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ISSN 2432-0250

Vol.42 No.2
JournalJournal of the Magnetics Society of Japan
Electronic Journal URL: https://www.jstage.jst.go.jp/browse/msjmag

2018

Magnetic Recording
Basic Study of Electric Field Induced Magnetization Reversal of Multiferroic (Bi1­xBax)FeO3 Thin Films
発行（公社）日本磁気学会 〠101‒0052 東京都千代田区神田小川町2丁目8番地
三井住友海上小川町ビル5階☎(03)5281-0106
印刷（株）国 際 文 献 社 〠169‒0075 東京都新宿区高田馬場3の8の8☎(03)3362-9741

at Room Temperature for Magnetic Recording Technology
S. Yoshimura, Y. Sugawara, G. Egawa, and H. Saito…011

Writing Field Amplitude in Heat-Assisted Magnetic Recording
T. Kobayashi, Y. Nakatani, K. Enomoto, and Y. Fujiwara…015

Hard and Soft Magnetic Materials
Relationship Between Bulk Coercivity and Coercivity of Surface Layer in Nd-Fe-B-Based Sintered Magnet
T. Maki, R. Ishii, M. Natsumeda, and T. Nishiuchi…020

Inﬂuence of Misorientation Angle Between Adjacent Grains on Magnetization Reversal in Nd-Fe-B-Based
Sintered Magnet
T. Maki, R. Uchikoshi, R. Ishii, M. Natsumeda, T. Nishiuchi, and M. Takezawa…024

Thin Films, Fine Particles, Multilayers, Superlattices
Structure Characterization of Fe, Co, and Ni Thin Films Epitaxially Grown on GaAs(111) Substrate
T. Soda, S. Minakawa, M. Futamoto, M. Ohtake, and N. Inaba…030

Spin Electronics
Theoretical Study on Magnetic Tunneling Junctions with Semiconductor Barriers CuInSe2 and CuGaSe2
Including a Detailed Analysis of Band-resolved Transmittances
K. Masuda and Y. Miura…037

Biomagnetism / Medical Applications
Evaluation of Dispersibility in Liquid and AC Magnetization Properties of Polyion Complex-coupled
Magnetic NanoParticles
G. Shi, S. B. Trisnanto, K. Nakai, S. Yusa, T. Yamada, S. Ota, and Y. Takemura…041

Power Supply for Medical Implants by Wiegand Pulse Generated from Magnetic Wire
K. Takahashi, A. Takebuchi, T. Yamada, and Y. Takemura…049

Dynamic Hysteresis Measurement of Magnetic Nanoparticles with Aligned Easy Axes
R. Takeda, S. Ota, T. Yamada, and Y. Takemura

広告◇ネオアーク株式会社◇東英工業株式会社◇計測エンジニアリングシステム株式会社

…055

42

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

2

2018

INDEX

Journal of the Magnetics Society of Japan
Vol. 42, No. 2
Electronic Journal URL: https://www.jstage.jst.go.jp/browse/msjmag

CONTENTS
Magnetic Recording
Basic Study of Electric Field Induced Magnetization Reversal of Multiferroic (Bi1−xBax)FeO3 Thin Films
at Room Temperature for Magnetic Recording Technology……………………………………………………… S. Yoshimura, Y. Sugawara, G. Egawa, and H. Saito

11

Writing Field Amplitude in Heat-Assisted Magnetic Recording………………………………………………… T. Kobayashi, Y. Nakatani, K. Enomoto, and Y. Fujiwara

15

Hard and Soft Magnetic Materials
Relationship Between Bulk Coercivity and Coercivity of Surface Layer in Nd-Fe-B-Based Sintered Magnet…
……………………………………………………………… T. Maki, R. Ishii, M. Natsumeda, and T. Nishiuchi

20

Influence of Misorientation Angle Between Adjacent Grains on Magnetization Reversal in Nd-Fe-B-Based
Sintered Magnet……………………… T. Maki, R. Uchikoshi, R. Ishii, M. Natsumeda, T. Nishiuchi, and M. Takezawa

24

Thin Films, Fine Particles, Multilayers, Superlattices
Structure Characterization of Fe, Co, and Ni Thin Films Epitaxially Grown on GaAs(111) Substrate……………………………………………T. Soda, S. Minakawa, M. Futamoto, M. Ohtake, and N. Inaba

30

Spin Electronics
Theoretical Study on Magnetic Tunneling Junctions with Semiconductor Barriers CuInSe2 and CuGaSe2
Including a Detailed Analysis of Band-resolved Transmittances……………………………………………………………………………………… K. Masuda and Y. Miura

37

Biomagnetism / Medical Applications
Evaluation of Dispersibility in Liquid and AC Magnetization Properties of Polyion Complex-coupled
Magnetic NanoParticles………………………G. Shi, S. B. Trisnanto, K. Nakai, S. Yusa, T. Yamada, S. Ota, and Y. Takemura

41Power Supply for Medical Implants by Wiegand Pulse Generated from Magnetic Wire……………………………………………… K. Takahashi, A. Takebuchi, T. Yamada, and Y. Takemura

49

Dynamic Hysteresis Measurement of Magnetic Nanoparticles with Aligned Easy Axes…………………………………………………………… R. Takeda, S. Ota, T. Yamada, and Y. Takemura

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

55

INDEX
J. Magn. Soc. Jpn., 42, 11-14 (2018)<LETTER>

Basic Study of Electric Field Induced Magnetization Reversal ofMultiferroic (Bi1-xBax)FeO3 Thin Films at Room Temperature
for Magnetic Recording Technology
S. Yoshimura1,2, Y. Sugawara1, G. Egawa1, and H. Saito1

Research Center for Engineering Science, Graduate School of Engineering & Resource Science, Akita University,

1

1-1 TegataGakuen-machi Akita, Akita 010-8502, Japan

PRESTO, Japan Science and Technology Agency (JST), Honcho Kawaguchi, Saitama 332-0012, Japan2(Bi1-xBax)FeO3 multiferroic thin films with ferromagnetism and ferroelectricity were fabricated and applied tocreate magnetic recordings using an electric field. The (001)-oriented (Bi1-xBax)FeO3 thin films of which electric
polarization direction is perpendicular to the film plane were fabricated onto a non-single-crystalline substrate witha Ta seedlayer / (111)-oriented Pt underlayer at a low substrate temperature of 500 ℃. A very high frequency
plasma irradiation was applied during sputtering deposition of (Bi1-xBax)FeO3 to accelerate the crystallization at thelow substrate temperature. The fabricated (Bi0.6Ba0.4)FeO3 film exhibited hysteresis curves indicating ferromagnetic
and ferroelectric behavior. The saturation magnetization of the film was approximately 60 emu/cm3 and thecoercivity was approximately 2.5 kOe, respectively. Magnetic Force Microscopy analysis of the (Bi0.6Ba0.4)FeO3 film
confirmed that the magnetization was reversed by applying only a local electric field. The multiferroic film describedhere is expected to be useful for electric field-driven magnetic devices.
Key words: Multiferroic thin film, control of magnetization direction, application of local electric field
1. Introduction
Magnetic recording using an electric field is apromising technology for future recording devices due to
its lower power consumption and ease of recording. For
example, magnetic field writing in hard disk drives(HDDs) will gradually become more difficult due toincreased magnetic field required as the switching fieldof the bits in order to prevent thermal agitation of therecording media. Recently, energy-assisted magnetic
recording technologies for next-generation HDDs, suchas thermally-assisted magnetic recording1) andmicrowave-assisted magnetic recording2), have been
widely studied in order to decrease the switching field.
However, energy-assisted magnetic recording devices
consume a large amount of power and the structure of
the magnetic field writing head is complex due to theadditional elements (a plasmon antenna forthermally-assisted magnetic recording or a spin-torqueoscillator for microwave-assisted magnetic recording).
On the other hand, electric field writing can be used for
HDDs with multiferroic recording media. The power
requirement for switching the magnetization direction
by this process is very low, the structure of the writing
head is simple (involving a needle-shaped conductive
element), and a very high electric field can be easily
applied to the multiferroic layer because it is thin.
Therefore, the control of magnetization using
electric fields is now being widely studied and several
techniques have been reported. (1) The magnetoelectriceffect
has
been
used
on
piezoelectricandmagnetostrictive laminate composites3) or multilayer
structures4). (2) An electric field can be used to control
the magnetic anisotropy of a thin metallic magnetic

layer in magnetic tunnel junctions with three layers: athin metallic magnetic layer, a thin insulating layer,
and a metallic magnetic layer5),6). (3) -Cr2O3 has bothdielectric and antiferromagnetic properties so a notableexchange interaction occurs between layers of -Cr2O3
and a ferromagnetic material at room temperature7) andmagnetization switching occurs in the ferromagnetic
layer when an electric field is applied due toantiferromagnetic coupling8). (4) BiFeO3 is amultiferroic material with a high ferroelectric Curie
temperature of 1120 K and a high antiferromagnetic
Neel temperature of 640 K. Thus, magnetization
switching occurs in ferromagnetic layers fabricated on
BiFeO3 layers when an electric field is applied at room
temperature due to antiferromagnetic coupling9),10),11).
(5) (Dy,Tb)FeO3 is a multiferroic material that exhibits
ferromagnetism and ferroelectricity and undergoes
magnetization switching when an electric field isapplied12).
However, there are problems associated with theapplication of these systems to magnetic devices. In
system (1), rapid operation of magnetization is difficult
to achieve due to the use of a mechanical strain on thepiezoelectric and magnetostrictive materials. Thus, theapplication of this system to magnetic devices that
require a high-speed operation and high durability will
be difficult. In system (2), unique and accurate electricfield frequencies and DC magnetic field of several
hundred oersteds have to be applied to reverse themagnetization of the thin metallic magnetic layer.
Moreover, effective reversal of the magnetization can
only be achieved with very thin metallic magnetic layers
due to the interface effect between the metallic
magnetic layer and the insulating layer. As a result, theapplication of this system to magnetic devices that

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

11

INDEXat room temperature, 300 ℃ , and 400 – 500 ℃ ,respectively. The film thicknesses and the deposition
temperatures for the Ta seedlayer and Pt underlayer
were optimized to obtain strong (111) orientation in thePt underlayer18). Very-high-frequency (VHF) (40.68
MHz) plasma irradiation with an electric power of 5 W
was applied during the radio-frequency (RF) (13.56
MHz) sputter deposition of the (Bi1-xBax)FeO3 film; VHF
plasma irradiation with a low power density has
previously
beenreportedtoacceleratethecrystallization of the film19). The crystallographic
orientations and the crystalline structures of thefabricated (Bi1-xBax)FeO3 films were analyzed by X-ray
diffraction (XRD). The surface morphology of the
BiFeO3 films was examined by SPM with a Si tip. The
magnetization curves of (Bi1-xBax)FeO3 films and theCurie temperature Tc of the (Bi0.6Ba0.4)FeO3 film were
measured using a vibrating sample magnetometer(VSM) with application of a magnetic field
perpendicular to the film surface. The ferroelectric
hysteresis loops of the (Bi1-xBax)FeO3 films were
measured using a ferroelectric tester. The local electricfield was applied to the (Bi0.6Ba0.4)FeO3 film using an
SPM with a conductive, magnetic Co-Zr-Nb tip. The
electric and magnetic domain structures of the(Bi0.6Ba0.4)FeO3 film were analyzed by electric force
microscopy (EFM) and magnetic force microscopy(MFM), respectively, with a conductive, magnetic
Co-Zr-Nb tip.
3. Results and Discussions
Fig. 1 shows the XRD profiles of Bi-Fe-O films
fabricated with various sputtering conditions on athermally oxidized Si substrate with a Ta/Pt layer (theXRD profile of Ta/Pt layer is also illustrated). The Pt
underlayer was found to have a strong (111) orientationand the full width at half maximum of the rocking curvefor the Pt(111) diffraction was 2.0 degrees. The Bi-Fe-O
film fabricated on the (111)-oriented Pt underlayer at a
substrate temperature of 400℃ was found to have a(001) orientation. Moreover, the Bi-Fe-O film fabricated(001)

require large magnetization switching and a low error
rate without a magnetic field will be difficult. In system(3), the antiferromagnetic Neel temperature of -Cr2O3
is not high (307 K) and not only an electric field but alsoa magnetic field are needed for magnetization switching.
Further, the magnetization reversal can only beachieved with a very thin metallic ferromagnetic layer
due to interface effect between the -Cr2O3 andferromagnetic layer. Therefore, the application of thissystem to magnetic devices that must withstand hightemperatures and require large magnetization
switching without a magnetic field will be difficult. In
system (4), the exchange interaction between the
BiFeO3 and ferromagnetic layer is small so theapplication of this system to magnetic devices thatencounter large external fields and thus require large
magnetization switching will be difficult. In system (5),the magnetic Curie temperature of almost multiferroic
materials are very low, making the application of thissystem to practical magnetic devices difficult.
Therefore, a new system that can undergo magnetic
switching in a single phase with ferromagnetism andferroelectricity and with a high Curie temperature isneeded for the magnetic recording via electric fieldwriting. Several suitable multiferroic materials withferromagnetism
andferroelectricity
atroom
temperature,suchas(Bi1-xBax)FeO313)
andBi(Fe1-xCox)O314),15), have been reported. In this study,
we propose a novel magnetic recording technology basedon electric field writing with a (Bi1-xBax)FeO3 thin filmthat can be used to develop new magnetic recording
devices with low power consumption In HDDs withmultiferroic recording media, the proposed system offers
low power consumption for writing, a simple structure
for the writing head, and a very high electric field forwriting. The standard fabrication method for
multiferroic films is pulsed-laser deposition onto
single-crystalline substrates at a high temperature ofmore than 600 ℃ 16,17). However, this method is notsuitable for device manufacturing. In this study, wefabricated (001)-oriented (Bi1-xBax)FeO3 thin films onto
non-single-crystalline substrates at a relatively low
substrate temperature of less than 500℃ by magnetron
sputtering. The obtained thin film had electric
polarization direction perpendicular to the film plane
and was suitable for device applications. Moreover, wedemonstrated the control of the magnetization direction
on the micrometer scale by applying a local electric field
on a scanning probe microscope (SPM) with a Co-Zr-Nb
coated conductive tip.
2. Experimental ProcedureThree layers were deposited onto a thermally
oxidized Si wafer using a ultra-high-vacuum (UHV)
sputtering system: 5 nm Ta, 100 nm Pt, and 100 nm(Bi1-xBax)FeO3, in which the Ba concentration, x, wasvaried to 0, 0.1, 0.2, and 0.4. The layers were deposited

12

250 nmFig. 1 XRD profiles of Ta seedlayer /Pt underlayer
and Bi-Fe-O films on Ta/Pt layer fabricated withvarious sputtering conditions.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX

on a (111)-oriented Pt underlayer at a substratetemperature of 400℃ with a VHF plasma irradiation
power of 5 W was found to exhibit a strong (001)diffraction peak. The atomic force microscopy (AFM)
images of the Bi-Fe-O films are also shown in this figure.
The grain size of the Bi-Fe-O film fabricated in thepresence of VHF plasma irradiation was found to belarger than that of the Bi-Fe-O film fabricated in theabsence of VHF plasma irradiation. These results
indicate that the use of a highly-oriented (111) Pt
underlayer and VHF plasma irradiation are the key toobtaining strong (001) orientation and crystal grain
growth in BiFeO3 thin films at a relatively low
substrate temperature.
Fig. 2 depicts the dependence of the saturation
magnetization Ms and electric polarization at 750 kV/cm
P750kV/cm on the Ba concentration in the (001)-oriented(Bi1-xBax)FeO3 film fabricated on an amorphous Ta
seedlayer on a highly (111)-oriented Pt underlayer inthe presence of VHF plasma irradiation. The maximum
electric field used to measure the electric polarization
was 750 kV/cm due to the insulating destruction. The
magnetization hysteresis curve (major loop) and theferroelectric hysteresis curve (minor loop) of the(Bi0.6Ba0.4)FeO3 (x=0.4) film are also illustrated in thisfigure. As the Ba concentration is increased, Ms
increased. Based on the clear hysteresis in themagnetization curve and ferroelectric loops of(Bi0.6Ba0.4)FeO3 film, the Ms was found to be about 60
emu/cm3; this is approximately the same as that of(Bi1-xBax)FeO3 films fabricated by pulsed-laser
deposition on single-crystalline substrates at atemperature above 600 ℃ . The coercivity Hc of the(Bi0.6Ba0.4)FeO3 film was about 2.5 kOe and thesquareness ratio Mr/Ms was 0.6. Based on the minor
loop of the magnetic hysteresis loop (not shown), themagnetization reversal is thought to be dominated by anucleation process. The relationship between themagnetic and electric properties will be discussed in a

50
0

-50

-100
-10

-5

0

H (kOe)

5

40

10

2

P (C/cm )

10

20

5
0
-5

-10
-1000 -500

0500

E (kV/cm)

002P750 kV/cm (C/cm )

40

60

100

M (emu/cm3)3

Ms (emu/cm )60

0.1 0.2 0.3 0.4

x [Ba/(Bi+Ba)]

1000

20

0
0.5

Fig. 2 Dependence of saturation magnetization andelectric polarization at 750 kV/cm on Ba
concentration for (Bi1-xBax)FeO3 films.

