
input voltage V1 was set to 48 Vrms.
Since the measured value of the relative
permeability in Fig. 2 was a small signal characteristic
and tan was regarded as constant value, the calculated
power loss may be estimated to be smaller than theactual power loss. However, for the purpose ofcomparing the differences in characteristics due to thewinding layout, the essential problem was considered tobe small.
3.2 Electrical characteristics
Fig. 6 shows the equivalent series inductance,
equivalent series resistance and coupling coefficient of
the primary side of the transformer at the time of
no-load. In Fig. 6, the plots are measured values byusing an impedance/gain-phase analyzer (Agilent;
4194A) and the solid lines are calculated ones. The

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

135

INDEX

Table 2 Measured and calculated DC resistance of eachwinding of transformers.

type-B
type-C

Measured26.6 m14.5 m28.8 m14.8 m27.4 m
7.2 m

Calculated
24.0 m12.1 m24.1 m12.0 m24.7 m
6.3 m

99
Efficiency [%]type-A

Primary
SecondaryPrimary
SecondaryPrimary
Secondary

100

98
97

Output voltage V2 [V]f =1 MHz: type-A: type-B: type-C

24.1

23.9

0

2
4
6
8
Output current I2 [A]10

Fig. 7 Analysis results of load current dependence of
output voltage for three types of transformers.
coupling coefficient was obtained by the secondary side
open/short circuit method.
As shown in Fig.6, the equivalent series inductanceof the fabricated transformer was about 4.6 Hregardless of the winding layout. The self-resonancefrequency of type-B and type-C was lower than that of
type-A because of the parasitic capacitance increase dueto the winding layout5).
Table 2 shows the measured and calculated DC
resistance of each winding of the transformers. As
shown in Table 2, the DC resistance of the primarywindings was almost equal regardless of the winding
layout. However, since two windings were connected inparallel, the DC resistance of the secondary winding oftype-C was almost half.
Although there was the difference between themeasured values and the calculated ones of eachparameter, this was considered to be due to thepositional deviation of the winding in the thickness
direction of the magnetic core5) and the influence of theparasitic parameters of the lead wire from the winding
inside the composite magnetic core not considered in theanalytical model. Also, since 2D-FEA was only magnetic
field analysis, the influence of self-resonance did notappear in the calculation results. However, thetendency of the electrical characteristics of the threetransformers generally agrees with the measured values
and the calculated ones, it was considered that it waspossible to compare the difference in characteristics dueto the winding layout.

136

0

2
4
6
8
Load current I2 [A]10

Fig. 8 Analysis results of load current dependence ofefficiency for three types of transformers.

24

23.8

: type-A: type-B: type-C

96
95

24.2

f = 1MHz

3.3 Load current dependences of transformer
Fig. 7 shows the calculated load current dependenceof the output voltage V2 of transformer. As shown in Fig.
7, the output voltage V2 of type-A exceeded 24 V in the
no-load and light load, which was higher than the value
calculated from the turn ratio of 8:4. It was considered
to be due to the low relative permeability of thecomposite magnetic core. Since the relative permeability
was low, the main magnetic flux was considered not
only to have the component in the composite magnetic
core body but also to include the leakage component in
the region near to the primary winding. As a result, theeffective inner diameter of the secondary winding wasconsidered to be enlarged and the inductance also
increased. This was equivalent to the increase of thenumber of turns of the secondary winding. Hence, it wasconsidered that the output voltage became higher thanthe value calculated from the turn ratio of 8:4. This
operation suggests that the operating point of thetransformer may unintentionally change from theoperating point designed based on the specifications of
the converter. On the other hand, in type-B and type-C,
since the innermost conductors of the primary winding
and the secondary winding were adjacent to each otherand there was almost no difference in inner diameter,
the output voltage was almost equal to the calculated
value from the turn ratio of 8:4.
The output voltage of type-A transformer decreased
more than the other two types of transformers. It wasconsidered because type-A had lower coupling
coefficient between primary and secondary winding
than the other two types of transformers, and its
leakage inductance was larger than the other two
transformers.
3.4 Power efficiency and loss
Fig. 8 shows the calculated power efficiency of thetransformer at 1 MHz. The maximum efficiency of 99%
or more in any transformer was obtained. Since the loss
due to the excitation current was mainly at no-load andlight load, the type-A transformer with the lowest

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

frequency decreases with increasing load current, there
is a possibility that high efficiency can be maintained in
the wide load condition when the composite magnetic
core transformer is applied to such converter.