Fig. 3 Dependence of saturation magnetization on
measuring temperature and the magnetization
curve at the measuring temperature of 250 ℃ for(Bi0.6Ba0.4)FeO3 film.
future study because only the minor loop of the electric
property was measured here.
Fig. 3 shows the temperature dependence of Ms andthe magnetization curve obtained at a temperature of
250℃ for the (Bi0.6Ba0.4)FeO3 film. A clear hysteresis
was observed and Tc was estimated to be approximately
400 ℃ , which is useful for application to practicalmagnetic devices.
Figs. 4 (a), (b), and (c) show topographic, EFM, and
MFM images, respectively, of the (Bi0.6Ba0.4)FeO3 film
prior to the application of a DC voltage to the film. The
MFM images were captured with a Co-Zr-Nb magnetic
tip that was magnetized to saturation at the end of thetip such that the magnetization direction would be
perpendicular to the sample surface and the direction of
the detected magnetic field from the sample would also
be vertical to the sample surface. For the EFM imaging,
the Pt underlayer was grounded and a DC voltage of-1.0 V (-100 kV/cm) was applied to the Co-Zr-Nb
conductive tip. Because this DC electric field wassmaller than the coercive electric field of the(Bi0.6Ba0.4)FeO3 film, as depicted in Fig. 2, polarization
reversal did not occur during this imaging. The MFM
image shows the demagnetized domain structure, andthe EFM image shows similar domain structure.
While writing on the (Bi0.6Ba0.4)FeO3 film using a
local electric field, the surrounding area (3 m squareregion) was scanned by the Co-Zr-Nb conductive tip on
which a DC voltage of +6.5 V was applied and the centerarea (1 m square region) was scanned by the Co-Zr-Nb
conductive tip on which a DC voltage of -6.5 V wasapplied. The DC voltage of ±6.5 V (±650 kV/cm) for the
local electric field used for writing was larger than thecoercive electric field of the (Bi0.6Ba0.4)FeO3 film so it
was sufficient to reverse the electric polarization
direction of the film. Figs. 4 (d), (e), and (f) show
topographic, EFM, and MFM images of the(Bi0.6Ba0.4)FeO3 film after a DC voltage of ±6.5 V wasapplied. In the EFM image obtained with a with the tip
voltage of -1.0 V, an attractive force between tip andsample was observed in the center of the ferroelectric

Journal of the Magnetics Society of Japan Vol.42, No.2, 201813

INDEX

(a)

(b)

(c)

(e)

(f)

2 m

(d)

Fig. 4 (a) topographic, (b) EFM, and (c) MFM imagesof (Bi0.6Ba0.4)FeO3 film before applying DC voltage,and (d) topographic, (e) EFM, and (f) MFM images ofthat film after applying DC voltage.
domain structure, which indicates a positive electriccharge in this area. Conversely, in the surrounding area,
a repulsive force between tip and sample was seen in
the ferroelectric domain structure, indicating a negativeelectric charge. In an EFM image obtained with a tip
voltage of +1.0 V (not shown here), the attractive force
and repulsive force in the center area and thesurrounding area were reversed. These images showthat clear micrometer-scale ferroelectric domains wereformed upon application of the local electric field. In the
MFM image obtained with the tip magnetized at theend, a repulsive force between the tip and the sample
was observed in the center of the ferromagnetic domain
structure, indicating that the magnetization direction inthis area was upward from the film surface. However, inthe surrounding area, an attractive force between the
tip and the sample was observed in the ferromagnetic
domain structure, indicating that the magnetization
direction was downward toward film surface. In an
MFM image obtained with a tip voltage of 0 V, therelationship between the attraction and repulsion in thecenter area and the surrounding area did not change
regardless of both before and after EFM measurementswith tip voltages of -1.0 V and +1.0 V. Therefore, it canbe concluded that micrometer-scale magnetization
reversal was achieved by applying a local electric field
and that the directions of P and M are parallel. Thus,
the proposed multiferroic films are expected to be usefulin novel magnetic recording devices driven by electricfield writing.
4. Conclusion
In this study, we fabricated (Bi1-xBax)FeO3
multiferroic thin films and demonstrated local
magnetization reversal by the application of an electricfield. The Ta seedlayer on the (111)-oriented Pt
underlayer
enabledsuccessful
fabricationof(001)-oriented (Bi1-xBax)FeO3 thin films with electric
polarization perpendicular to the film plane on a
non-single-crystalline substrate. Further, sputtering
deposition with VHF plasma irradiation was found to14

facilitate the crystallization of the (Bi1-xBax)FeO3 thinfilms at a relatively low substrate temperature(compared to the standard method) of 500 ℃ . The(Bi0.6Ba0.4)FeO3 film, in which the Ba concentration was
optimized, was found to have a high saturation
magnetization of about 60 emu/cm3, a high coercivity of
about 2.5 kOe, and a high Curie temperature of about
400℃. MFM images revealed that the magnetization
direction in the (Bi0.6Ba0.4)FeO3 film could be controlled
on a micrometer scale by simply applying an electricfield. Based on these results, the proposed multiferroic
film is expected to be useful for applications in electricfield-driven magnetic devices.
Acknowledgements This work was partially supported
by Iketani Science and Technology Foundation, Yazaki
Memorial Foundation for Science and Technology, and
JST/PRESTO (No. JPMJPR152C, ID: 15655293).
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Received Aug. 08, 2017; Accepted Jan. 18, 2018

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J. Magn. Soc. Jpn., 42, 15-19 (2018)<Paper>                    
          
             

              
                                                                                    
                                                   
              
     
           
      
                           
                  
    
                                     
                   
             
                               
                                                           
                 
                                                             

Journal of the Magnetics Society of Japan Vol.42, No.2, 201815

INDEX                   
                                   



                                     
                                    
                                                                                
                                                                                                   
 

 
                               
     
                                           16          
                          
       
                                                                                                                                              


           
           
                   
                                                                                                                      



                                                                             
                                    Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX                
                            








      
                                    
                                                          
                                        
                         




 



                                                       

                       
        
           
         
                                                                                                    

Journal of the Magnetics Society of Japan Vol.42, No.2, 201817

INDEX                           
       
                          
                                                                    
                                                          
 
        
                                                        

18                      
                                                     
         
       
                                                                        
                                                                                                                                                 

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX                                                                  
                                                                                     
                                                                                                           
     
              
                                   
                     
        
                



            

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018  

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J. Magn. Soc. Jpn., 42, 20-23 (2018)<LETTER>

Relationship between bulk coercivity and coercivity of surface layer in
Nd-Fe-B-based sintered magnet
T. Maki, R. Ishii, M. Natsumeda, and T. Nishiuchi

Magnetic Materials Research Laboratory, Magnetic Materials Company, Hitachi Metals, Ltd., 2-15-17 Egawa, Shimamoto-cho,Mishima-gun, Osaka 618-0013, JapanThe bulk coercivity and coercivity of the polished surface layer of a Nd-Fe-B-based sintered magnet were
measured systematically in order to discuss the effectiveness of observing the magnetic domain on the surface of themagnet. It was revealed that the bulk coercivity and coercivity of the polished surface layer were in a proportional
relationship even though the coercivity of the surface layer was much smaller than the bulk coercivity. This suggeststhat the magnetization reversal that occurred in the surface layer has a mechanism similar to that inside of themagnet. The relationship between the bulk coercivity and coercivity of the surface layer was not proportional when
the degree of alignment decreased, which can be explained by the increase in the local demagnetizing field on thesurface of the magnet.
Key words: Nd-Fe-B magnet, coercivity, surface layer, degree of alignment, demagnetizing field
1. Introduction
The coercivity of a Nd-Fe-B-based magnet is an
important property, especially for high temperature use,and partially substituting Nd with heavy rare earth
elements such as Dy and Tb is an effective way to
improve the coercivity. However, decrease in the usageof Dy and Tb has been required recently because of
their limited production. The coercivity of a Nd-Fe-B-based magnet is greatly changed not only by theaddition of Dy and Tb but also by the microstructure. To
understand directly the relationship between the
microstructure and magnetic domain structure in themagnet, several methods for evaluating the magnetic
domain structure on the surface of the magnet arestudied, such as the magneto-optical Kerr microscope1, 2),magnetic force microscope (MFM)3, 4) and spin-polarizedscanning electron microscope (spin-SEM)5, 6). However,on a polished surface of a Nd-Fe-B-based magnet,
magnetization reversal progresses with a clearlysmaller magnetic field than the bulk coercivity. For thisreason, a method for suppressing the reduction insurface coercivity was studied recently by using
fractured surface7). If we could identify the similarities
and differences in the magnetic reversal process
between the surface and inside of a magnet, it would bevery useful for discussing the coercivity. In this study,
to clarify the correlation between the surface and insideof the Nd-Fe-B-based sintered magnet, the bulk
coercivity and coercivity of the polished surface layer
were measured, and the influence of the Dy content,
annealing condition, and degree of alignment on these
coercivities was investigated.

(30.2-x)Nd-xDy-67.6Fe-1.0B-0.9Co-0.1Al-0.1Cu-0.1Ga(mass%) (x = 0, 1.0, 2.0, 5.0) were prepared by using an
ordinary process except for pressing, which applied a
different magnitude of magnetic field. Degree ofalignment  was defined as the following formula8, 9),
 = Br / Js
(1)
where Br and J s are the remanence and saturation
magnetization of each sample. For example, the degreeof alignment of HA and MA magnets for x = 0 was  =0.97 and 0.90, respectively. The annealing temperature
after sintering was changed from 460 to 540°C. Eachsample was set to 7 × 7 × 7 mm under the same
grinding condition by using a surface grinder and thedemagnetization curve was measured by using a BH
tracer. The coercivity of the surface layer was obtained
from the differential value of the step appearing in thesecond quadrant of the demagnetization curve, as
shown in Fig. 1. It is supposed that this step in the
demagnetization curve is derived from a-b planes:

2. Experimental method
Highly aligned (HA) and moderately aligned (MA)sinteredmagnetsthathadcompositionsof20

Fig. 1 Illustration of bulk coercivity and coercivity ofsurface layer.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

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Fig. 2 Bulk coercivity and coercivity of surface layerof highly aligned (HA) magnets when changing
annealing temperature. Solid lines are bulk
coercivity and dashed lines are coercivity of surface
layer.

Fig. 3 Relationships between bulk coercivity andcoercivity of surface layer of highly aligned (HA)
and moderately aligned (MA) magnets.

Fig. 4 (a) Model of Nd-Fe-B sintered magnet for calculation with 3D-FEM and (b) distribution of permeancecoefficient Pc when YX and YZ were changed from 0 to 45°.

surface planes located parallel to the orientation
direction, which demagnetized independently of theinterior of the magnet10). When the coercivity was 1600
kA/m or higher, the bulk coercivity was measured byusing a pulsed BH tracer. To understand the influenceof the demagnetizing field on the magnet surface, athree-dimensional finite element method (3D-FEM) wasused to simulate the distribution of the permeancecoefficient.

3. Result and discussion
Figure 2 shows the bulk coercivity and coercivityof the surface layer of the HA magnets when changing
the Dy content and annealing temperature. The bulk
coercivity increased as the Dy content increased andreached the maximum value when the annealing
temperature was 500°C or 520°C in each composition. It
was found that the coercivity of the surface layer also
varied depending on these conditions with the sametendency as the bulk coercivity, even though the

Journal of the Magnetics Society of Japan Vol.42, No.2, 201821

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Fig. 5 Average permeance coefficient Pc on surfaceof 1 × 1 × 1 mm particle when YX and YZ werechanged from 0 to 45°.
coercivity of the surface layer was much smaller thanthe bulk coercivity. Figure 3 shows the relationshipbetween the bulk coercivity and the coercivity of thesurface layer when the degree of alignment, the Dycontent, and the annealing temperature were changed.
The coercivity of the surface layer changed inproportion to the change in the bulk coercivity
depending on the Dy content and the annealing
temperature under the same degree of alignment. Thissuggests that the coercivity of the surface layer also
depends on the magnetocrystalline anisotropy of the(Nd,Dy)2Fe14B main phase and the microstructure,especially around the grain boundary. This resultmeans that the magnetization reversal occurring in thepolished surface layer has a similar mechanism to theinside of the magnet. Comparing each HA magnet andMA magnet under the same composition and annealing
condition, the coercivity of the surface layer had about
the same value, even though the bulk coercivity in the
MA magnet was larger than that in the HA magnet.
When the degree of alignment decreases, it is expectedthat the number of grains that incline its easy axis ofmagnetization in the out-of-plane direction increases on
the magnet surface. This may lead to an increase in thelocal demagnetizing field and induce magnetization
reversal in the magnet surface.
To clarify the change in the demagnetizing field
due to the inclination of the easy axis of magnetizationon the magnet surface, we assumed the simple model
shown in Fig. 4(a) and calculated the demagnetizing
field by 3D-FEM. First, a 3 × 3 × 3 mm of cubic Nd-Fe-B
magnet that was fully magnetized in the Y direction
was created. The mesh size had a 0.1 mm pitch, and thetotal number of units was 27000. Each unit was affectedby the static magnetic field from the surrounding units,
but there was no exchange interaction among units.

22

Next, we focused on a 1 × 1 × 1 mm of cubic Nd-Fe-B
particle that existed in one plane that was parallel tothe easy magnetization direction in the 3 × 3 × 3 mm of
cubic Nd-Fe-B magnet. The direction of themagnetization of this particle changed independently in
the in-plane direction (Y to X) and out-of-plane direction(Y to Z) at angles of YX and YZ, respectively. Figure4(b) shows distributions of permeance coefficient Pc
when YX and YZ were changed from 0 to 45°. When YX= YZ = 0, the distribution of Pc in a surface area of 1
mm square, indicated by a square line in Fig. 4(b), wasrelatively even. When YZ was increased, thedistribution of Pc in the area changed greatly comparedwith the case in which YX was increased. Figure 5
shows the change in average Pc in the area with respect
to YX and YZ. The average Pc decreased as YZ
increased, whereas it hardly changed when YX varied.
From the relation of N = 1 / (1 + Pc), where N is ademagnetizing factor, it is suggested that the local
demagnetizing field becomes large when the easy axisof magnetization is tilted in the out-of-plane direction
on the magnet surface. Therefore, the increase in thelocal demagnetizing field promotes magnetization
reversal and may decrease the coercivity on the magnet
surface, which is the reason for the non-proportional
relationship between the bulk coercivity and the surface
layer coercivity when comparing HA and MA magnets.
4. Conclusion
To clarify the correlation between the surface andthe inside of a Nd-Fe-B-based sintered magnet, the bulk
coercivity and coercivity of the polished surface layer
were measured systematically. It was revealed that thebulk coercivity and coercivity of the surface layer werein a proportional relationship, which suggests that themagnetization reversal occurring in the surface layer
has a mechanism similar to the inside of the magnet.
The relationship between the bulk coercivity andcoercivity of the surface layer was not proportional
when the degree of alignment decreased, which can be
explained as the increase in the local demagnetizing
field due to the low degree of alignment promoting themagnetization reversal and decreasing the coercivity on
the magnet surface.
References
1) D. Li and K. J. Strnat: J. Appl. Phys., 57, 4143 (1985).
2) A. Fukuno and R. C. O’Handley: J. Appl. Phys., 65, 4959
(1989).
3) T. Yamaoka, H. Tsujikawa, R. Hirose, A. Ito, H. Kawamura,and T. Sakon: J. Magn. Soc. Jpn., 35, 60 (2011) [inJapanese].
4) J. Thielsch, T.G. Woodcock, L. Schultz, and O. Gutfleisch: J.
Appl. Phys., 111, 103901 (2012).
5) T. Kohashi, K. Motai, T. Nishiuchi, T. Maki, and S.
Hirosawa: J. Magn. Soc. Jpn., 33, 374 (2009).
6) H. Suzuki, Y. Satsu, T. Kohashi, K. Motai, and M. Komuro:
J. Appl. Phys., 109, 07A746 (2011).

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

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7) K. Toyoki, Y. Kotani, Y. Senba, D. Billington, H. Okazaki, A.
Yasui, W. Ueno, H. Ohashi, Y. Shiratsuchi, S. Hirosawa, K.
Hono, and T. Nakamura: Proc. 24th Int. Workshop on
Rare-Earth and Future Permanent Magnets and their
Applications, Darmstadt, 2016, 129 (2016).
8) D. Harimoto, Y. Matsuura, and S. Hosokawa: J. Jpn. Soc.
Powder Powder Metallurgy, 53, 282 (2006) [in Japanese].
9) Y. Matsuura, J. Hoshijima, and R. Ishii: J. Magn. Magn.

Mater., 336, 88 (2013).
10) K. Kobayashi, M. Nakamura, and K. Urushibara: J. Appl.
Phys., 117, 173909 (2015).
Received Oct. 17, 2017; Revised Nov. 27, 2017; Accepted Dec.
29, 2017

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23

INDEX
J. Magn. Soc. Jpn., 42, 24-29 (2018)<Paper>Influence of misorientation angle between adjacent grains onmagnetization reversal in Nd-Fe-B-based sintered magnet
T. Maki, R. Uchikoshi*, R. Ishii, M. Natsumeda, T. Nishiuchi, and M. Takezawa*

Magnetic Materials Research Laboratory, Magnetic Materials Company, Hitachi Metals, Ltd., 2-15-17 Egawa, Shimamoto-cho,Mishima-gun, Osaka 618-0013, Japan*Department of Applied Science for Integrated System Engineering, Kyushu Institute of Technology, 1-1 Sensui-cho, Tobata-ku,

Kitakyushu-shi, Fukuoka 804-8550, JapanTo clarify the difference between the degree of alignment dependence of coercivity and the angular dependenceof coercivity, the crystal orientation distribution and demagnetization curve of Nd-Fe-B-based sintered magnets
with different degrees of alignment were compared. It is suggested that the increase in coercivity due to a low
degree of alignment cannot be explained only by the angular dependence of coercivity. A crystal orientation analysis
and in-situ observation of magnetic domains that were performed in the same area clarified that the ratio of thegrain boundary where the magnetization reversal stopped became larger when the misorientation angle between
adjacent grains became larger. This suggests that a grain boundary having a larger misorientation angle is one of
the factors that suppresses magnetization reversal.
Key words: Nd-Fe-B magnet, coercivity, degree of alignment, Kerr microscope, magnetic domain
1. Introduction
Nd-Fe-B-based magnets used in motors for electricvehicles (EV) and hybrid electric vehicles (HEV) are
required to have a high coercivity for maintaining
thermal stability. Generally, heavy rare earth elements
such as Dy and Tb are added to increase themagnetocrystalline anisotropy of the Nd2Fe14B
compound to obtain a high coercivity. However, theseelements are recognized as materials with a highsupply risk, and there is a strong demand for Nd-Fe-B-based magnets to achieve a high coercivity with a lower
amount of heavy rare earth elements.
The coercivity of the Nd-Fe-B-based magnet isgreatly changed not only by the addition of heavy rareearth elements but also by the microstructure. It waspointed out that Nd-rich phases existing at the grainboundary play an important role in coercivity1)-3), and it
was observed that a Nd-rich phase inhibited
propagating magnetic domains4,5). The coercivity of the
Nd-Fe-B-based sintered magnet also strongly depends
on the degree of alignment of the Nd2Fe14B phase, and
it is experientially known that the coercivity increases
as the degree of alignment decreases6,7). It was also
reported that the coercivity increases when the angle
between the easy magnetization direction of the
Nd-Fe-B sintered magnet and the direction of theapplied magnetic field becomes large, which is referred
to as the angular dependence of coercivity8). Thisdependency is explained by the magnetic domain wallmotion model, which shows that the coercivity increases
at a rate of 1/cos. Attempts have been made to explain
the dependence of the coercivity on the degree ofalignment by combining the angular dependence andorientation distribution7), but it is not fully understood.
When the degree of alignment decreases, it is easy to

24

expect that the distribution of the orientation angle 
between the applied magnetic field H and the easy axisof magnetization of each Nd2Fe14B grain broadens asshown in Fig. 1. However the distribution of themisorientation angle  between the easy axis ofmagnetization of adjacent grains will change at thesame time. There is a possibility that this change in the
distribution of  may be related to magnetic domain
wall motion and coercivity, but no example has beenreported experimentally.
To understand directly the relationship between
the microstructure and magnetic domain structure in
the Nd-Fe-B-based magnet, various methods forevaluating the magnetic domain structure on thesurface of the magnet are studied, such as themagneto-opticalKerrmicroscope9),10),
Lorentztransmission electron microscope (Lorentz-TEM)4),5),
electronic holography 11) , magnetic force microscope