Power loss [W]

10

1

0.1

I2 = 0 A: type-A: type-B: type-C

Power loss [W]

0.01
10

: type-A: type-B: type-C

1

I2 = 4 A: type-A: type-B: type-C

0.1
10Power loss [W]I2 = 2 A

I2 = 6 A: type-A: type-B: type-C

1

I2 = 8 A

0.1
100k: type-A: type-B: type-C

I 2 = 10 A
: type-A: type-B: type-C

1M
10M
1M
10M 100k
Frequency f [MHz]Frequency f [MHz]Fig. 9 Analysis results of frequency dependence ofpower loss for three types of transformers.

equivalent series resistance of primary winding at 1
MHz had the highest power efficiency. On the other
hand, at heavy load, type-C with lowest secondary
winding resistance and highest coupling coefficient
between windings had the highest power efficiency.
Fig. 9 shows the calculated values of the frequency
dependence of the power loss of the transformer. The
frequency at the minimum power loss varied with theload current value, and the composite magnetic coretransformers under investigation in this study exhibitedthe minimum power loss obtained at 500 kHz to 2 MHz
according to the load current intending for the PFM
controlled quasi-resonant type flyback converter in thisstudy. The power loss in the low frequency increased
because the winding loss increased and the input power
factor decreased, which was considered to be owing to
the low inductance and the increase of excitation
current with lowering frequency. On the other hand, it
was considered that the power loss increased in the highfrequency due to the skin effect and proximity effect of
the windings.
With increasing load current I2, the frequency at
the minimum power loss shifted to the low frequency
side in any transformer. In the quasi-resonant mode
flyback-type DC-DC converter, since the switching

3.5 Magnetic flux density distribution in transformer
Fig. 10 shows the calculated magnetic flux density
distribution of the three types of transformers when thefrequency of 1 MHz and load current of 10 A. In Fig. 10,
the in-phase flux component with the exciting current
and 90-degrees phase-lag component are shown, where
the stored energy in the transformer corresponds to thein-phase flux component and the energy loss
corresponds to the 90-degrees phase-lag component.
Since the magnetic loss due to tan (=r”/r’) of thecomposite magnetic core was considered to be smallenough at 1 MHz, the energy loss was considered to bemainly due to the windings.
From upper figure in Fig. 10, it was found that thein-phase magnetic flux density with the exciting current
occurred mainly inside the winding and its spatial
distribution was almost same in three types of winding
layouts. On the other hand, as shown in the lower figure
in Fig. 10, 90-degrees phase-lag magnetic flux density
component shows a difference depending on the winding
layout. In the type-A, 90-degrees phase-lag magnetic
flux density distribution was generated between the
conductors where the primary winding and thesecondary winding were adjacent to each other. In
particular, the magnetic flux was also distributed in thecomposite magnetic core at the beginning and the end of
the secondary winding. Also in type-B, 90-degrees
phase-lag magnetic flux was generated between the
conductors where the primary winding and thesecondary winding were adjacent to each other, but no90-degrees phase-lag magnetic flux was generated in
the composite magnetic core. Type-C had little
90-degrees phase-lag magnetic flux in the composite
magnetic core and the 90-degrees phase-lag magnetic
flux density between windings was smaller than type-B.
It was considered that the 90-degrees phase-lag
magnetic flux density component was due to the energy
loss owing to the difference between the primarymagnetomotive force and the secondary magnetomotive
force. It was suggested that offsetting both themagnetomotive forces was incomplete in type-A.
Furthermore, the leakage magnetic flux due to theoffset magnetomotive force was considered to be a causeof an increase in eddy current loss in the winding.
Therefore, it was inferred that type-A had larger
winding loss at the large offset magnetomotive force
when the heavy load, compared with the other two typesof transformers. It was considered that the difference
between the losses of type-B and type-C was not only
due to the winding resistance of the secondary winding
but also due to the influence of the eddy current loss
caused by the leakage magnetic flux.