Fig. 1 Schematic images of orientation angle θ
between easy axis of magnetization of each grain
and applied magnetic field H and misorientation
angle between easy axis of adjacent grains.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX(MFM)12),13), and spin-polarized scanning electron
microscope (spin-SEM)14),15). Among them, themagneto-optical Kerr microscope can perform in-situ
observation in a high magnetic field, so it is an effective
method for directly observing the magnetization
behavior of the Nd-Fe-B magnet16). Takezawa et al.
reported the in-situ magnetic domain observation of a
Nd-Fe-B-based sintered magnet in a magnetic field.
Several grains were simultaneously reversed along thedirection of the easy axis of magnetization, and when
the applied magnetic field increased, the magnetization
reversal propagated17),18). It is important to observe awide area to understand the propagation ofmagnetization reversal across multiple grain groups,
and the magneto-optical Kerr microscope is also
suitable in this respect.
In this study, first, to clarify the difference between
the degree of alignment dependence of coercivity andthe angular dependence of coercivity, we investigatedthe  distribution and  distribution of Nd-Fe-B-based
sintered magnets with different degrees of alignment.
The coercivity was measured by changing the angle
between the magnetic field and sample. Subsequently,
the magnetization process was analyzed by using recoil
curves20). Next, to experimentally evaluate therelationshipbetweenthedifferenceinthemisorientation angle between adjacent grains and themagnetic domain structure in the Nd-Fe-B-based
sintered magnet, a scanning electron microscope with
an electron backscatter diffraction (SEM/EBSD)
analysis and the in-situ magnetic domain observationin the magnetic field by using a magneto-optical Kerr
microscope were conducted in the same observed area.
2. Experimental method
Highly aligned (HA) and moderately aligned (MA)sintered magnets that had a composition of
30.2Nd-67.6Fe-1.0B-0.9Co-0.1Al-0.1Cu-0.1Ga (mass%)
were prepared by using an ordinary process except forpressing, which applied a different magnitude ofmagnetic field. The remanence Br and the coercivity HcJof the HA and MA magnets were Br = 1.44 T, HcJ = 1038
kA/m, Br = 1.33 T, and HcJ = 1126 kA/m, respectively.
The distribution of the degree of grain alignment was
analyzed by observing the polished surface plane thatwas parallel to the easy magnetization direction of themagnet with SEM/EBSD. The sintered magnets wereshaped into spheres for magnetic measurement. Aftermagnetization with a pulse magnetic field of 5.6 MA/m,the demagnetization curves were measured in the rangeof 1.6 to −1.6 MA/m with a vibrating sample
magnetometer (VSM). In the case of changing the angleof the applied magnetic field with respect to the easymagnetization
directionofthesample,theelectromagnet of the VSM was rotated at a 1° pitch withrespect to the sample. Each demagnetization curve wascorrected by setting the demagnetizing factor N of thesphere to N = 0.33. The recoil curve measurement was
repeated by increasing the demagnetization field by 40
kA/m in the second and third quadrants of thedemagnetization curve and returning to 0 kA/m, as
described in the previous paper20).
The sintered magnets used for magnetic domain
observations had two compositions (30.2-x)Nd-xDy-67.6Fe-1.0B-0.9Co-0.1Al-0.1Cu-0.1Ga (x = 0, 5.0) withdifferent Dy content, and the MA magnets were used.
The magnetic property of a Dy doped (x = 5.0) magnet
was Br = 1.24 T and HcJ = 1958 kA/m. Each sample wascut into 3 × 3 × 3 mm, and the plane that was polished
was parallel to the easy magnetization direction for an
observation surface. First, SEM/EBSD analyses were
performed at the center of the polished plane, and themagnetic domain was then observed with themagneto-optical Kerr microscope. An SiO film wasdeposited on the surface of each sample by vacuum
evaporation for antireflection17). Each sample wasmagnetized in the orientation direction with a pulsed
magnetic field of 4.0 MA/m and then set in themagneto-optical Kerr microscope. A DC magnetic field
was applied to the samples in the range of 1.6 to −1.6
MA/m. Magnetization reversal was detected from thechange in contrast in an observation image, and thereversed area was extracted by image processing18). To
obtain the effective magnetic field in the observationarea, the demagnetizing field at the center of the planethat is parallel to the easy magnetization direction of
the cubic Nd-Fe-B magnet was calculated by using athree-dimensional finite element method (3D-FEM).
The effective magnetic field Heff was obtained from therelationship of Heff = Hex − Hd, where Hex and Hd are theexternal magnetic field and the demagnetizing fieldfrom a sample, respectively. For example, in the case of
x = 0, Heff = −250 kA/m when Hex = 0 kA/m.
3. Results and discussion
3.1 Distribution of  and 
Figure 2 shows inverse pole figure maps of the MA
and HA magnets with the horizontal direction as anorientation direction. Compared with the HA magnet,
the MA magnet had a large dispersion of crystal
orientation. The orientation angle  between the
direction of the applied magnetic field H and the
direction of the easy axis of magnetization in each grain
was calculated from the direction cosine of the RD
direction and <001> direction of each grain. The
misorientation angle  between the easy axis ofadjacent grains that was defined regardless of therelative position of adjacent grains was calculated fromthe direction cosine of <001> directions of adjacentgrains. Figures 3(a)-(d) show the distribution of  and in each sample. Compared with the HA magnet, both ofthe distributions of  and  in the MA magnet werebroad, and the frequencies on the higher angle werelarge. It was found that not only the orientation angles

Journal of the Magnetics Society of Japan Vol.42, No.2, 201825

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Fig. 2 Inverse pole figure maps of Nd2Fe14B phases of (a) highly aligned (HA) magnet and(b) moderately aligned (MA) magnet.

Fig. 3 (a) to (d) Distributions of  and  of HA and MA magnets, (e) distribution of ' of MA magnet
calculated from direction tilted 23° from orientation direction, and (f) schematic image of  and ’.
but also the misorientation angles between adjacentgrains had a wide distribution when the degree ofalignment decreased.
3.2 Relationships between distribution of ,  andcoercivity
Figure 4 shows demagnetization curves of the HA
magnet, MA magnet, and HA magnets that wereinclined at 23° and 48° with respect to the direction of
the applied magnetic field. When the HA magnet was
tilted at an angle of 23°, the coercivity was clearly lower
than that of the MA magnet, whereas the squareness
Jr/Js of each sample was almost the same. When trying
to obtain the same coercivity as the MA magnet, the HA
magnet must be tilted up to 48° with respect to the26

applied magnetic field. Figure 3(e) shows the
distribution of ’, which was obtained by calculating theangle ’ between the direction of the easy axis ofmagnetization in each grain and the direction that was
intentionally deviated 23° from the RD direction. The
distribution of ’ was similar to the distribution of  ofthe MA magnet. Therefore, the MA magnet and HA
magnet tilted at 23° had approximately the sameorientation distribution as the direction of the appliedfield, which was consistent with the fact that Jr/Js was
the same in the demagnetization curve. These results
suggest that the large increase in coercivity due to thelow degree of alignment cannot be explained only by the
angular dependence of coercivity.
Figure 5(a) shows whole recoil curves of the MA

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Fig. 4 Demagnetization curves measured for HA
magnet, MA magnet, and HA magnet inclined at 23°
and 48° with respect to magnetic field direction(spherical sample, demagnetizing factor N = 0.33).

Fig. 6 Reversed region in each effective magnetic
field Heff extracted from Kerr microscope image indemagnetization process of (a) Dy undoped (x = 0)
and (b) Dy doped (x = 5.0) MA magnets.

Fig. 5 (a) Recoil curves of MA magnet and HA
magnet measured by tilting 23° with respect tomagnetic field direction and (b) comparison of two
recoil curves returning to near origin position.

Fig. 7 Changes in ratio of unreversed area withregards to Heff in demagnetization process of each offour observed fields for x = 0 and x = 5.0 samples
extracted from Kerr microscope image.

magnet and HA magnet tilted at 23°, and Fig. 5(b)shows recoil curves of both samples that returned tonear the origin. The curvature of the recoil curve of theMA magnet was smaller than that of the tilted HAmagnet. In the previous paper20), it was pointed out thatthe recoil curve of a Nd-Fe-B sintered magnet that has

an average grain size of several m contains domain
wall motion in its magnetization process. The small
curvature of the recoil curve suggests that the magnetic
domain wall motion in the MA magnet was small. As
shown in Fig. 3(d), the MA magnet had a large
di s tri b u ti o n o f , wh i ch m a y b e re l a te d to th e

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

27

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Fig. 8 Ratio of GBa and GBb with respect to misorientation angle of adjacent grains in (a) Dy
undoped magnet (x = 0) and (b) Dy doped magnet (x = 5.0).
suppression of the domain wall motion.
3.3
Relationshipbetweenpropagationofmagnetization reversal and misorientation angle Figure 6 shows magnetically inversed areas in eachmagnetic field Heff extracted from the magnetic domain
observation results for the demagnetization process of
the Dy undoped (x = 0) and Dy doped (x = 5.0) MA
magnets. These samples had almost the same
distribution of  and . Grain boundaries in the sample
are indicated by white dotted lines to facilitate
recognition. The unreversed areas were considered to beundetected areas because of a significantly weak change
in contrast and were excluded from the evaluation.
Several grains were reversed simultaneously along themagnetic easy axis direction, which is similar to thecase of highly aligned Nd-Fe-B sintered magnets17). To
estimate the repeatability of the demagnetization
process in the same Heff, we prepared four samples forboth x = 0 and x = 5.0 under the same condition andobserved these samples at the center of each polishedplane with the same method discussed above. Figure 7
shows changes in the ratio of the unreversed area withrespect to Heff in the demagnetization process of each of
the four observed fields with x = 0 and x = 5.0. In thesurface layer for x = 5.0, which had a large amount of
Dy, the magnetization reversal propagated at a higher
magnetic field. It is thought that the magnetization
reversal on the surface depended on the magnitude ofmagnetocrystalline anisotropy of the (Nd,Dy)2Fe14B
phase, even though the coercivity of the polishedsurface layer was lower than the interior of the magnet
due to the deterioration and lack of a Nd-rich grainboundary phase. When a group of grains was reversed
in a certain applied magnetic field, positions where the28magnetization reversal stopped often existed along thegrain boundary. To clarify how the misorientation angle
between two grains on both sides of the grainboundary affects the magnetization reversal, the grainboundaries were divided into two groups: GBa and GBb.
GBa is a grain boundary where the magnetization
reversal stopped, which is indicated by the borders ofthe different colored regions in Fig. 6. GBb is a grainboundary where the magnetization reversal did not stopthat is located inside the colored region. In this study,
the relationship between GBa and Heff, as well as GBb
and Heff, was not considered. The grain boundaries thatwere clearly separated by a large Nd-rich phase were
excluded from the evaluation. Figure 8(a) shows theresult of dividing about 900 grain boundaries into GB aand GBb for the x = 0 samples and the ratio of GBa and
GBb in each angular range of . It was found that theratio of GBa became larger when  became larger. Thistendency was similar when about 1000 grainboundaries were analyzed in the samples with x = 5.0,
as shown in Fig. 8(b). These results experimentallysuggest that a grain boundary with a larger
misorientation angle of adjacent grains  has the ability
to suppress magnetization reversal. It is also consideredthat the influence of the misorientation angle on
suppressing magnetization reversal does not differ
much even when the saturation magnetization andmagnetocrystalline anisotropy of the main phase aredifferent.
4. Conclusion
To clarify the difference between the degree ofalignment dependence of coercivity and angulardependence of coercivity, the crystal orientation

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

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distribution and demagnetization curve of Nd-Fe-B
based sintered magnets with different degrees ofalignment were compared. It was shown that anincrease in coercivity due to a low degree of alignment
cannot be explained only by the angular dependence of
coercivity and that domain wall movement is small in amagnet that has a low degree of alignment. A crystal
orientation analysis and in-situ magneto-optical Kerr
microscope observation performed in the same area
clarified that the ratio of the grain boundary where themagnetization reversal stopped became larger when themisorientation angle between adjacent grains became
larger. This suggests that a grain boundary having alarger misorientation angle is one of the factors thatsuppresses magnetization reversal. From the above
results, it can be said that the increase in coercivity by
the decrease in the degree of alignment is caused by the
suppression of magnetization reversal at the grainboundary due to the increase in the misorientation
angle between adjacent grains. The reason that themisorientation angle suppresses magnetization reversal
is considered to be the reduction in magnetic interaction
between grains or a change in the microstructural
relationship between the main phase and grainboundary phase. Further understanding of the grainboundary structure and magnetization process isnecessary in the future.
References
1) K. Hiraga, M. Hirabayashi, M. Sagawa, and Y. Matsuura:
Jpn. J. Appl. Phys., 24, L30 (1985).
2) R. K. Mishra, J. K. Chen, and G. Thomas: J. Appl. Phys., 59,
2244 (1986).
3) F. Vial, F. Joly, E. Nevalainen, M. Sagawa, K. Hiraga, andK. T. Park: J. Magn. Magn. Mater., 242-245, 1329 (2002).

4) Y. Shinba, T. J. Konno, K. Ishikawa, and K. Hiraga: J. Appl.
Phys., 97, 053504 (2005).
5) H. Sepehri-Amin, T. Ohkubo, T. Shima, and K. Hono: Acta.
Mater., 60, 819 (2012).
6) A. Handstein, K. -H. Muller, D. Eckert, and P. Northnagel:
J. Magn. Magn. Mater., 101, 382 (1991).
7) Y. Matsuura, J. Hoshijima, and R. Ishii: J. Magn. Magn.
Mater., 336, 88 (2013).
8) D. Givord, P. Tenaud, and T. Viadieu: J. Magn. Magn.
Mater., 72, 247 (1988).
9) D. Li and K. J. Strnat: J. Appl. Phys., 57, 4143 (1985).
10) A. Fukuno and R. C. O’Handley: J. Appl. Phys., 65, 4959
(1989).
11) Y. -G. Park and D. Shindo: J. Magn. Magn. Mater., 238, 68(2002).
12) T. Yamaoka, H. Tsujikawa, R. Hirose, A. Ito, H.
Kawamura, and T. Sakon: J. Magn. Soc. Jpn., 35, 60 (2011)[in Japanese].
13) J. Thielsch, T.G. Woodcock, L. Schultz, and O. Gutfleisch:
J. Appl. Phys., 111, 103901 (2012).
14) T. Kohashi, K. Motai, T. Nishiuchi, T. Maki, and S.
Hirosawa: J. Magn. Soc. Jpn., 33, 374 (2009).
15) H. Suzuki, Y. Satsu, T. Kohashi, K. Motai, and M. Komuro:
J. Appl. Phys., 109, 07A746 (2011).
16) M. Takezawa, T. Shimada, S. Kondo, S. Mimura, Y.
Morimoto, T. Hidaka, and J. Yamasaki: J. Appl. Phys., 101,
09K106 (2007).
17) M. Takezawa, Y. Kimura, Y. Morimoto, and J. Yamasaki:IEEE Trans. Magn., 49, 3262 (2013).
18) M. Takezawa, H. Ogimoto, Y. Kimura and Y. Morimoto: J.
Appl. Phys., 115, 17A733 (2014).
19) K. Toyoki, Y. Kotani, Y. Senba, D. Billington, H. Okazaki,
A. Yasui, W. Ueno, H. Ohashi, Y. Shiratsuchi, S. Hirosawa,
K. Hono, and T. Nakamura: Proc. 24th Int. Workshop on
Rare-Earth and Future Permanent Magnets and their
Applications, Darmstadt, 2016, 129 (2016).
20) T. Maki and S. Hirosawa: J. Appl. Phys., 103, 043904(2008).
Received Oct. 2, 2017; Accepted Jan. 18, 2018

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J. Magn. Soc. Jpn., 42, 30-36 (2018)<Paper>

Structure Characterization of Fe, Co, and Ni Thin Films
Epitaxially Grown on GaAs(111) Substrate
Takahiro Soda1, Shigeyuki Minakawa1, Masaaki Futamoto1, Mitsuru Ohtake1,2, and Nobuyuki Inaba3
Faculty of Science and Engineering, Chuo University, 1-13-27 Kasuga, Bunkyo-ku, Tokyo 112-8551, Japan
2Faculty of Engineering, Kogakuin University, 2665-1 Nakano, Hachioji, Tokyo 192-0015, Japan
3Faculty of Engineering, Yamagata University, 4-3-16 Jyonan, Yonezawa, Yamagata 992-8510, Japan

1

Fe, Co, and Ni films of 40 nm thickness are prepared on GaAs(111) single-crystal substrates at room
temperature by using a radio-frequency magnetron sputtering system. The film growth behavior and the
crystallographic properties are investigated by in-situ reflection high-energy electron diffraction and pole-figure
X-ray diffraction. bcc single-crystals of (111) orientation are formed on the substrates for all the film materials,
though the bcc structure is metastable for Co and Ni materials. The metastable structure is stabilized throughhetero-epitaxial growth. Fe films possess bcc structure for the investigated thickness range. On the contrary, thebcc-Co and the bcc-Ni crystals, respectively, start to transform into hcp and fcc structures, as the thickness isincreased_ beyond_2 nm. The phase_transformations occur through atomic displacements from the close-packed planesof bcc(110), bcc(101), and bcc(011), which are perpendicular to the substrate surface, to hcp(0001) and fcc(111)close-packed planes. The crystallographic orientation relationships of hcp and fcc crystals with respect to bcc crystal
are similar to the Kurdjumov-Sachs relationship.
Key words: 3d transition metals, epitaxial growth, thin films, GaAs(111) substrate, metastable bcc structure, phasetransformation
1.

Introduction

Hybrid structure of ferromagnetic metal andsemiconductor materials has been investigated for spinelectronics applications like spin-dependent field effecttransistors, etc. bcc phase is stable for Fe, whereas thatis metastable for Co and Ni at room temperature (RT).
It has been reported that magnetic tunnel junction
elements prepared by using Co films with metastablebcc structure show high tunnel magnetoresistanceratios1,2). The magnetic and electronic properties varydepending on the crystal structure. It is thus importantto understand the formation conditions of magnetic
films with metastable bcc structure.
Metastable bcc phase formation has beenrecognized for very thin Co3–5) and Ni6) films on GaAs
single-crystal substrates of (100) and (110) orientations
around RT deposited by molecular beam epitaxy (MBE).
With increasing the thickness, most of the bcc-Co andthe bcc-Ni crystals transformed into more stable hcp or
fcc structure and the resulting films involved large
volumes of hcp or fcc crystals. Recently, we succeeded in
the formation of Co and Ni films with bcc structure on
GaAs(100)7,8) and GaAs(110)9) substrates by sputtering,which is more suitable than MBE for mass-production
applications. The thickness stability of bcc phase andthe transformation process to hcp or fcc structure were
investigated. With increasing the thickness beyond 2
nm, the bcc crystals formed on GaAs(100) andGaAs(110) substrates, respectively, started to transform
into stable structures through atomic displacements
parallel to theclose-packed planes and the_ six bcc{110}_four bcc(101 ), bcc(011 ), bcc(011), and bcc(101) planes
which are 60° inclined from the substrate surface. The
slide planes, where the phase transformation occurs,
differ depending on the substrate orientation.
In our previous study10), Co films were prepared on
GaAs(111) substrates by varying the substratetemperature from RT to 600 °C. The film structure was30

investigated by reflection high-energy electron
diffraction (RHEED), where_ the incident electron beam
was parallel to only GaAs[110]. In the early stage of film
growth at temperatures lower than 200 °C, Co(111)crystals with bcc structure were formed. With
increasing the thickness, the bcc(111) crystals started totransform into the stable phase, similar to the cases of
Co films formed on GaAs(100)7,8) and GaAs(110)9)
substrates. When the substrate temperature was higher
than 400 °C, Ga atoms of substrate seemed to diffuse
into the Co films and an ordered alloy with bcc-based B2
structure was formed.
In the present study, Fe, Co, and Ni films areprepared on GaAs(111) substrates at RT under similar
deposition conditions. The detailed growth behaviors
are studied by RHEEDof incident_ using two kinds_
electron beam, GaAs[110] and GaAs[112]. The resulting
structure is characterized by out-of-plane, in-plane, andpole-figure X-ray diffractions (XRDs). The influences offilm material and thickness on the crystallographic
properties are systematically investigated.
2.

Experimental Procedure

A radio frequency (RF) magnetron sputteringsystem equipped with an RHEED facility was employed.
The base pressure was lower than 4 × 10 –7 Pa.
GaAs(111) substrates were heated at 600 °C before
deposition to obtain clean surfaces. Figures 1(a) and 2(a)show the RHEED patterns observed for a GaAs
substrate after annealing.
The incident electron beam is_parallel to GaAs[11_ 0] in Fig. 1(a), whereas that isparallel to GaAs[112] in Fig. 2(a). Diffraction patterns
consisting of spots are observed, which may suggest thatthe substrate surface involves an atomic level roughness.
Figures 1(b) and 2(b) show the schematic diagrams ofdiffraction patterns of a B3(111) single-crystal surface
calculated by using the two kinds of electron beam
direction. The experimental data of Figs. 1(a) and 2(a)Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX(b) B3(a) GaAs(f) bcc(a) GaAs(f) bcc(b) B3222
222004111111
000

220

112
002

_

110

(c-1) Fe(1 nm)(d-1) Co(1 nm)(e-1) Ni(1 nm)(c-1) Fe(1 nm)(d-1) Co(1 nm)(e-1) Ni(1 nm)(c-2) Fe(2 nm)(d-2) Co(2 nm)(e-2) Ni(2 nm)(c-2) Fe(2 nm)(d-2) Co(2 nm)(e-2) Ni(2 nm)(c-3) Fe(10 nm)(d-3) Co(10 nm)(e-3) Ni(10 nm)(c-3) Fe(10 nm)(d-3) Co(10 nm)(e-3) Ni(10 nm)(c-4) Fe(40 nm)(d-4) Co(40 nm)(e-4) Ni(40 nm)(c-4) Fe(40 nm)(d-4) Co(40 nm)(e-4) Ni(40 nm)

111
220 000 220_000

110 000 110_

Fig. 1 (a) RHEED pattern observed for a GaAs(111)substrate. (b) Schematic diagram of RHEED pattern
simulated for a B3(111) substrate. (c)–(e) RHEED
patterns observed for (c) Fe, (d) Co, and (e) Ni films
deposited on GaAs substrates at RT. The film
thicknesses are (c-1)–(e-1) 1, (c-2)–(e-2) 2, (c-3)–(e-3) 10,
and (c-4)–(e-4) 40 nm. (f) Schematic diagram of abcc(111) single-crystal with reconstructed surface_ ofp(3×3). The incident electron beam is parallel to [110].

Fig. 2 (a) RHEED pattern observed for a GaAs(111)substrate. (b) Schematic diagram of RHEED pattern
simulated for a B3(111) substrate. (c)–(e) RHEED
patterns observed for (c) Fe, (d) Co, and (e) Ni films
deposited on GaAs substrates at RT. The film
thicknesses are (c-1)–(e-1) 1, (c-2)–(e-2) 2, (c-3)–(e-3) 10,
and (c-4)–(e-4) 40 nm. (f) Schematic diagram of abcc(111) single-crystal with reconstructed surface_ ofp(3×3). The incident electron beam is parallel to [112].

agree with the simulation results of Figs. 1(b) and 2(b),respectively. A clean GaAs(111) surface with B3
structure is thus obtained. A chemical analysis isnecessary to determine the kind of atom terminated at
the surface.
Fe, Co, and Ni films were prepared on thesubstrates at RT by varying the thickness in a range
from 1 to 40 nm. The distance between target andsubstrate was fixed at 150 mm. The Ar gas pressure
was kept constant at 0.67 Pa. Fe, Co, and Ni targets of 3
inch diameter were employed and the respective RF
powers were fixed at 48, 54, and 50 W, where thedeposition rate was 0.02 nm/s for all the materials.
The surface structure was studied by RHEED. The
resulting film structure was investigated by XRDs with
Cu-Kα radiation (λ = 0.15418 nm). The magnetization
curvesweremeasuredbyvibratingsamplemagnetometry.

RHEED patterns simulated for a bcc(111) single-crystal,
where the open circle spots correspond to the reflections
from reconstructed surface of p(3×3). The RHEED
patterns observed for Fe film are in agreement with thesimulated patterns without open circle spots, whereas
those observed for Co and Ni films are matching withthe simulated patterns involving open circle spots. bcc
structure is thus stabilized not only for the Fe film but
also for the Co and the Ni films through hetero-epitaxial
growth on GaAs(111) substrate, similar to the cases offilms prepared on GaAs(100)7,8) and GaAs(110)9)
substrates. The epitaxial relationship is determined by
RHEED as__bcc(111)[110] || GaAs(111)[110].

3.

Results and Discussion

Figures 1(c-1)–(e-1) and 2(c-1)–(e-1) show the
RHEED patterns observed for Fe, Co, and Ni films of 1
nm thickness prepared on GaAs(111) substrates. Clear
RHEED patterns are recognized for all the film
materials. The films grow epitaxially on the substrates.
Figures 1(f) and 2(f) show the schematic diagrams of

Figures 1(c-2)–(c-4) and 2(c-2)–(c-4) show the
RHEED patterns observed for Fe films thicker than 2
nm. RHEED patterns corresponding to the diffraction
patterns simulated for bcc(111) surface remainunchanged till the end of 40-nm-thick film formation.
Figures 1(d-2)–(e-4) and 2(d-2)–(e-4) show the RHEED
patterns observed for Co and Ni films thicker than 2 nm.
As the thickness increases, the RHEED spots become
broader and diffraction spots different from those ofbcc(111) surface appear. The result indicates that aphase transformation is taking place.
When a bulk bcc material transforms into fcc

Journal of the Magnetics Society of Japan Vol.42, No.2, 201831

INDEXbcc(110)(a)bcc(101)bcc(011)NW

KS-1

ABAB...
__[111]KS-1__[111]KS-2_[111]KS-2

_

_

_

___

__

5°

___

__

bcc{110}
Matching

ABAB...
ABAB...
hcp(0001)[1010] ||hcp(0001)[0110] ||bcc(110)[111]bcc(110)[111]bcc(101)[111]bcc(101)[111]bcc(011)[111]bcc(011)[111]bcc(110)[111]bcc(110)[111]bcc(101)[111]bcc(101)[111]bcc(011)[111]bcc(011)[111]
_

_

__ _
_ _
_
__ _

___ _
___
_
___ _
_ _
__
__
_

ABAB...
hcp(0001)[1120] ||bcc(110)[001]bcc(101)[010]bcc(011)[100]bcc(110)[001]bcc(101)[010]bcc(011)[100]Matching__

Matching

_

_

_

_

Fig. 3(a) Configuration of bcc{110} close-packed
planes in bcc crystal with the (111) plane parallel to thesubstrate surface. (b) Crystallographic orientation
relationships of bcc-hcp phase transformation.
structure, there are two possibilities in the
crystallographic orientation relationship,
Nishiyama-Wasserman (NW)11,12) and Kurdjumov-Sachs(KS) 13) . The phase transformation occurs throughatomic displacements from six bcc{110} close-packed
planes to fcc(111) close-packed plane. The
crystallographic orientation relationship of bcc-hcp
phase transformation seems similar to that of bcc-fcc
transformation, as shown in Fig. 3. The phasetransformation in Co film is considered to occur throughatomic displacements from all the six {110} close-packed
planes, from the three close-packed planes of bcc(110),bcc(101), and bcc(011) which are 35° canted from thesubstrate _surface, or_from the three close-packed
planes_o f bc c ( 11 0 ) , b c c ( 1 01), an d bc c ( 011 ) w h i c h a r eperpendicular to the substrate surface. Figures 4(a) and5(a) show the schematic diagrams of RHEED patterns
calculated for hcp crystals transformed through atomic
displacements from the bcc(110), bcc(101), and bcc(011)planes. Figures 4(b) and 5(b) show the schematic
diagrams of RHEED patterns simulated for hcp crystals
transformed
throughatomic displacementsfrom the___bcc(110), bcc(101), and bcc(011) planes. The RHEED
patterns observed for Co films thicker than 2 nm [Figs.
1(d-2)–(d-4), 2(d-2)–(d-4)] agree with the simulated
patterns of Figs. 4(b-3) and 5(b-3), as shown in Fig. 6.
The result shows that the phase transformation is nottaking place through atomic displacements from the six
bcc{110} planes
but_from the three close-packed
planes__of bcc(11 0), bcc(1 01), and bcc(011 ) in the KS-2

32

Fig. 4
Schematic diagrams of RHEED patterns
simulated for hcp crystals with the atomic stacking
sequence of ABAB… along hcp[0001] transformed froma bcc(111) crystal epitaxially grown on GaAs(111)substrate through atomic displacements (a) frombcc(110), bcc(101),
and_ bcc(011) close-packed
planes and__(b) from bcc(110), bcc(101), and bcc(011) planes in [(a-1),(b-1)] the NW, [(a-2), (b-2)] the KS-1, and [(a-3), (b-3)]
KS-2 relationships.
_ The incident electron beam isparallel to GaAs[110].

NW

KS-1

ABAB...

ABAB...

bcc(110)+(101)+(011) → hcp(0001)
(a-1)
(a-2)

bcc(110)+(101)+(011) → hcp(0001)
(b-1)
(b-2)
_

5°

KS-2

_

_ _

hcp(0001)

KS-1

_

_

NW

(b-3)

[111]KS-1[111]KS-1 [111]KS-1(b)[111]KS-2 [100]NW___[111]KS-2_[010]NW

bcc(110)+(101)+(011) → hcp(0001)
(b-1)
(b-2)
_

bcc(011)bcc(101)_bcc(110)[111]KS-2(a-3)[010]NW [100]NW[111]KS-2[001]NW

KS-2

ABAB...

_

(111)
c a
b

bcc(110)+(101)+(011) → hcp(0001)
(a-1)
(a-2)

[111]KS-1_ _[111]KS-1[001]NW

ABAB...

KS-2

ABAB...
(a-3)

(b-3)

Fig. 5
Schematic diagrams of RHEED patterns
simulated for hcp crystals with the atomic stacking
sequence of ABAB… along hcp[0001] transformed froma bcc(111) crystal epitaxially grown on GaAs(111)substrate through atomic displacements (a) frombcc(110), bcc(101),
and_ bcc(011) close-packed
planes and__(b) from bcc(110), bcc(101), and bcc(011) planes in [(a-1),(b-1)] the NW, [(a-2), (b-2)] the KS-1, and [(a-3), (b-3)]
KS-2 relationships._ The incident electron beam isparallel to GaAs[112].
(a) Co(10 nm)(b) Co(10 nm)

Fig. 6 RHEED patterns observed for a 10-nm-thick Co
film prepared on GaAs(111) substrate [Figs. 1(d-3),2(d-3)] overlapped with reflection spots simulated for
hcp crystals
through atomic
displacements_ transformed__
from bcc(110), bcc(101), and bcc(011) planes in the KS-2
relationship [Figs. 4(b-3), 5(b-3)].
The incident _electron_
beam is parallel to (a) GaAs[110] or (b) GaAs[112].

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEXKS-2
5°bcc{110}ABCABC...
fcc(111)[011] ||bcc(110)[111]bcc(101)[111]bcc(011)[111]bcc(110)[111]bcc(101)[111]bcc(011)[111]ACBACB...
fcc(111)[011] ||bcc(110)[111]bcc(101)[111]bcc(011)[111]bcc(110)[111]bcc(101)[111]bcc(011)[111]
_

_
_

_

(a-3)

ACBACB...

bcc(110)+(101)+(011) → fcc(111)
(b-1)
(b-2)

(b-3)

__ _
_ _
_
__ _

KS-1

bcc(110)+(101)+(011) → fcc(111)
(a-1)
(a-2)

ACBACB...

bcc(110)+(101)+(011) → fcc(111)
(b-1)
(b-2)
_

_

_

_

_

_

_

_

_

_

_

The crystallographic relationships are slightly different
from the results reported for Co film growth on
GaAs(111) substrate10). In the present study, twodifferent incident electron beam directions are employedto accurately determine the relationships, since broad
RHEED spots may include position errors.
Phase transformation in Ni film is also considered
to occur through atomic displacementsfrom the three___close-packed planes of bcc(110), bcc(101), and bcc(011) tofcc(111) close-packed plane. However, there are twokinds of atomic stacking sequence of fcc(111) plane
along the fcc[111] direction, ABCABC… and ACBACB…,
for an fcc crystal, while that of hcp(0001) plane along
the hcp[0001] direction is only ABAB… Figure 7
summarizes the crystallographic orientation
relationships of bcc-fcc phase transformation. Figures
8(a) and 9(a) show the schematic diagrams of RHEED
patterns simulated for fcc crystals with the atomic
stacking sequence of ABCABC… transformed__ throughatomic _displacements from the bcc(110), bcc(101), andbcc(011 ) planes. Figures 8(b) and 9(b) show theschematic diagrams of RHEED patterns calculated forfcc crystals with the ACBACB… stacking sequencetransformed__through atomic _displacements from thebcc(110), bcc(101), and bcc(011) planes. Figures 10(a)and (b) show the RHEED patterns observed for the
10-nm-thick Ni film [Figs. 1(e-3), 2(e-3)] overlapped withthe simulated patterns of fcc crystals with stacking

NW

ABCABC...
_

__ _
_ _
_
__ _

___ _
_
___
___ _
_ _
__
__

_

relationship. Therefore, the Co films thicker than 2 nm
involve three hcp variants whose c-axes are lying in thefilm plane and rotated around the film normal by 120°each other. The transformation direction seems to be
influenced by the strain caused by accommodation of
the lattice misfit between the film and the substrate.
The crystallographic orientation relationships are
determined as follows,__ __hcp(0001)[0110] || bcc(110)[111],___ _hcp(0001)[0110] || bcc(101)[111],__ __hcp(0001)[0110] || bcc(011)[111].

_

_

___ _
_
___
___ _
_ _
__
__
_

Fig. 7 Crystallographic orientation relationships ofbcc-fcc phase transformation.

Fig. 8
Schematic diagrams of RHEED patterns
simulated for fcc crystals with the atomic stacking
sequences of (a) ABCABC… and (b) ACBACB… alongfcc[111] transformed from a bcc(111) crystal epitaxiallygrown on GaAs(111) _substrate_ through atomic_
displacements from bcc(1 1 0), bcc( 1 01), and bcc(01 1 )planes in [(a-1), (b-1)] the NW, [(a-2), (b-2)] the KS-1,
and [(a-3), (b-3)] KS-2 relationships.
The incident_
electron beam is parallel to GaAs[110].

_

_

_
_

ABCABC...
fcc(111)[101] ||bcc(110)[111]bcc(101)[111]bcc(011)[111]bcc(110)[111]bcc(101)[111]bcc(011)[111]ACBACB...
fcc(111)[101] ||bcc(110)[111]bcc(101)[111]bcc(011)[111]bcc(110)[111]bcc(101)[111]bcc(011)[111]KS-2

Matching_

ABCABC...
fcc(111)[110] ||bcc(110)[001]bcc(101)[010]bcc(011)[100]bcc(110)[001]bcc(101)[010]bcc(011)[100]ACBACB...
fcc(111)[110] ||bcc(110)[001]bcc(101)[010]bcc(011)[100]bcc(110)[001]bcc(101)[010]bcc(011)[100]Matching

KS-1

bcc(110)+(101)+(011) → fcc(111)
(a-1)
(a-2)

_

Matching

NW

ABCABC...

_

5°

_

KS-1

_

fcc(111)

_

NW

KS-2
(a-3)

(b-3)

Fig. 9
Schematic diagrams of RHEED patterns
simulated for fcc crystals with the atomic stacking
sequences of (a) ABCABC… and (b) ACBACB… alongfcc[111] transformed from a bcc(111) crystal epitaxiallygrown on GaAs(111) _substrate_ through atomic_
displacements from bcc(1 1 0), bcc( 1 01), and bcc(01 1 )planes in [(a-1), (b-1)] the NW, [(a-2), (b-2)] the KS-1,
and [(a-3), (b-3)] KS-2 relationships.
The incident_
electron beam is parallel to GaAs[112].
(a) Ni(10 nm)(b) Ni(10 nm)

Fig. 10 RHEED patterns observed for a 10-nm-thickNi film prepared on GaAs(111) substrate [Figs. 1(e-3),2(e-3)] overlapped with reflection spots simulated for fcc
crystals with the stacking sequences of ABCABC... and
ACBACB... along fcc[111] _transformed
though atomic__
displacements from bcc(1 1 0), bcc( 1 01) and, bcc(01 1 )planes in the KS-2 relationship [Figs. 8(a-3), 8(b-3),
9(a-3), 9(b-3)].
beam is parallel to_ The incident electron_(a) GaAs[110] or (b) GaAs[112].

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

33

INDEX

34

GaAs(333)

GaAs(444)

Intensity (arb. unit)

bcc-Fe(222)(a) Fe(b) Co(c) Ni
40

20

60

80
100
2θ (deg.)

120

140

Fig. 11 Out-of-plane XRD patterns measured for (a) Fe,(b) Co, and (c) Ni films of 40 nm thickness formed on
GaAs(111) substrates. The intensity is shown inlogarithmic scale.
GaAs(220)bcc-Fe(110)_WL(a) FeIntensity (arb. unit)

The slide planes and the thickness, where thetransformation from bcc to hcp or fcc phase starts, aresimilar between Co and Ni films. Furthermore, thecritical thickness of around 2 nm is almost similar to
those observed for Co and Ni films prepared on
GaAs(100)7,8) and GaAs(110)9) substrates, though the
crystallographic orientation relationships of hcp or fcc
crystals with respect to bcc crystal are different in thecases of film deposition on GaAs(100), GaAs(110), andGaAs(111) substrates.
Figures 11(a) and 12(a) show the out-of-plane andthe in-plane XRD patterns of 40-nm-thick Fe film
formed on GaAs(111) substrate. Here, the scattering_vector of in-plane XRD is parallel to GaAs[1 1 0].
Fe(222)bcc reflection is clearly observed in addition to
GaAs (444) reflectionin the out-of-plane pattern,_whereas Fe(1 1 0)bcc reflection is recognized in thein-plane pattern. The out-of-plane and the in-plane
XRDs confirm the epitaxial orientation relationshipdetermined by RHEED. Figures 11(b) and (c) show theout-of-plane XRD patterns of 40-nm-thick Co and Ni
films, respectively. Reflections from the hcp and the fcc
crystals transformed from the bcc structure are notrecognized for the Co and the Ni films, since the low
index planes of transformed hcp-Co and fcc-Ni crystals
are not parallel to the substrate surface. Figures 12(b)and (c) show the in-plane XRD patterns of 40-nm-thickCo and Ni films, respectively. hcp(0002) reflection isobserved for the Co film, whereas fcc(111) reflection isrecognized for the Ni film. The out-of-plane and thein-plane XRDs _confirm _that the transformations
occur_
from the bcc(110), bcc(101) and, bcc(011) slide planes
which are perpendicular to the substrate surface.
Figures 13(a)–(c) show the pole-figure XRD
patterns of 40-nm-thick Fe, Co, and Ni films measured
by fixing the diffraction_ angle of 2θB at 47.5º, where
GaAs{220} and hcp{1011} reflections are expected to be
detectable. The signals with the intensities smaller than
10 cps are subtracted. Figures 13(d) and (e), respectively,
show the schematic diagrams of pole-figure XRD
patterns simulated for a GaAs(111) substrate and thehcp crystals transformed from bcc structure in theorientation relationships determined by RHEED. The
pole-figure XRD patterns of Fe and Ni films show
three-fold symmetry, where GaAs{220} reflections arerecognized at the tilt angle, α, of 55º and the rotation
angles, β, of 120º, 240º, and 0º. The_ _ pole-figure XRD
pattern of Co film shows six hcp{1101} reflections from

GaAs(111)_

sequences of ABCABC... andwhich are
_ ACBACB...
__transformed from the bcc(110), bcc(101), and bcc(011)planes in the KS-2 relationship [Figs. 8(a-3), 8(b-3),
9(a-3), 9(b-3)]. The observed RHEED patterns are wellmatchedwith
thecalculated
patterns.
The
crystallographic orientation relationships are thus
determined as follows,__ __fcc(111)[011] || bcc(110)[111],___ _fcc(111)[011] || bcc(101)[111],__ __fcc(111)[011] || bcc(011)[111],__ __fcc(111)[011] || bcc(110)[111],___ _fcc(111)[011] || bcc(101)[111],__ __fcc(111)[011] || bcc(011)[111].

hcp-Co(0002)(b) Cofcc-Ni(111)(c) Ni
20

30

40
50
2θχ (deg.)