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017137

INDEX95

10

Efficiency  [%]

90
85
1
80
75
70

Fig. 11 Circuit diagram of PFM controlled
quasi-resonant mode flyback-type DC-DC converter.
4. Application to PFM controlled quasi-resonant mode
flyback-type DC-DC converter
4.1 Evaluation circuit diagram and conditions
Since the PFM controlled quasi-resonant mode
flyback-type DC-DC converter7) used in this study
operates under the critical current mode and soft
switching operation, it is expected to realize both higher
conversion efficiency and higher switching frequency.
Fig. 11 shows the circuit diagram of the PFM
controlled quasi-resonant mode flyback-type DC-DC
converter with the composite magnetic core transformer
under investigation. While the drain-source voltage
waveform of the MOSFET was being confirmed with theoscilloscope, the switching frequency and on-time ratio
were adjusted by the function generator (YOKOGAWA;
FG420) so that the output voltage was constant at 12 V
even if the output current was changed.
4.2 Experimental results and discussion
Fig. 12 shows the measured power conversion

138: type-A: type-B: type-C0

10 20 30 40 50 60 70
Output power POUT [W]

Switching frequency fsw [MHz]type-A
type-B
type-C
Fig. 10 Magnetic flux density distributions of three types of transformers. Upper figure represents in-phase flux
component with exciting current, and lower figure shows flux component with 90-degree phase-lag from exciting
current when frequency was 1 MHz and load current was 10 A.

0.1

Fig. 12 Power conversion efficiency and switchingfrequency of converter using composite magnetic coretransformers with different winding layouts.
efficiency and switching frequency of the converter
using three types of composite magnetic coretransformers. In this evaluation condition, theswitching frequency varied from approximately 400 kHz
to 2 MHz, which was roughly in agreement with thefrequency band in which the power loss of the composite
magnetic core transformer was small in theabove-described operation analysis for the transformers.
As shown in Fig. 12, when type-C was used, themaximum power conversion efficiency was 92.5% at an
output power of around 28 W, and an efficiency of 90%
was obtained at an output power of 60 W.
Power conversion efficiency at the light load washighest when type-A was used. However, withincreasing the output power, the power conversion
efficiency decreased significantly. Main reason of theefficiency decrease at large output power condition wasconsidered to be owing to the increase of eddy current
loss in the winding, where the leakage flux generated by
the difference of the magnetomotive forces between

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

INDEX

primary and secondary winding became large whenheavy load condition. On the other hand, in type-B andtype-C, the decrease in power conversion efficiency was
small even if the output power increases. These trends
were consistent with the results of power loss analysis
for the transformers.
5. Conclusion
In general, the low permeability magnetic core isnot suitable for transformer because of large leakage
inductance owing to the low coupling coefficient
between windings. Though the surface-oxidized
carbonyl-iron powder/epoxy composite magnetic core
has a low relative permeability of about 6, its magnetic
loss was small below 10 MHz in frequency.
This study was done to develop the surface-oxidized
carbonyl-iron powder/epoxy composite magnetic coretransformer with high coupling coefficient and to apply
it to the PFM controlled quasi-resonant mode
flyback-type DC-DC converter. The obtained results are
as follows.
(1) In order to obtain high coupling coefficient, theembedded winding structure was introduced to thecomposite magnetic core transformer design, and thethree kinds of winding layouts were investigated on thebases of 2D-FEA numerical simulation and experiments.
The adjacent layout of the primary and secondary
winding was best choice for highest coupling coefficient
over 0.99.
(2) The output voltage of the transformer, in which
the innermost of the primary winding and thesecondary winding were not adjacent to each other andhad low coupling coefficient, was higher than the value
calculated from the turn ratio at the no-load and lightload. This operation suggests that the operating point of
the transformer may unintentionally change from theoperating point designed based on the specifications of
the converter.
(3) The frequency at the minimum power loss
shifted to the low frequency side with increasing load
current in any transformer under investigation in thisstudy, and the transformers exhibited the minimum
power loss obtained at 500 kHz to 2 MHz according to
the load current intending for the PFM controlled
quasi-resonant type flyback converter in this study.
Since the switching frequency of the converter decreases
with increasing load current, there is a possibility thathigh efficiency can be maintained in the wide loadcondition when the composite magnetic coretransformer is applied to such converter.