60

70Fig. 12 In-plane XRD patterns of (a) Fe, (b) Co, and (c)Ni films of 40 nm thickness formed on GaAs(111)substrates measured_ by making the scattering vector
parallel to GaAs[1 1 0]. The intensity is shown inlogarithmic scale.
hcp crystals at the tilt angle, α , of 44º. The hcp
reflections are not observed for the Ni film. Figures
14(a)–(c) show the pole-figure XRD patterns of
40-nm-thick Fe, Co, and Ni films measured by fixing thediffraction angle of 2θB at 51.8º, where GaAs{311} andfcc{001} reflections are expected to be detectable. The
signals with the intensities smaller than 10 and 1 cps
are subtracted for the respective films. Figures 14(d)and (e), respectively, show the schematic diagrams ofpole-figure XRD patterns simulated for a GaAs(111)substrate and the fcc crystals transformed from bcc
structure in the orientation relationships determined by
RHEED. Reflections which originate from the GaAs
substrate, are recognized around α = 10°, 32°, and 61°
for all the films [Figs. 14(a-2), (b-2), and (c-1)].
Reflections, which originate from the fcc crystals, areobserved around α = 9°, 38° and, 51° for the Ni film [Fig.
14(c-1)]. On the other hand, the fcc reflections are not

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX(a) Fe(b) Coβ = 90°(c) Niβ = 0°

Intensity(cps)ββ = 180°

90°
60°
30°

α = 0°

103
101

Threshold10–1

β = 270°(e) hcp

__

__

__

_

_

_ _

_

_

_

__

_ _

__

_

_

_

_

(220)

__

(202)

(022)

hcp(0001)[0110] || bcc(110)[111]hcp(0001)[0110] || bcc(101)[111]hcp(0001)[0110] || bcc(011)[111]
_
_
_

_

_

(022)

__

(220)

_

(202)

(1101)(0111)(0111) (1011)(1101)(0111) (1101)(1011) (1011)(0111)(1011) (1011)(1101)(1101)(1011)(0111)(0111)(1101)

(022)

_

(202)

__ _
_ _
_
__ _

(220)

_

(d) B3-GaAs_

Fig. 13 (a)–(c) Pole-figure XRD patterns measured of (a) Fe, (b) Co, and (c) Ni films of 40 nm thickness prepared on
GaAs(111) substrates measured by fixing the diffraction angle of 2θB at 47.5º, where signals with the intensities
smaller than 10 cps are subtracted. [(d), (e)] Schematic diagrams of diffraction patterns simulated for (d) GaAs(111)substrate and (e) hcp crystals transformed from bcc structure in the crystallographic orientation relationships
determined by RHEED.
(c-1) Niβ = 0°

β = 180°

90°
60°

Intensity(cps)(b-1) Coβ = 90°β103Intensity(cps)(a-1) Fe10330°

α = 0°

101
10–1

Thresholdβ = 270°(a-2) Fe(b-2) Co(e) fcc(311)_(010)ACB_(001)ACB(010)ABC(001)ABC (100)ABC(100)ACB_(010)ACB(001) ACB___(100) ACB(010)ABC (001) ABC
_

_

_

__

_

(113)

(100)ABC (010)ABC

fcc(111)[011] & [011] || bcc(110)[111]fcc(111)[011] & [011] || bcc(101)[111]fcc(111)[011] & [011] || bcc(011)[111]
_
_
_

(113)

_
_
_

(113)
_

_

(131)
(131)
(131)

(100)ACB(001)ACB(100)ABC(001)ABC(010) ACB

(113)

(131)

10–1

__ _
_
_ _
__ _

__

(311) (311)Threshold

_

(311)

101

_

(d) B3-GaAs(c-2) Ni_

_ _Fig. 14 (a)–(c) Pole-figure XRD patterns of (a) Fe, (b) Co, and (c) Ni films of 40 nm thickness prepared on GaAs(111)substrates measured by fixing the diffraction angle of 2θB at 51.8º, where signals with the intensities smaller than(a-1)–(c-1) 1 and (a-2)–(c-2) 10 cps are subtracted. [(d), (e)] Schematic diagrams of diffraction patterns simulated for(d) GaAs(111) substrate and (e) fcc crystals transformed from bcc structure in the crystallographic orientation
relationships determined by RHEED.
observed for the Co film. Therefore, the bcc crystals in
Co and Ni films are apparently transforming into hcp
and fcc structures in the orientation relationships
determined by RHEED, respectively.
Figure 15(a) shows the magnetization curves of
10-nm-thick bcc-Fe(111) single-crystal film_measured by_applying the magnetic field along GaAs[112], GaAs[101],

__
_
_
_
GaAs[211], GaAs[110], GaAs[121], or GaAs[011]. The
magnetization curves are almost isotropic in thein-plane measurements. The easy magnetization axes ofbcc[100], bcc[010], and bcc[001] are not parallel to thesubstrate surface but 35° canted from the substratesurface. Therefore, the effective easy magnetization__
directions are considered to be observed along bcc[211],

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

35

INDEX

Ms = 1.5

Ms = 1.3
kemu/cm3(b)

1.0
kemu/cm3
kemu/cm3

Ms = 0.5
kemu/cm3(c)_

GaAs[011] GaAs[121] GaAs[110] GaAs[211] GaAs[101]_GaAs[112]

(a)

__

_

4.

_
_
–2000

0

2000

–2000
0
2000
Applied field (Oe)

–2000

0

2000

Fig. 15 Magnetization curves measured for (a) Fe, (b)Co, and (c) Ni films of 10 nm thickness prepared on
GaAs(111) substrates.
[211]
__

In
-p

[100]

lan
e

[001] [112]
__

[111]

Conclusion

Fe, Co, and Ni films are prepared by sputter
deposition on GaAs(111) substrates at RT by varying
the thickness in a range from 1 to 40 nm. Nucleation of
crystal with metastable structure and phasetransformation process into stable structure areinvestigated by RHEED and XRDs. bcc(111)single-crystals are formed in the early stages of film
growth of not only Fe but also Co and Ni. The
metastable structure is stabilized for Co and Ni films
through hetero-epitaxial growth on GaAs(111) substrate.
With increasing the thickness beyond 2 nm, the bcc-Co
and the bcc-Ni crystals, respectively, start to transform
into hcp and fcc structures__ through atomic_
displacements in the bcc(11 0), bcc(1 01) and bcc(011 )close-packed planes.
Acknowledgements A part of this work was supported
by Chuo University Grant for Special Research.

Tilt angle (deg.)
030 60 90

[010]
[121]
_ _

Fig. 16 Pole figure map of bcc(111) single-crystal film
showing angles of typical crystallographic planes tilted
from the film plane.
_ ___bcc[121], and bcc[112] due to the demagnetization field,
as shown in Fig. 16. As a result, the in-plane magnetic
anisotropy seems to show a six-fold symmetry. It should
be noted that a contribution to in-plane magnetic
anisotropy decreases with increasing the tilt angle, andthat the angle of bcc<100> with respect to the film plane
is as large as 35°. Therefore, the Fe(111) single-crystal
film prepared in the present study is considered to show
an almost isotropic in-plane magnetic property. The
saturation magnetization value of Fe film is less thanthat of bulk Fe crystal (1.73 kemu/cm3). The reason ispossibly due to surface oxidation of the film sample withno overcoat.
Figures 15(b) and (c), respectively, show themagnetization curves measured for Co and Ni films of
10 nm thickness. The films, which include hcp and fcc
crystals, show almost isotropic magnetic properties. The

36

easy magnetization axes for hcp and fcc crystals are
hcp[0001] and fcc<111>. These crystallographic__
directions correspond to the GaAs[10 1 ] and [1 1 0],
respectively. When the distributions of multiple easy
axes and the three-fold crystallographic symmetries of
Co and Ni films on GaAs(111) substrate are considered,the magnetic anisotropy of whole film will be observed
as nearly isotropic. Such nearly isotropic magnetic
properties observed in the present study indicate thatthe magnetocrystalline anisotropies of three hcp-Co andsix fcc-Ni crystals are overlapped. The magnetization
property is different from the cases of Co and Ni films
formed on GaAs(110) substrates where the films consistof variants with two-fold symmetrical distributions9).

References
1) X. -G. Zhang and W. H. Butler: Phys. Rev. B, 70, 172407 (2004).
2) S. Yuasa, A. Fukushima, H. Kubota, Y. Suzuki, and K. Ando: Appl.
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Lett., 62, 2480 (1989).
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A. C. Bland, and L. M. Brown: Phys. Rev. B, 52, 14704 (1995).
6) C. S. Tian, D. Qian, D. Wu, R. H. He, Y. Z. Wu, W. X. Tang, L. F.
Yin, Y. S. Shi, G. S. Dong, X. F. Jin, X. M. Jiang, F. Q. Liu, H. J.
Qian, K. Sun, L. M. Wang, G. Rossi, Z. Q. Qiu, and J. Shi: Phys.
Rev. Lett., 94, 137210 (2005).
7) M. Ohtake, S. Minakawa, and M. Futamoto: Key Eng. Mater., 605,
478 (2014).
8) S. Minakawa, M. Ohtake, D. Suzuki, N. Inaba, and M. Futamoto: J.
Magn. Soc. Jpn., 38, 185 (2014).
9) T. Soda, S. Minakawa, M. Ohtake, M. Futamoto, and N. Inaba:IEEE Trans. Magn., 51, 2000904 (2015)10) S. Minakawa, D. Suzuki, M. Ohtake, N. Inaba, and M. Futamoto:
EPJ Web Conf., 75, 01004 (2013).
11) G. Wasserman: Arch. Eisenhüttenwesen, 16, 647 (1933).
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13) G. Kurdjumov and G. Sachs: Z. Phys., 64, 325 (1930).
Received Dec. 28, 2015; Accepted Dec. 21, 2017

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J. Magn. Soc. Jpn., 42, 37-40 (2018)<Paper>

Theoretical study on magnetic tunneling junctions with semiconductor
barriers CuInSe2 and CuGaSe2 including a detailed analysis ofband-resolved transmittances
K. Masuda1 and Y. Miura1,2

1Research Center for Magnetic and Spintronic Materials, National Institute for Materials Science (NIMS),

1-2-1 Sengen, Tsukuba 305-0047, Japan

2Electrical Engineering and Electronics, Kyoto Institute of Technology, Kyoto 606-8585, Japan

We study spin-dependent transport properties in magnetic tunneling junctions (MTJs) with semiconductor
barriers, Fe/CuInSe2/Fe(001) and Fe/CuGaSe2/Fe(001). By analyzing their transmittances at zero bias voltage on thebasis of the first-principles calculations, we find that spin-dependent coherent tunneling transport of Δ1 wave
functions yields a relatively high magnetoresistance (MR) ratio in both the MTJs. We carry out a detailed analysisof the band-resolved transmittances in both the MTJs and find an absence of the selective transmission of Δ1 wave
functions in some energy regions a few eV away from the Fermi level due to small band gaps in CuInSe2 andCuGaSe2.
Key words: magnetic tunneling junctions, magnetoresistance ratios, semiconductors, ab initio calculations
1. Introduction
Both high magnetoresistance (MR) ratios and low
resistance area products (RA) are required for MR
devices to realize read heads of ultrahigh-density harddisk drives and Gbit-class spin transfer torque magnetic
random access memories (STT-MRAMs) 1). Althoughmagnetic tunneling junctions (MTJs) with insulator
MgO barriers have high MR ratios 2),3), these MTJs also
have high RA. To reduce the RA, many studies haveused ultrathin MgO barriers (~1nm), by which
relatively low RA ~ 1 Ωμm2 have been obtained 4)-6).
However, MR ratios also decrease by reducing the
thicknesses of the barriers. In addition, ultrathin
barriers do not have sufficient controllability infabrication
processes.
On
the
other
hand,
current-perpendicular-to-plane giant magnetoresistive(CPP-GMR) devices have quite low RA because metallic(not insulating) spacers are sandwiched between
ferromagnetic electrodes. Although the MR ratio in
CPP-GMR devices is increased up to 82% at room
temperature (RT) by using Co-based Heusler alloys aselectrodes 7), this value is still insufficient for
applications. Recently, Kasai et al. used a compound
semiconductor CuIn0.8Ga0.2Se2 (CIGS) as a barrier layerof MTJs to achieve high MR ratios and low RA 8). Theyobtained relatively high MR ratios ~ 40% at RT and low
RA ranging from 0.3 to 3 Ωμm2 in the MTJs where CIGS
is sandwiched between Co2FeGa0.5Ge0.5 (CFGG) Heusler
compounds. Since the barrier thicknesses of these MTJs
are twice as those of the above-mentioned MgO-based
MTJs with low RA, we can expect high controllability inthe fabrication processes.
In our previous work 9), we theoretically studied
spin-dependent transport properties of two MTJs withdifferent semiconductor barriers, CuInSe2 (CIS) andCuGaSe2 (CGS), to understand the origin of the highMR ratios observed in the CIGS-based MTJs 8). By
analyzing their complex band structures and k//
dependences of the transmittances, we found thatspin-dependent coherent transport of Δ1 wave functions
occurs in both the CIS- and CGS-based MTJs, which
can explain the high MR ratios in the CIGS-based MTJs.
In our present study, we carry out a more detailed
analysis of the CIS- and CGS-based MTJs to
understand their transport properties more deeply. We
provide the values of calculated conductances in both
the MTJs for both the parallel and antiparallel
magnetization cases of Fe electrodes. We also analyze
the band-resolved transmittances in these MTJs, fromwhich we find an absence of the selective transmissionof Δ1 wave functions in some energy regions a few eVaway from the Fermi level due to the small band gaps
in the CIS and CGS barriers.
2. Method
In this work, we analyzed CIS- and CGS-based MTJs
with Fe electrodes: Fe/CIS/Fe(001) and Fe/CGS/Fe(001).
Although Heusler alloys were used as electrodes in theexperiments on the CIGS-based MTJs, we selected Fe
electrodes to understand more simply the transport
properties in CIS- and CGS-based MTJs. Since thea-axis lengths of CIS (CGS) is close to twice that of bcc
Fe, the lattice mismatch between them is considered tobe small. We first prepared supercells of Fe/CIS/Fe(001)and Fe/CGS/Fe(001), each of which has two unit cells of
CIS (CGS) and one unit cell of Fe on both sides of the
CIS (CGS) barrier. The in-plane lattice constant a of thesupercell was fixed to 0.5782 nm (0.5614 nm), which isan a-axis length of the bulk CIS (CGS) 10). In both thesupercells, we optimized all the atomic positions andthe distance between the barrier and electrode using
the density-functional theory within the generalized

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37

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Table 1 Values of conductances per in-plane unit cell area (a2), RAP (RA in parallel magnetization states), andMR ratios in (a) Fe/CIS/Fe(001) and (b) Fe/CGS/Fe(001) for different Coulomb interactions U. The majority-spin(minority-spin) conductance in the parallel magnetization state is represented as GP,maj. (GP,min.), and the total-spin
conductance in the antiparallel magnetization state is represented as GAP,tot.
(a) Fe/CIS/Fe(001)
U [eV]

GP,maj.-spin [e2/h]GP,min.-spin [e2/h]

GAP,tot [e2/h]

RAP [Ωμm2]

MR ratio [%]01.923×10-2

9.609×10-3

1.885×10-2

0.299

52.9

5

1.662×10-2

4.540×10-3

1.303×10-2

0.408

62.3

10

1.427×10-2

2.500×10-3

9.322×10-3

0.515

79.9

U [eV]

GP,maj.-spin [e2/h]GP,min.-spin [e2/h]

GAP,tot [e2/h]

RAP [Ωμm2]

MR ratio [%]01.685×10-2

5.796×10-4

5.854×10-3

0.467

197.7

5

1.168×10-2

2.795×10-4

2.983×10-3

0.680

300.9

10

7.830×10-3

1.643×10-4

1.673×10-3

1.018

377.8

(b) Fe/CGS/Fe(001)gradient approximation implemented in the Vienna ab
initio simulation program (VASP) 11),12). In the processesof such structure optimizations, we found that the Se
layer is always energetically favored as the termination
layer of the barrier in both the CIS- and CGS-based
supercells. We studied transport properties of both the
MTJs at zero bias voltage with the help of the quantum
code ESPRESSO 13). This code was applied to thequantum open system composed of the above-mentioned
supercell attached to the left and right semi-infinite
electrodes of bcc Fe. In the present work, we considered
the Coulomb interaction U in the Cu 3d states in thebarrier. Since the band gap Eg of the barrier increases
as U increases, considering U is useful to systematicallystudy the relationship between the band gap and MR
ratio. The detailed calculation conditions are the sameas our previous study 9).
Due to the two-dimensional periodicity of our systems
in the plane parallel to the layers, the scattering states
can be classified by an in-plane wave vector k// = (kx,ky).
For each k// and spin index, we solved scattering
equations derived from connection conditions of the
wave functions and their derivatives at the boundaries
between the supercell and electrodes. Transmittances T
were obtained from the transmission amplitudes of thescattering wave functions 14),15). We obtained
conductances G by substituting the transmittances T
into the Landauer formula G = (e2/h) × T. In this work,
we adopted the usual definition of the optimistic MR
ratio: MR ratio (%) = 100 × (GP-GAP)/GAP, where GP (GAP)
is the sum of the majority- and minority-spintransmittances averaged over k// in the Brillouin zone in
the case of parallel (antiparallel) magnetization of Fe
electrodes. The values of RA in the parallel
magnetization states, RAP, were calculated using thein-plane lattice constants a of the supercells andconductances GP per in-plane unit cell area (a2) shownin Table 1.

38

3. Results and Discussion
Table 1 shows the values of conductances, RAP, andMR ratios in the CIS- and CGS-based MTJs fordifferent Coulomb interactions U. We obtained MR
ratios over 50% (190%) for the CIS-based (CGS-based)
MTJs. These relatively high MR ratios are due to thespin-dependent coherent transport of Δ1 wave functions,which was directly confirmed by the k// dependences of
the transmittances 9). We also analyzed complex band
structures of the CIS and CGS barriers 9), which also
indicated selective transmission of Δ1 wave functions.
The RA values in both the MTJs are much smaller thanthose of the well-known MgO-based MTJs 9), which
originates from the small band gaps in CIS and CGS
barriers.
From Table 1, we see that the MR ratio and RA
increase as the Coulomb interaction U increases inboth the MTJs. As we mentioned in the previous section,
the band gap Eg,CIS (Eg,CGS) of CIS (CGS) increases by
increasing U from 0 to 10 eV. Actually, we confirmedthat Eg,CIS = 0.31, 0.39, and 0.42 eV (Eg,CGS = 0.71, 0.84,
and 0.93 eV) for U = 0, 5, and 10 eV, respectively, by
analyzing the local density of states of Fe/CIS/Fe(001)and Fe/CGS/Fe(001) 16). Thus, it is found that asemiconductor barrier with a larger band gap gives ahigher MR ratio and a higher RA. This is also supportedby the fact that the CGS-based MTJ has a higher MR
ratio and a higher RA than the CIS-based one for thesame value of U (see Table 1). From the viewpoint oftunneling transport, such an increase in Eg corresponds
to an increase in the barrier height φ = Eg – EF, where EF
is the Fermi level of the system. Using theabove-mentioned band gaps (Eg,CIS and Eg,CGS) and the
Fermi levels in the Fe/CIS/Fe(001) and Fe/CGS/Fe(001),
we obtained the following values of barrier heights: φCIS
= 0.080, 0.100, and 0.111 eV (φCGS = 0.263, 0.301, and 0.330eV) for U = 0, 5, and 10 eV, respectively. Note here that thetransmittance through the barrier is approximately

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

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Fig. 1 Transmittances at k//=(0,0) in Fe/CGS/Fe(001) MTJ with U = 5 eV. (a) The energy dependence of thetotal transmittance and (b) the band-resolved transmittance in the [001] direction for the majority-spin channel
in the parallel magnetization state. (c) and (d) The same as (a) and (b) but for the minority-spin channel in theparallel magnetization state. (e) The energy dependence of the total transmittance in the antiparallel
magnetization state.
proportional to exp(-2κd), where κ = (2mφ/ℏ2 - |k//|2)1/2 and dis the barrier thickness 17). Therefore, increasing Eg (i.e.,
increasing φ) in the semiconductor barrier has roughly
the same effect as increasing the barrier thickness d,
both of which sharpen the tunneling feature andenhance the MR ratio.
In order to find characteristic features in thetransmittances of the CIS- and CGS-based MTJs, weanalyzed energy- and band-resolved transmittances atk//=(0,0) in the CGS-based MTJs 18). Here, we show theresults only for the CGS case, but qualitatively thesame results were also obtained in the CIS case.
Figures 1(a) and 1(c) show the energy-resolved
transmittances for majority-spin and minority-spinchannels, respectively, in the parallel magnetization
state of the electrodes. At the Fermi level (E=EF), thetransmittance in the majority-spin channel is morethan one order higher than that in the minority-spinchannel. In Fig. 1(e), we also show the energy-resolved
transmittance in the antiparallel magnetization state of
the electrodes. The transmittance at the Fermi level is
one order smaller than the majority-spin transmittancein the parallel magnetization case. From all thesebehaviors in the transmittances, we can understandnaturally the high MR ratio (~300%) in this system (seeTable 1). Figures 1(b) and 1(d) show the majority-spin
and minority-spin transmittances resolved into each(kz,E) contribution on the band structures of the Feelectrode. The energy-gap region is -0.539 eV ≦ E-EF ≦0.301 eV for the CGS with U = 5 eV in our calculations 19).
In the majority-spin channel [Fig. 1(b)], the Δ1 band
mainly from d3z2-r2 state gives high values oftransmittance in the energy-gap region. Such a selective
transmission of the Δ1 wave functions is consistent withsmall imaginary wave vectors in the Δ1 state obtainedin our previous work [see Fig. 2(b) of Ref. 9]. The
selective transmission of Δ1 wave functions also occurs

outside of the energy-gap region (0.3 eV ≲ E-EF ≲ 1.1 eV
and -1.0 eV ≲ E-EF ≲ -0.54 eV), which is because theconduction and valence bands of CGS in these regions
have Δ1 symmetry. In some other regions (e.g., a region
around E-EF = 1.5 eV and that around E-EF = 2.2 eV, etc.),
the Δ5 band mainly from px and py states gives higher
transmittances than the Δ1 band, which is in contrast to
the
case
of
the
band-insulator
barrier,
Fe/MgAl2O4/Fe(001), where the selective transmissionof the Δ1 wave functions persists to high energies [seeFig. 3(b) of Ref. 20 for comparison]. Note again thatsuch energy regions are outside of the band gap of the
CGS barrier. Thus, in these regions, tunnelingtransport is not possible and selective transmission of
Δ1 wave functions do not necessarily occur. We
confirmed that the conduction bands of CGS in theseregions include Δ5 component, which is consistent withthe high transmittances of Δ5 states (px and py states)
mentioned above. The similar feature can also be seen
in lower-energy regions with E-EF < -1 eV, where the Δ5and Δ2 bands from dyz (dzx) and dxy states give relatively
high transmittances. This is also supported by the factthat valence bands of CGS in these regions contain Δ5
and Δ2 components. Around the Fermi level in theminority-spin channel shown in Fig. 1(d), the Δ5 and Δ2
bands give similar values of transmittance as thed3z2-r2-based Δ1 band, i.e., clear selective transmission of
Δ1 wave functions does not occur unlike themajority-spin case. Since the in-plane lattice constant of
CGS is set to be twice that of the bcc Fe as mentioned
above, the original bands of bcc Fe is folded in the k//
Brillouin zone. Usually, in such a case, the folded
minority-spin Δ1 band crossing the Fermi level provides
relatively large transmittances and decreases the MR
ratio, as discussed in previous studies on
Fe/MgAl2O4/Fe(001) 21),22). However, in our present case,the folded minority-spin Δ1 band crossing the Fermi

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

39

INDEX

level does not give large transmittance compared to theother Δ5 and Δ2 bands and therefore does not decrease
MR ratio drastically. The configurations of interfacial
Se atoms in Fe/CGS/Fe(001) do not have fourfold
rotational symmetry along the c-axis of the supercell,which is clearly different from the configurations ofinterfacial O atoms in Fe/MgAl2O4/Fe(001). Thus, in thecase of Fe/CGS/Fe(001), although the bands areformally folded in the k// Brillouin zone, the folded
minority-spin band crossing the Fermi level has adifferent character from that of Fe/MgAl2O4/Fe(001),
which may be the reason for the ineffectiveness of the
band-folding effect in Fe/CGS/Fe(001).
The present analysis focused only on thetransmittances at zero bias voltage. On the other hand,in actual experiments, a small bias voltage is applied to
MTJs to obtain electron transmission. Since the bandgaps in the present study are small (Eg,CIS ≲ 0.42 eV and
Eg,CGS ≲ 0.93 eV), such a bias voltage gives a
non-negligible
spatialgradient
inthesemiconductor-barrier potentials, which can affect thetransport properties in the CIS- and CGS-based MTJs.
However, the small values of theoretical band gaps are
artifacts due to the generalized gradient approximation
used in this study. Although the present band gaps are
small, we adopted them as alternatives of experimental
band gaps of CIS and CGS (Eg,CIS = 1.0 eV and Eg,CGS = 1.7eV), on which the effect of the spatial gradient from thesmall bias voltage is not so large.
4. Summary
We investigated spin-dependent transport propertiesof magnetic tunneling junctions with semiconductor
barriers, Fe/CIS/Fe(001) and Fe/CGS/Fe(001). We
clarified that spin-dependent coherent tunnelingtransport of Δ1 wave functions leads to relatively highMR ratios in both the MTJs. From a detailed analysis of
the band-resolved transmittances in both the MTJs, wefound an absence of the selective transmission of Δ1wave functions in some energy regions a few eV awayfrom the Fermi level, which is a characteristic feature of
the MTJs with small band gaps in the barrier layers.
Acknowledgements The authors would like to thankK. Hono, S. Kasai, and K. Mukaiyama for giving ususeful information about experimental results on the
CIGS-based MTJs. KM also thanks K. Nawa forvaluable discussion on theoretical aspects. This work
was in part supported by Grants-in-Aid for Scientific
Research (S) (Grant No. 16H06332) and (B) (Grant No.
16H03852) from the Ministry of Education, Culture,
Sports, Science and Technology, Japan, by NIMS MI2I,
and also by the ImPACT Program of Council for Science,
Technology and Innovation, Japan.
References
1)40

B. Dieny, R. B. Goldfarb, and K. J. Lee: Introduction tomagnetic random-access memory (John Wiley & Sons,

2)

3)
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5)

6)

7)
8)

9)
10)
11)
12)
13)

14)
15)
16)

17)
18)

19)
20)
21)
22)

2016).
S. S. P. Parkin, C. Kaiser, A. Panchula, P. M. Rice, B.
Hughes, M. Samant, and S.-H. Yang: Nat. Mater. 3, 862(2004).
S. Yuasa, T. Nagahama, A. Fukushima, Y. Suzuki, and K.
Ando: Nat. Mater. 3, 868 (2004).
Y. S. Choi, Y. Nagamine, K. Tsunekawa, H. Maehara, D.
D. Djayaprawira, S. Yuasa, and K. Ando: Appl. Phys. Lett.
90, 012505 (2007).
S. Isogami, M. Tsunoda, K. Komagaki, K. Sunaga, Y.
Uehara, M. Sato, T. Miyajima, and M. Takahashi: Appl.
Phys. Lett. 93, 192109 (2008).
H. Maehara, K. Nishimura, Y. Nagamine, K. Tsunekawa,T. Seki, H. Kubota, A. Fukushima, K. Yakushiji, K. Ando,and S. Yuasa: Appl. Phys. Express 4, 033002 (2011).
J. W. Jung, Y. Sakuraba, T. T. Sasaki, Y. Miura, and K.
Hono: Appl. Phys. Lett. 108, 102408 (2016).
S. Kasai, Y. K. Takahashi, P.-H. Cheng, Ikhtiar, T.
Ohkubo, K. Kondou, Y. Otani, S. Mitani, and K. Hono:
Appl. Phys. Lett. 109, 032409 (2016).
K. Masuda and Y. Miura: Jpn. J. Appl. Phys. 56, 020306(2017).
J. E. Jaffe and A. Zunger: Phys. Rev. B 29, 1882 (1984).
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G. Kresse and D. Joubert: Phys. Rev. B 59, 1758 (1999).
S. Baroni, A. Dal Corso, S. de Gironcoli, and P. Giannozzi,
s: Quantum ESPRESSO package. For more information,
see http://www.pwscf.org.
H. J. Choi and J. Ihm: Phys. Rev. B 59, 2267 (1999).
A. Smogunov, A. Dal Corso, and E. Tosatti: Phys. Rev. B
70, 045417 (2004).
Note that these values of band gaps are different fromthose of bulk CIS and CGS shown in Ref. 9, which is dueto the hybridization between Se and Fe states at theinterfaces of Fe/CIS/Fe(001) and Fe/CGS/Fe(001)supercells. These band gaps are more suitable than thebulk ones for understanding the transport properties in
Fe/CIS/Fe(001) and Fe/CGS/Fe(001).
J. M. MacLaren, X.-G. Zhang, and W. H. Butler: Phys.
Rev. B 56, 11827 (1997).
Here, we focus only on the transmittances at k// = (0,0),
since these components provide dominant contributions to
the total transmittance. Note, however, that MR ratios
and RA are calculated from the transmittances integrated
over k// at the Fermi level (E=EF).
In the case of CIS with U = 5 eV, the energy-gap region is-0.290 eV ≦ E-EF ≦ 0.100 eV.
K. Masuda and Y. Miura: Phys. Rev. B 96, 054428 (2017).
Y. Miura, S. Muramoto, K. Abe, and M. Shirai: Phys. Rev.
B 86, 024426 (2012).
H. Sukegawa, Y. Miura, S. Muramoto, S. Mitani, T.
Niizeki, T. Ohkubo, K. Abe, M. Shirai, K. Inomata, and K.
Hono: Phys. Rev. B 86, 184401 (2012).

Received Oct. 18, 2017; Accepted Jan. 18, 2018

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

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J. Magn. Soc. Jpn., 42, 41-48 (2018)<Paper>

Evaluation of dispersibility in liquid and AC magnetization properties ofpolyion complex-coupled magnetic nanoparticles
Guannan Shi, Suko Bagus Trisnanto, Keita Nakai*, Shin-ichi Yusa*, Tsutomu Yamada,
Satoshi Ota**, and Yasushi Takemura
Department of Electrical and Computer Engineering, Yokohama National University,

79-5 Tokiwadai, Hodogaya-ku, Yokohama 240-8501, Japan*Department of Applied Chemistry, Graduate School of Engineering, University of Hyogo,

2167 Shosha, Himeji, Hyogo 671-2280, Japan**Department of Electrical and Electronic Engineering, Shizuoka University,

3-5-1 Johoku, Naka-ku, Hamamatsu 432-8501, Japan

Medical applications such as those using magnetic nanoparticles (MNPs) for hyperthermia and magnetic particle
imaging (MPI) require suitably designed particles with distinct characteristics. However, it is challenging to develop
such particles with a high degree of biocompatibility. In this study, a cationic diblock copolymer (PMPC 100-bMMAPTAC100: P100M100) composed of poly(2-(methacryloyloxy)ethyl phosphorylcholine) (PMPC) and poly(3(methacryloylamino)propyl trimethylammonium chloride) (PMAPTAC) was synthesized via a controlled radical
polymerization technique to obtain particles with high biocompatibility and antithrombogenicity. Magnetic polyion
complex (PIC) aggregate (M-300/P100M100) is an aggregate of magnetic Fe3O4 nanoparticles (M-300), in which theiranionic surface is electrostatically coated with cationic PMPC 100-b-MMAPTAC100 (P100M100). We investigated thestability of the magnetic PIC aggregate in an ionic solution by evaluating the relationship between the particle
diameter and salt concentration. We then estimated the intrinsic loss power (ILP) from the areas of the alternating
current (AC) hysteresis loops and measured the AC magnetization of the magnetic PIC aggregate. The peakfrequencies from the Brownian relaxation of M-300 and M-300/P100M100 were 9 kHz and 245 Hz, respectively. Whenthe third harmonic was evaluated for use in MPI, the signal intensity was found to be comparable to that of M-300 ina fixed state.
Keywords: magnetic nanoparticles, hyperthermia, magnetic particle imaging (MPI), biocompatibility, polyion
complex (PIC), magnetic relaxation
1. Introduction
In recent years, several medical treatment
applications have come to require magnetic
nanoparticles with high biocompatibility. Research into
magnetic nanoparticles coated with polyion complex(PIC) organic materials is ongoing due to their wideclinicalapplicability1–3),
and
the
sizes
and
characteristics of core-shell structures of PIC micelles
have rendered them useful as carriers of oligonucleotides
and plasmid DNA for human gene therapy4).
In the medical field, magnetic nanoparticles (MNPs)
are employed in a variety of applications including
magnetic hyperthermia5) and magnetic particle imaging(MPI)6,7). Magnetic relaxation loss is one of the heating
mechanisms of MNPs used in the research of
hyperthermia. In this regard, superparamagnetism in
MNPs results from two distinct magnetic relaxation
processes. The first occurs through a type of particle
rotation called Brownian relaxation, whereas the otheroccurs through a magnetic moment rotation referred toas Néel relaxation. The Brownian (τB) and Néel (τN)
relaxation times can be defined as follows8,9):

B 

3VH
,
k BT

(1)

 KVM , k BT  N   0 exp (2)

where η is the viscosity, VH is the hydrodynamic MNP
volume, 𝜏𝜏0 is the attempt time of ~10−9 s associatedwith gyromagnetic procession, K is the magnetic
anisotropic constant, VM is the core particle volume, kB is
Boltzmann’s constant (1.38×10−23 J/K), and T is thetemperature in Kelvin7). In theory, Brownian and Néel
relaxations occur in parallel, and the effective relaxation
time, τ, can be expressed as8):

1





1

B



1

N

.

(3)

The effective frequency for heat dissipation can be
estimated from measuring the dependency of the
intrinsic loss power (ILP) relative to the frequency. The
ILP is the intrinsic heat dissipation independent of theapplied field and is calculated from the specific loss
power (SLP)10), which is defined as:

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41

INDEX

Table 1. Concentration calculated from catalog data of M-300. Values in parentheses were used in calculations inthis study.
H2O

Fe3O4

RSO3NaM-300

Composition (wt%)

55–66 (55)

30–40 (40)5100Specific gravity (g/mL)

0.998 (1)

4.8–5.1 (5.1)

1

1.3–1.4 (1.4)

CAS No.

−1309-38-2

1106-21-0−

M.W. (g/mol)−

231.53298.42−Concentration (g/L)

−

560

70.0

−



ΔT c
,
Δt mFe(4)where ΔT/Δt is the time rate of change of temperature, cis the heat capacity of the fluid per unit mass of fluid,
and mFe is the iron mass in the fluid per unit mass of fluid.
Then, the ILP is given as10):
ILP 

SLP
,
H2 f

(5)

2. Materials and Methods
2.1 Materials
The characteristics of M-300 used for particle
synthesis in this study, along with its catalog data, areshown Table 1. As calculated from the catalog data, theconcentrations of Fe3O4 and RSO3Na in M-300 were 560
and 70.0 g/L, respectively. M-300 was subjected to
dynamic light scattering (DLS), and the zeta potential of
the solution diluted 100 times in pure water was

42

50

106

40

105

30

104

20

103

10

102

0
10-4

10-3

10-2

10-1

100

Fe3O4 concentration (g/L)
Fig. 1SI (kcps)

where H is the amplitude of the magnetic field.
In this study, we synthesized a diblock copolymer(PMPC100-b-PMAPTAC100:P100M100)
composedof
biocompatiblepoly(2-(methacryloyloxy)ethyl
phosphorylcholine) (PMPC) and cationic poly(3(methacryloylamino)propyl
trimethylammonium
chloride) (PMAPTAC) via a controlled radical
polymerization technique. PMPC shows excellent
biocompatibility and antithrombotic properties4) because
it has hydrophilic pendant phosphorylcholine groups,
which are structurally identical to the hydrophilic
surfaces of biological membranes. In our study, wecharacterized the chemical structure of the diblock
copolymer and evaluated the dependence of particle sizeon salt concentration to confirm its stability in an ionicsolution. We then measured the AC magnetization
properties of a magnetic PIC aggregate (M-300/P100M100)
comprising anionic Fe3O4 nanoparticles (M-300)
electrostatically coated with cationic P100M100. Finally,
we calculated the ILP and MPI signals from the AC
hysteresis loops and compared the results with those of
unmodified M-300.

measured (Fig. 1). After dilution, the concentrations of
Fe3O4 (CF) and RSO3Na (Cs) were 5.60 and 0.700 g/L,
respectively. In this case, the hydrodynamic radius (Rh)obtained from the DLS measurement was 21.2 nm andthe zeta potential was −52.0 mV. M-300 had a very low
scattering light intensity (SI) when diluted 1,000,000
times; although Rh was not calculated at this dilution
level, it remained constant to 100,000 times dilution (CF= 0.00560 g/L). In contrast, the zeta potential of the
10,000-fold diluted solution (CS = 0.00700 g/L) was −47.4
mV. Overall, it is clear that M-300 exhibited highstability, as evidenced by constant value of Rh versus
dilution in pure water.
PMPC100-CTA (Mn = 29800, Mw/Mn = 1.08 (acetate
buffer as the eluent), 1.50 g, 0.0503 mmol), MAPTAC(1.11 g, 5.03 mmol), and V-501 (4.59 mg, 0.0201 mmol)
were dissolved in 5.03 mL of pure water, and the mixture
was degassed by purging with Ar gas for 30 min.
Polymerization was performed at 343.15 K for 5 h, andthe diblock copolymer (PMPC100-block-PMAPTAC,
P100M100, 2.36 g, 90.4%) was recovered by freeze-drying
after the reaction mixture was dialyzed against purewater for 2 days. The M-300/P100M100 complex was
prepared by adding a P100M100 aqueous solution to an M300 aqueous solution so that the mixing ratio (R =P100M100/M-300, v/v) was 0.0739. The concentration of
Fe3O4 was 5.31 g/L (3.84 mg-Fe/mL) in the M-

Rh (nm)

SLP 

101
101



Dependence of Rh (open circles) and SI (opensquares) on concentration of Fe3O4 in waterbased M-300.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX

Table 2. Degree of polymerization (DP), number-average molecular weight (Mn), and molecular weight distribution(Mw/Mn) of P100M100.

DP(theo)aMn(theo)aDP(NMR)b

Mn(NMR)b

99

51,700100

51,900

Mw/𝑀𝑀𝑛𝑛𝑐𝑐
1.10

aTheoretical values estimated from the feed ratio and conversion. bEstimated from 1H NMR. cEstimated from GPC.

300/P100M100 complex. For comparison, we also diluted
water-based Fe3O4-containing sodium α-olefin sulfonatecoated nanoparticles, which are commerciallydistributed as M-300 by Sigma Hi-Chemical Inc.
(Chigasaki, Kanagawa, Japan), 100 times in pure waterto give a concentration of 5.31 g/L. The core diameter was
11 ± 3 nm, which is identical to that of M-300/P100M100.