(4) Application of the fabricated composite
magnetic core transformer with the best winding layout
described in above (1) to the 48 V input/12 V output
PFM controlled quasi-resonant mode flyback-type
DC-DC converter exhibited a maximum power
conversion efficiency of 92.5% at an output power ofaround 28 W, and 90% efficiency was obtained at an
output power of 60 W. The obtained efficiency
characteristics were consistent with the trend of power
loss analysis for the composite magnetic coretransformers.
Acknowledgements This study was supported in partby Grant-in-Aids for Japan Society for the Promotion of
Science (15H02233), New Energy and Industrial
Technology Development Organization (15101041-0),
Japan, and Japan Science and Technology Agency(Nagano Satellite Cluster, Kyoto Super Cluster
Program).
References
1) A. León-Masich, H. Valderrama-Blavi, J. M. Bosque-Moncusí,and L. Martínez-Salamero: IET Power Electronics, 8, 6,
pp.869-878 (2015).
2) H. Wang, A. M. H. Kwan, Q. Jiang, and K. J. Chen: IEEE
Trans. Electron Devices, 62, 4, pp.1143-1149 (2015).
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USA (2016).
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N. Zettsu, K. Teshima and M. Mizusaki: AIP Advances, 6,055932-1-8 (2016).
5) D. Shibamoto, A. Ueno, K. Sugimura, R. Hirayama, K. Sato,T. Sato, and M. Sonehara; The Papers of Joint Technical
Meetingon“Magnetics”
and“Linear
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6) K. Sato, K. Sugimura, T. Sato, M. Sonehara, and H.
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40, pp.35-38 (2016).
Received June 13, 2017; Revised July 10, 2017; Accepted July24, 2017

Journal of the Magnetics Society of Japan Vol.41, No.6, 2017

139

INDEX

Editorial Committee Members・Paper Committee Members
K. Kobayashi and T. Ono (Chairperson),(T. Kato, K. Koike and T. Taniyama (Secretary)
A. Fujita
H. GotoH. Hashino
S. Honda
S. Inui
S. Kasai
A. Kikitsu
H. Kikuchi
T. Kimura
T. Kubota
T. Nagahama
H. Naganuma
M. Naoe
M. Ohtake
N. PhamT. SatoT. SatoS. Seino
K. Sekiguchi
M. Sekino
Y. Shiratsuchi
M. Sonehara
T. Tanaka
K. Yamamoto
H. Yuasa
N. Adachi
N. Inaba
S. Kokado
T. Nakagawa
K. Tajima
K. Yamazaki

K. Bessho
S. Isogami
Y. Kota
H. Nakayama
M. Takezawa
S. Yoshimura

M. DoiK. KamataT. MakiT. Narita
T. Takura

T. DoiH. KatoE. Miyashita
D. Oyama
M. Tsunoda

Y. Kanai
K. Miura
Y. Sasayama
T. Shima

M. EndoK. KatoT. Morita
T. Saito
S. Yabukami

T. Hasegawa
T. KodaS. Muroga
R. Sugita
T. Yamamoto

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