2.2 Methods
The physical properties of M-300/P100M100 were
evaluated with nuclear magnetic resonance spectroscopy(NMR: Bruker, DRX500), dynamic light scattering (DLS),
zeta potential measurements (Malvern, Zetasizer
nanoZS), and transmission electron microscopy (TEM:
Jeol, JEM-2100). During this process, we also evaluatedthe NaCl concentration dependence of M-300/P100M100.
The AC hysteresis loops were measured in thefrequency range of 100 Hz–100 kHz at a field intensity



c
de

f

g

3. Results and Discussion

3.1 Synthesis, structure, and characterization
NMR spectroscopy was used to identify the chemical
structure of P100M100 synthesized by the above method,and it was confirmed that the peaks of PMPC and
PMAPTAC P100M100 were present (Fig. 2). On the basisof the consumption of the vinyl group before and after
polymerization, the degree of polymerization (DP (theo))of the monomer MAPTAC was determined to be 99. In
addition, we determined that the degree ofpolymerization (DP (NMR)) from the integrated
intensity ratio of the PMPC and PMAPTAC peaks in the
NMR spectrum of P100M100 after recovery was 100. The
molecular weight distribution (Mw/Mn) obtained from gelpermeation chromatography (GPC) was as narrow as
1.10, and a structure-controlled polymer was obtained(Table 2). We therefore concluded that P100M100 was
synthesized by the cationic monomer MAPTAC and thebetaine monomer MPC, as observed from the 1H NMR
spectrum. We also confirmed the mixing ratio

600

l

j m



a i

k

Rh (nm)waterof 4 kA/m, and the DC hysteresis loops were measured
using a vibrating sample magnetometer (VSM: Toei
Kogyo, TEM-WFR7) at 298.15 K. For thesemeasurements, 0.2 mL of the sample encapsulated incylindrical tubes with a diameter of 10 mm were used.
MNPs dispersed in water were used as the liquid sample,
and MNPs fixed with epoxy bonds were used as the solid
sample. The ILP and MPI values of the samples were
calculated, and the results were compared with those of
M-300.

bh

400
200

5.0 4.5 4.0 3.5 3.0 2.5 2.0 1.5 1.0

Chemical shift (ppm)Fig. 2

1H NMR spectrum measured for P100M100 in

D2O at room temperature. The symbols
indicate the states of 1H corresponding toeach peak of the spectrum.

0Fig. 3

0

0h
72 h

0.1

0.2

P100M100/M-300 (R)0.3Relationship of Rh with mixing ratio of M300/P100M100 immediately after preparation
and 72 h after preparation in pure water.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

43

INDEX

dependence of the M-300/P100M100 complex. The
concentration of M-300 anions was identical to theconcentration of the anionic surfactant sodium α-olefin
sulfonate in solution, although the exact concentrationof the anions is unknown. Therefore, in order to confirm
the neutralization point of the complex, we investigatedthe changes in Rh and the zeta potential of the aggregatewith respect to the mixing ratio of M-300 and P100 M100
in the complex. Figure 3 plots the Rh values obtained by
DLS measurement versus the mixing ratio (R = P100M100/ M-300, v/v) of M-300 and P100 M100 in pure water after
72 h. When P100M100 was mixed with M-300, Rh increased
as the value of R increased until it reached a maximum
at R = 0.0474. Increasing R beyond this point decreased
both Rh and SI. In addition, in M-300/P100M100 with R =0.0474, the SI decreased sharply after 72 h, suggestingthat precipitation had occurred.

20

ζ (mV)

0

R = 0.0739

-20
-40
-60

Fig. 4



0

0.1

0.2

0.3

P100M100/M-300 (R)Zeta potential for mixing ratio (R =P100M100/M-300, v/v) of water-based M300/P100M100.

By comparison, we concluded that the aggregates werestable since M-300/P100M100, other than that for R =0.0474, neither precipitated nor demonstrated a change
in Rh over the course of 72 h. Figure 4 shows the zeta
potential versus the mixing ratio of the synthesized M300/P100M100 after 96 h. The M-300/P100M100 with R =0.0474 precipitated; thus, the zeta potential of thesupernatant was measured. As the R value of M300/P100M100 increased, the zeta potential also increased
as the solution concentration of P100M100 increased.
Previously, it was thought that the charge would
neutralize with a zeta potential of zero for M300/P100M100 with R = 0.0739. It was also thought thatlarger absolute values of the zeta potential indicatedstronger repulsive forces between the particles andhence higher particle stability. However, the results ofthis study indicate that the surface charges were
neutralized to completion, and they were stably
dispersed due to steric hindrance by the polymer
material outside the particles. In this case, the Na+
cations from sodium tetradecane sulfonate and the Clˉanions from P100M100 were ionically bonded in water,while the N+ cations from P100M100 and the Oˉ anionsfrom tetradecane sulfonate were electrostatically bonded.
In this way, the magnetic PIC aggregate was synthesized
as shown in Fig. 5.
Moreover, we confirmed that the surface of the M300/P100M100 complex (R = 0.0739) was covered by a
PMPC shell based on the dispersion stability (i.e., the Rh
and SI) versus the NaCl concentration in solution (Fig.
6). The Rh and SI values of M-300/P100M100 wereindependent of the NaCl concentration ([NaCl]); however,the Rh and SI values of M-300 increased when [NaCl]
was increased to 0.2 M (mol/L). These results indicatethat M-300/P100M100 was covered by a PMPC shell since
the dispersion stability improved upon complexation,
which was synthesized with P100M100. Further,aggregates appeared in M-300 for [NaCl] = 0.2 M whenleft overnight. In contrast, M-300/P100M100 did notprecipitate but was maintained as a stable dispersion.
For a NaCl concentration of 0.2 M, specific quantities of
M-300/P100M100 and M-300 aqueous solutions were

80

1

40

0.5

0


Fig. 544

Schematic representation of M300/P100M100 preparation.

1.5

M-300/P100M100
M-300

Fig. 6

0

0.05

0.1

0.15

[NaCl] (M)SI×10-5 (kcps)

Rh (nm)

120

0
0.2



Change in Rh and SI with respect to NaCl
concentration of M-300 (open squares) and
M-300/P100M100 (open circles).



Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX

(a)

(b)

(c)

200 nm50 nm
Fig. 7

100 nm

TEM images of (a) M-300 and (b, c) M-300/P100M100. From TEM image, it was confirmed that M-300 and
M-300/P100M100 are aggregates of Fe3O4.

allowed to stand still for one week to confirm the changes
over time. The Rh value of M-300/P100M100 was 63.2 nm
and remained unchanged over time, whereas that of M300 increased to 144 nm. Furthermore, since M-300
precipitated by salting-out, we concluded that the PMPC
shell of the complex inhibited this process. Based onthese results, we inferred that as the sodium
concentration in the solvent increased, the sodium ionson the M-300 surface ionically bonded with the chlorine
ions, resulting in aggregation. However, the magnetic
PIC aggregate surface was not affected by the salt
concentration since it was charge-neutralized. From thisproperty, we conclude that magnetic PIC aggregates canbe stably dispersed in biological bodies.
Figure 7 shows the TEM images of M-300 and M300/P100M100. On this basis, we confirmed that M-300

was an aggregate of Fe3O4 (Fig. 7 (a)) and that M300/P100M100 was the same aggregate as M-300 (Fig. 7 (b,c)). Therefore, we concluded that M-300/P100M100 was
packed with anionic M-300 by the cationic diblock
copolymer P100M100, and the estimated radius of M300/P100M100 was ~50 nm. Moreover, the hydrodynamic
radii of M-300 and M-300/P100M100 were 20.2 nm and47.8 nm, respectively, as determined by DLS
measurements (Fig. 8). The characteristics of M-300 and
M-300/P100M100 are listed in Table 3.



3.2 Magnetic properties
Figure 9 shows the DC hysteresis loops of M-300 and
M-300/P100M100 for the fixed and liquid samples at field
intensities of 0–800 kA/m. We also measured the DC

1.2

Table 3. Characterization of M-300 and
M-300/P100M100.

0.6
SI (kcps)

PDI

M-300

20.2

4.21 × 105

0.120

M-300/P100M100

47.8

1.70 × 106

0.119

0
Magnetization [M/Ms ]Rh (nm)

20.2 nm

-1.2
1.2
0.6

M-300
M-300/P

100

M

(b)
100

0

47.8 nm

101

-1.2
-1000

102

Rh (nm)103

Hydrodynamic radii of M-300 and M300/P100M100.

DC
Liquid800 kA/m-0.6

M-300/P100M100

Fig. 8DC
Fixed
800 kA/m

-0.6

M-300

100

(a)

M-300
M-300/P100M100

Fig. 9-500
0500Magnetic Field [kA/m]

1000


DC hysteresis loops of (a) fixed sample and(b) liquid sample at magnetic field
amplitude of 0–800 kA/m. Concentration of
Fe was 3.84 mg/mL.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

45

INDEX

0.2

100

M

0.2

100

-0.1
-0.2
0.3
0.15

DC
Fixed
4 kA/m(b)

M-300
M-300/P

100

M

100

M-300

-0.2
-0.4
0.2
0.1

-0.15
0

3

-0.2
-6

6


DC hysteresis loops of (a) fixed sample and(b) liquid sample in magnetic field
amplitudes ranging from 0–4 kA/m.
Concentration of Fe was 3.84 mg/mL.

Fig. 11M-300/P100M100
Liquid4 kA/m
3.84 mg-Fe/ml-3
0
3Magnetic Field [kA/m]

6

AC hysteresis loops at 1, 10, and 100 kHz
for M-300 and M-300/P100M100.

hydrodynamic particle size in M-300/P100M100 decreased
the degree of particle rotation because of dipole
interactions11).
Figure 13 (a) shows the dependence of ILP on
frequency in the liquid sample, as estimated from theareas of the AC hysteresis loops. Note that the observed
ILP peaks derived from Brownian relaxation did notappear for the fixed sample (Fig. 13 (b))12). Furthermore,
we calculated the theoretical Brownian relaxation timeusing Eq. (1), and determined the viscosities of M-300
and M-300/P100M100 at 297.15 K as 1.081 mPa∙s and1.037 mPa∙s, respectively. The experimental Brownian

0.4
Magnetization [M/Ms ]-3hysteresis loops for field intensities of 0–4 kA/m to better
understand the magnetization process (Fig. 10). As
observed from the results of DC measurements, the
MNPs exhibited superparamagnetic characteristics
because the coercive force and remanence magnetization
could not be confirmed from the hysteresis loops. In thiscase, the fixed samples had the same magnetization
because only the Néel relaxation from the magnetic
moment rotation of the core particles operated in the
fixed state. However, the magnetization of M-300 was
higher than that of M-300/P100M100 in the liquid state.
We concluded that the action of Brownian relaxation
caused by particle rotation was very weak in M300/P100M100 because their Rh value was larger than thatof M-300.
Figure 11 shows the dependence of the AC hysteresis
loops on frequency at 3.84 mg-Fe/mL for the fixed andliquid samples. As seen, the magnetization decreased as
the frequency increased because the particle rotationand magnetic moments gradually became delayed. We
also confirmed that the magnetization of M-300/P100M100
was lower than that of M-300 in the liquid samples
under an AC magnetic field of 10 kHz (Fig. 12). Based onthese results, we concluded that the influence ofBrownian relaxation was small in the liquid state of themagnetic PIC aggregate due to the difference in thehydrodynamic particle size in M-300. The increase of the

46

DC
1 kHz
10 kHz
100 kHz-0.1

DC
Liquid4 kA/mMagnetic Field [kA/m]
Fig. 10Liquid4 kA/m
3.84 mg-Fe/ml

0

0

-0.3
-6

DC
1 kHz
10 kHz
100 kHz00

Magnetization [M/Ms ]Magnetization [M/Ms ]

0.1

0.4

(a)

M-300
M-300/P

0.2

M-300
M-300/P

M

100

100

0
Liquid4 kA/m
10 kHz

-0.2
-0.4
-6

Fig. 12-303Magnetic Field [kA/m]

6

AC hysteresis loops of liquid samples inapplied magnetic field with amplitude of 4
kA/m. Frequency was 10 kHz.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018



INDEX

Theoretical f B Theoretical f B
Experimental f B
Experimental f B

4

M-300
M-300/P

3

M

100

0.04

M-300/P100M100

0.01
0

0.01

M-300
M-300/P

M

100

100

Fixed
4 kA/m

0.005

0

1

10
Frequency [kHz]

100





Fig. 14 Dependence of the third-harmonic signal on
frequency in liquid and fixed samples.

Liquid
100

0.1

M-300
M-300/P

M

M-300
M-300/P

M

100

100

Liquid4 kA/m

1

0.05

0
4

Fixed

M-300
M-300/P100M100

M3/M1

3
2

0
0.1

1

100

0
0.1

Fig. 13

Liquid4 kA/m0.02

2

ILP [nHm2/kg]

M-300

0.03

M3 [M/Ms ]relaxation frequencies of M-300 and M-300/P100M100
were 9.0 kHz and 250 Hz, respectively. The measured
frequency of the Brownian relaxation peaks was in good
agreement with the calculated frequency. The frequencyof the Brownian relaxation peak in M-300 was
significantly higher than that in M-300/P100M100. Thus,
particle rotation derived from Brownian relaxation,
which increases the degree of the magnetization rotation,clearly occurred in M-300 relative to M-300/P100M100 13,
14). A higher ILP was obtained for M-300 in the liquidstate, whereas a comparable ILP was obtained in the
fixed state. The frequency of the Néel relaxation peak, fN,
was not observed in the range of 0.1–100 kHz for the
fixed samples. Note that the theoretical Néel relaxation
peak was higher than 450 kHz [K ≤ 41 kJ/m3]15).
Finally, we assessed the suitability of the magnetic
PIC aggregate for applications in MPI. Figures 14 and
15 show the third-harmonic intensities of M-300 and M300/P100M100. The third harmonic signal intensity of M300 decreased with increasing frequency. We concluded
that this was due to a decrease in magnetization caused
by magnetic relaxation16). In contrast, the magnetic PIC
aggregate changed very little under the same conditions.
However, it was confirmed that M-300 and M-

1
10Frequency [kHz]

0.05

100

Dependence of ILP on frequency for theliquid and fixed samples. Dotted lines
indicate theoretical Brownian relaxation
frequencies of fB = 5.8 kHz (M-300) and fB= 460 Hz (M-300/P100M100). Solid lines
indicate
experimental
Brownianrelaxation frequencies of fB = 9.0 kHz (M300) and fB = 250 Hz (M-300/P100M100). No
Néel relaxation peaks were observed.

100

Fixed
4 kA/m0

1

10Frequency [kHz]

100




Fig. 15

Dependence of ratio of the third-harmonic
signal to fundamental signal on frequency
in liquid and fixed samples.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

47

INDEX

300/P100M100 had comparable MPI signal intensities in
the fixed state, whereas stronger signal intensities wereobtained for M-300 in the liquid state.
The harmonic signal used for MPI is normallyconcerned with the nonlinear component ofmagnetization property. The difference in this nonlinear
component of magnetization with and without anapplied DC bias field is significant for spatial resolutionof the constructed image. As shown in Figs. 10 and 11,
the nonlinear component of the magnetization loops is
not very large. Although the surface modification using
P100M100 was successfully achieved, optimization of themagnetic core is inevitable to lead to a higher ILP andharmonic signal.

Reference
1)
2)
3)

4)
5)
6)

4. Conclusion

7)

In this study, we successfully fabricated P100M100 witha biological membrane structure and evaluated the salt
concentration dependence of M-300/P100M100 whencoated with P100M100. When the mixing ratio was 0.0739,
the surface of the synthesized particles was neutralized
by the charge, and the magnetic PIC aggregate wasstable and dispersible in pure water and ionic solutions.
Based on this result, we infer that the magnetic PIC
aggregates can be stably dispersed in a biological body.
The results of magnetization measurements allowed us
to characterize the influence of different hydrodynamic
radii on the magnetization. As the hydrodynamic
particle size of the magnetic PIC aggregate increased,
the influence of Brownian relaxation decreased due tomagnetic interactions. Further, the ILP was measured toinvestigate the applicability of M-300 and M300/P100M100 to hyperthermia. The particles had
comparable ILP values in the fixed state, but that of M300 was higher in the liquid state. In addition,comparable MPI signal intensities were confirmed in the
fixed state by evaluating the third harmonic.

8)

Acknowledgments This work was partially supportedby the JSPS KAKENHI Grant Numbers 15H05764,
17H03275, and 17K14693.

48

9)
10)

11)

12)
13)
14)
15)
16)

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Received Oct. 13, 2017; Accepted Dec. 21, 2017



Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX
J. Magn. Soc. Jpn., 42, 49-54 (2018)<Paper>Power supply for medical implants by Wiegand pulse generated
from magnetic wire
Katsuki Takahashi, Akitoshi Takebuchi, Tsutomu Yamada, and Yasushi Takemura
Department of Electrical and Computer Engineering, Yokohama National University,

79-5 Tokiwadai, Hodogaya-ku, Yokohama 240-8501, Japan

Implantable medical devices are utilized in the human body for maintaining good health. As these devices are
becoming increasingly functionalized, supplying power to them has become very important. Instead of batteries,
technologies for wireless power supplies are being developed, such as inductive coupling or piezoelectric elements.
However, we propose the use of magnetic wires for this purpose. Inside these wires, a fast magnetization reversal,
called a “large Barkhausen jump,” occurs due to an applied magnetic field. This reversal induces a large pulse voltage
in pick-up coils, which is called a “Wiegand pulse.” By applying this pulse as an electric source, a higher electric poweris expected compared with the conventional method using a sinusoidal excitation field. A wire core coil was prepared,
and open-circuit voltage was measured. In addition, DC voltage and electric power were measured by connecting thewire core coil to a rectifier. The experimental results confirmed the superiority of using a Wiegand pulse at an appliedmagnetic field intensity of 60 Oe and a frequency of lower than 10 kHz.
Keywords: magnetization reversal, large Barkhausen jump, FeCoV wire, Wiegand pulse, implantable medical device
1. Introduction
A large Barkhausen jump, which is known as the
Wiegand effect1), is a fast magnetization reversal in amagnetic wire that has bistable states. It induces a large
pulse voltage, called a Wiegand pulse, in the pick-up coilwhen a magnetic field is applied. In this study, a twisted
FeCoV wire made from an optimum material yielding
this effect was used. The Wiegand effect has certain
advantages: it does not need any external power supply,
and the amplitude of the pulse voltage is still large atboth high and low frequencies, while an ordinary
sinusoidal electromotive force depends on the frequencyof the applied magnetic field. This phenomenon has beenused in many sensor applications, such as rotation
sensors and speed sensors2,3). Further, vibration-type
energy-generating elements by exciting the magnetic
wire through a single magnet are being developed4,5).
In this study, a magnetic wire and Wiegand pulse wereused as the power source, which provides DC voltage tothe implant inside the human body. Ordinary inductive
power transfer systems typically use excitation
frequency ranging from hundreds of kilohertz to several
megahertz. At such high frequency, the excitation
voltage, current, and size of the coil tend to be too large
for applying enough magnetic field to deep the humanbody. Further, unexpected heating of the body surface
caused by the eddy currents cannot be ignored. To
eliminate these problems, we proposed the use of amagnetic wire. Using the wire as the core of the pick-up
coil and generating a large Wiegand pulse by applying amagnetic field, enough electric power can be achieved at
low excitation frequency. Then, reduction of not only thesize of both excitation coil and pick-up coil, but also therequired excitation voltage, can be realized.
2. Structure of Electricity Generation
2.1 FeCoV wire
An 11-mm-long FeCoV (Fe0.4Co0.5V0.1) wire was used inthis study. When torsion stress was applied to the wire,
the outer shell near the surface became magnetically soft.
When the stress was released, two layers appeared: theouter layer, called soft layer, which had a coercive forceof 20 Oe, and the inner layer, called the hard core, which
had a coercive force of 80 Oe. Details on the magnetic
properties of twisted FeCoV wires and torsion stressdependence have been reported by Abe et al.6,7). Thismagnetic wire exhibits a uniaxial magnetic anisotropy
along its length direction.
Figure 1 shows schematic diagrams of magnetic
structure and magnetization process of the soft layer in
the FeCoV wire used in this study. These characteristics
Antiparallelalignment

Parallel alignment

M

(A)

(B)

H
±20 Oe(C)

Hard core(D)
Soft layer

Fig. 1 Schematic diagrams of magnetic structure andminor hysteresis loop of FeCoV wire.

Journal of the Magnetics Society of Japan Vol.42, No.2, 201849

Output voltage [mV/turn]Output voltage [mV/turn]

50

5

0

0

-50

-5

-100
-500

-250

0
250
Time [ μs]

-10
500

Magnetic field intensity [Oe](a) Excitation frequency of 1 kHz.

100

40

50

20

0

0

-50

-20

-100
-50

-25

0
Time [ μs]

25

-40
50

(b) Excitation frequency of 10 kHz.
Fig. 3 Waveforms of applied AC magnetic field of 60 Oe
and induced voltage of wire core coil at field
frequencies of (a)1 kHz and (b)10 kHz.

0.6

observed in one period, each of which results from tworeversals of the soft layer: parallel to antiparallel andantiparallel to parallel. In the 1 kHz waveform shown in
Fig. 3(a), only a Wiegand pulse is mostly observed, whilein the 10 kHz waveform shown in Fig. 3(b), sinusoidal
components can be slightly seen, but are still much lessdominant compared to a Wiegand pulse.
In this study, a voltage-double rectifier was used, anda Wiegand pulse was converted to DC voltage assuming
the electric power supplied to capsule endoscopy or manyother implantable medical devices. Enough electric
power at a low frequency was aimed using a Wiegand
pulse generation. The excitation frequency was set tolower than 10 kHz in all experiments.

0.4

3. Experiment

1
0.8

0.2
0

0

50

100 150
TimeTime [s]
[μs]

200

250

Fig. 2 Wiegand pulse generated by magnet excitation.

50

10Voltage [V]2.2 Wiegand pulse in wire core coil
By using the wire as a coil core and applying amagnetic field larger than 20 Oe, a Wiegand pulse isinduced in the pick-up coil wound around the wire.
Figure 2 shows a waveform of the measured Wiegand
pulse during magnetization reversal of the soft layer
excited by approaching a magnet in extremely slow
movement. When the wire is excited by an AC magnetic
field, sinusoidal electromotive force is also induced in the
pick-up coil. Therefore, the waveform of the wire core coilby AC excitation combines these two waveforms, as
shown in Fig. 3. The pick-up coil of 3000 turn was usedin this measurement. Unlike the sinusoidal
electromotive force, the width of a Wiegand pulse does
not vary according to the excitation frequency. As seen in
the figure, positive and negative Wiegand pulses are100

Voltage [V]are essentially same with those reported in Ref. 5), in
which both of measured full and minor loops have beenshown. The magnetization alignment of the soft layer
and the hard core can be either in a parallel state or an
antiparallel state, as shown in Fig. 1. The magnetizationalignment takes a parallel state when an external
magnetic field of larger than 20 Oe is applied in the right
direction along its length, as shown in Fig. 1(A). When amagnetic field with an opposite direction to that ofmagnetization is applied and reaches −20 Oe, themagnetization of the soft layer is reversed steeply andtakes an antiparallel state, as indicated in Fig. 1(B). Thismagnetization reversal, called the large Barkhausen
jump, is due to a fast domain wall propagation. Thereby,a large pulse voltage, called a Wiegand pulse, is inducedin the pick-up coil around the wire. The strength ofmagnetization of the soft layer along the left direction
increases with the magnetic field applied in the left
direction from −20 Oe at (B). Then, the direction of themagnetic field switches at (C), and the wire starts
receiving the right magnetic field. On reaching 20 Oe,
magnetization reversal again occurs rapidly in the soft
layer, shown by (D) (from antiparallel state to parallelstate), and a Wiegand pulse is induced. If the maximum
value of the applied field intensity is within 80 Oe, themagnetization of the hard core does not reverse.

Magnetic field intensity [Oe]INDEX

3.1 Comparison of open-circuit voltage in wire and air
core coil
A pick-up coil of 9 mm length and 3000 turns waswound around the FeCoV wire (length: 11 mm). The
diameter of the wire was a few tenth of a millimeter. Its
magnetic properties were essentially same with those of

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX

Vp

Vopen(a) Configuration of wire core coil and measurementof Wiegand pulse.

Vopen

Vp

(b) Configuration of air core coil and measurement ofsinusoidal electromotive force.

Fig. 4 Open-curcuit voltage of wire core coil (Wiegand
pulse) and air core coil (sinusoidal electromotive
force).

with that of the air coil. In the measurement of magnetic
field intensity dependence, the field intensity ranged
from 0 Oe to 100 Oe and the excitation frequency was set
to 1 kHz. In the measurement of excitation frequency
dependence, the frequency ranged from 100 Hz to 10 kHz
and the field intensity was set to 60 Oe. As shown in Fig.
4, the peak of the Wiegand pulse was measured in thewire core coil, and the amplitude of the sinusoidal
electromotive force was measured in the air coil.
3.2 Measurement of electric power obtained from
Wiegand pulse
As shown in Fig. 5(a), the wire core pick-up coil wasconnected to a voltage-double rectifier, which consistedof two Schottky diodes and 1 μF capacitors and a load
resistor. The excitation coil was arranged around thepick-up coil, and the size and number of turns of eachcoil were the same as those mentioned in the previous
section. The Wiegand pulse was rectified and smoothedas shown in Fig. 5(b). Then, the load voltage and its
ripple components were measured with the resistancevalue varying from 100  to 1 M, and electric power
was calculated. Intensity of the AC magnetic field for
excitation was set to 60 Oe, and two frequencies wereused: 1 kHz and 10 kHz. The ripple rate was calculatedusing the following equation.

Ripple 

RL

VLoad

(a) Configuration of wire core coil, excitation coil,voltage rectifier circuit, and load resistor.

(b) Process through which Wiegand pulse is rectified
and smoothed by circuit above.

Fig. 5 Conversion of Wiegand pulse into DC voltage:
configuration and process.
the FeCoV wire of 0.25 mm diameter4,5). The volumes ofthe soft and hard core were 23% and 77%, respectively.
The excitation coil, with 25 mm length and 22 mm
diameter, applied the AC magnetic field to the entire
wire core coil. Each end of the core coil was connected toan oscilloscope and the peak of the open-circuit voltage
was measured; subsequently, this peak was compared

Vmax  Vmin
100
Vave

(1)

3.3 Comparison of electric power obtained from the wireand ferrite core coil
To confirm the usefulness of the Wiegand pulse at low
frequency, a high-permeability magnetic material, MnZnferrite, was used for the coil core and compared with thewire core coil. The ferrite core strengthened the magnetic
flux around the pick-up coil, and the sinusoidal
electromotive force itself increased. The size of the ferrite
core used in this experiment was 1 × 1 × 10 mm 3. The
ferrite core and the wire were each wound by a pick-up
coil of 400 turns. Each core coil was connected to thevoltage-double rectifier and load resistor, as described in
the previous section. Then, the load voltage wasmeasured and electric power generated from the wire
and the ferrite core was calculated and compared.
Because the ferrite core and wire had different sizes, themeasured data were normalized with their crosssectional area. Load resistance was 1 M andcapacitance was 1 μF. The excitation magnetic field
intensity was set to 60 Oe, which is enough to generatea Wiegand pulse. The excitation frequency ranged from100 Hz to 10 kHz.
4. Results and Discussion
4.1 Comparison of open-circuit voltage with changing
applied field intensity or frequency
Figure 6 shows the dependence of open-circuit voltage

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

51

INDEXon magnetic field intensity at 1 kHz of excitation
frequency and that of excitation frequency on 60 Oe ofapplied field intensity.
As shown in Fig. 6(a), when the excitation field was 20
Oe or more, the peak value of the wire core was far larger
than that of the air coil. This implies that the Wiegand
pulse was generated due to magnetization reversal in
the soft layer; its peak value increased with the appliedfield intensity. When the applied field was within 20 Oe,
because the Wiegand pulse was not generated, ordinary
electromotive force was measured. At an applied field of
10 Oe, the peak value of the wire core was 84.2 mV,
which was slightly larger than that of the air core (28.4mV). It was found that because of the magnetic
permeability of the wire, magnetic flux around the pickup coil strengthened, and the electromotive force
increased to a certain degree. The magnetization of thehard core was also reversed by the applied field above 80
Oe, which resulted in lower output voltage from thewire6). The voltage was expected to be saturated or moredecreasing even though the magnetic field exceeded 100
Oe. Magnetization reversal of the hard core is notnecessary for generating the Wiegand pulse, and thefield intensity should not be excessively strong. Then,the ideal intensity for generating the Wiegand pulse isaround 60–80 Oe. Therefore, in this study, the appliedfield was set to 60 Oe, which fixed the field intensity.
In Fig. 6(b), at all measured frequencies, the peakvalue of the Wiegand pulse was far larger than theelectromotive force of the air coil. As the frequency
increased, the amplitude of the wire core coil also
increased. We assumed that the sinusoidal component
increased at higher frequency, and then, the peak valueof the Wiegand pulse was increased.
From these results, higher electric power is expected to
be secured at lower excitation frequency during
conversion to DC voltage.
4.2 Electric power obtained from wire core coilFigure 7 shows the results of the measured voltage

Wire core8

Air core (x10)Power [mW]

10
6
4
2
0

1 kHz

0

20
40
60
80
100
Magnetic field intensity [Oe](a) Dependence of 1 kHz of excitation frequency on field
intensity.

1.2

10

1

8

0.8

0
1x102

1x103 1x104 1x105
Resistance []

0
1x106

(a) DC voltage and electric power with change inresistance.

25

Wire core20

Air core (x10)120
Ripple [%]Voltage [V]

2

0.2

150

15
10
5

60 Oe

1
Frequency [kHz]10(b) Dependence of 60 Oe of applied field intensity on
excitation frequency.
Fig. 6 Dependence of open-curcuit voltage of wire core
coil and air core coil on applied field intensity andexcitation frequency.

52

4

0.4

30

0
0.1

6

0.6

Voltage [V]Voltage [V]

12

90
60
30
0
1x102

1x103 1x104 1x105
Resistance []1x106(b) Ripple rate of load voltage.
Fig. 7 Dependence of DC voltage (average value) andelectric power on resistance and its ripple rate,
at 60 Oe of applied field intensity and 1 kHz of
excitation frequency.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

25

50

20

40

15

30

10

20

5

10

0
1x102

0
1x106

DC voltage and electric power with change inresistance.

25

15

40

10
5
0
1x102
(b)

1x103 1x104 1x105
Resistance []1x106

Ripple rate of load voltage.

Fig. 8 Dependence of DC voltage (average value) andelectric power on resistance, and its ripple rate
at 60 Oe of applied field intensity, and 10 kHz of
excitation frequency.
and electric power of load resistance (a) at 1 kHz
excitation frequency, and its ripple (b). The maximum
obtained electric power was 1.2 mW. The ripple rate
deteriorated at lower resistances. In the case where thewire core coil receives power from the excitation coil andprovides voltage to the implant device throughout
operation in the human body, the fluctuation in the load
voltage should be less, and thus, the capacitor should be
optimized.
Figure 8 shows same results as Fig. 7 but at an
excitation frequency of 10 kHz. The maximum electric
power was approximately 21 mW. Under this excitation
condition, there is enough possibility of feeding thepower required to operate implantable medical devices(for example, power consumption of capsule endoscopy is
25 mW8), and is expected to lessen in the future). Further,
the load voltage at the highest electric power was 1.56 V
at 1 kHz and 8 V at 10 kHz excitation. Considering thatthe voltage of a button battery, the main power source of
implantable medical devices, is around 1.5 V, DC voltage
from the Wiegand pulse was also sufficient.

Power [ μW/mm2]

Ripple [%]20

4.3 Comparison of electric power generated from wire
core and ferrite core
Figure 9 shows a comparison of the electric power
generated from the wire core coil and the ferrite core coil.
The excitation frequency was 1 kHz. The ripple almost
disappeared because the circuit parameters (capacitanceand load resistance) and the time constant were large
enough. Electric power from the Wiegand pulse
increased significantly from the applied field intensity of
20 Oe or more. Even considering the demagnetizing field
inside the ferrite, the power from the ferrite core coilremained smaller at almost all intensity values. The
tendency of the wire core coil resembles that in Fig. 6(a),in that the peak of the applied field intensity was in therange of 70–80 Oe.
Figure 10 shows a comparison of the electric power
generated from the wire core and ferrite core withchanging excitation frequency. In frequencies lower than
10 kHz, the wire core coil had an advantage. Thus, theWire core30

Ferrite core

20

1 kHz

10
0

0

20
40
60
80
100
Magnetic field intensity [Oe]Fig. 9 Comparison of electric power generated from wire
core and ferrite wire coil with change in appliedfield intensity at 1 kHz.

Power [ μ W/mm2](a)1x103 1x104 1x105
Resistance []Voltage [V]Power [mW]

INDEX

104
103
102
101
100
10-1
10-2
10-3
0.1

Wire core
Ferrite core

60 Oe

1
Frequency [kHz]10

Fig. 10 Comparison of electric power generated fromwire core and ferrite wire coil with change inexcitation frequecncy at 60 Oe.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

53

INDEX

usefulness of the Wiegand pulse and wire core coil atlower excitation frequency was confirmed. If theexcitation frequency is increased beyond 10 kHz, thesinusoidal electromotive force will increase, and theferrite core coil may become superior to the wire core coil.
5. Conclusion
This paper proposes a new method of power supply
using an 11-mm-long FeCoV wire. A Wiegand pulse was
induced in a pick-up coil wound around the wire by an
AC magnetic field of the excitation coil. A comparison of
the open-circuit voltage with the air core coil showedthat the Wiegand pulse has much larger peak than thesinusoidal electromotive force at frequencies lower than
10 kHz. Subsequently, the Wiegand pulse was rectified
and smoothed by a voltage-double rectifier, and it was
confirmed that the wire core coil can provide enoughelectric power to implantable medical devices. Further,
the wire core coil was compared to the ferrite core coil,which had high permeability, and higher electric power
was obtained by the Wiegand pulse, even though theelectromotive force was amplified by the ferrite core at
the measured frequency. In addition, it was found that
reduction in the size of pick-up and excitation coils, andreduction in the excitation frequency were expected.

54

Acknowledgments The authors express their gratitude
to Nikkoshi Co., Ltd., Japan, for supplying FeCoV wires.
References
1) J. R. Wiegand and M. Velinsky: U.S. Patent #3, 820, 090(1974).
2) R. Malmhall, K. Mohri, F. B. Humphrey, T. Manabe, H.
Kawamura, J. Yamasaki, and I. Ogasawara, IEEE Trans.
Magn., 23, 3242 (1987).
3) M. Vázquez, C. Gómez-Polo, D.-X. Chen, and A. Hernando,
IEEE Trans. Magn., 30, 907 (1994).
4) A. Takebuchi, T. Yamada, and Y. Takemura, J. Magn. Soc.
Jpn., 41, 34 (2017).
5) Y. Takemura, N. Fujinaga, A. Takebuchi, and T. Yamada,IEEE Trans. Magn., 53, 4002706, (2017).
6) S. Abe, T. Ikenaga, K. Suwa, and A. Matsushita, IEEE Trans.
Magn., 32, 4971 (1996).
7) S. Abe, A. Matsushita and M. Naoe, IEEE Trans. Magn., 33,
3916 (1997).
8) J. Thone, S. Radiom, D. Turgis, R. Carta, G. Gielen, and R.
Puers, Sensors and Actuators A: Physical, 156, 43 (2009).
Received Oct. 13, 2017; Accepted Dec. 21, 2017

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX
J. Magn. Soc. Jpn., 42, 55-61 (2018)<Paper>Dynamic hysteresis measurement of magnetic nanoparticles
with aligned easy axes
Ryoji Takeda, Satoshi Ota*, Tsutomu Yamada, and Yasushi Takemura
Department of Electrical and Computer Engineering, Yokohama National University,

79-5 Tokiwadai, Hodogaya-ku, Yokohama 240-8501, Japan*Department of Electrical and Electronic Engineering, Shizuoka University,

3-5-1 Johoku, Naka-ku, Hamamatsu 432-8501, Japan

Magnetic particle imaging (MPI) is a novel diagnostic imaging technique based on the use of magnetic
nanoparticles (MNPs). Investigating the magnetic properties of magnetic nanoparticles is important for achieving ahigh spatial and temporal resolution in MPI. In this study, -Fe2O3 nanoparticles (core diameter: dc = 4 nm), Fe3O4
nanoparticles (dc = 20–30 nm), and Resovist® were immobilized in a DC magnetic field with their easy axes aligned.
DC and AC magnetization curves were measured for the prepared MNPs. The measurements were performed byapplying fields parallel and perpendicular to the easy axis and evaluating the magnetic properties of the MNPs forthe easy and hard axes. The direction of magnetic moments under the AC magnetic field applied to the direction of
the easy axis or hard axis was evaluated by using both experimental results and numeric simulation using theLandau–Lifshitz–Gilbert equation to reveal magnetic relaxation property at wide frequency range and the effect ofcore size distribution of oriented MNPs. The effect of anisotropy in superparamagnetic nanoparticles, the relaxation
property depending on the anisotropy energy barrier, and the fast magnetization process of Néel relaxation weresuccessfully observed.
Keywords: magnetic nanoparticle, magnetic particle imaging, magnetic relaxation, magnetic easy axis,
magnetization curve
1 Introduction
Magnetic nanoparticles (MNPs) are promising
candidates for biomedical applications such as cancer
treatment and diagnostic imaging1). In particular,
magnetic nanoparticle imaging (MPI) is attracting
attention owing to its potential for novel diagnostic
applications2). MPI is an imaging technique that detects
MNPs directly through the nonlinearity of theirmagnetization. In vivo cancer imaging using MPI has
been reported with characteristic spatial resolution of
approximately 1 mm3). The MPI signal is strongly
dependent on the size or structure of the used particle4,5),
so the resolution of MPI can be improved by optimizingthese parameters. To achieve high spatial and temporal
resolution of MPI, the magnetic properties of MNPs
should be investigated, and optimal conditions related
to the applied field and the MNP parameters should bedetermined. In this study, we fabricated samples from
MNPs with aligned easy axes and measured theirmagnetic properties.
Orientation of the easy axes of MNPs in an AC fieldobtained by simulations and experiments has beenreported previously6–8). The orientation of the easy axes
should be considered in MPI, which uses an AC
magnetic field for diagnosis. It is important to clarify
the magnetic properties of MNPs in an AC field withregard to the degree of anisotropy. Magnetization
measurement at 20 kHz and its harmonic signal oforiented MNPs have been reported by Yoshida et al.9).
Wemeasuredmagnetization
curvesforsuperparamagnetic and ferromagnetic iron-oxide

nanoparticles with aligned easy axes under an AC
magnetic field with a frequency up to 500 kHz
considering their relaxation properties, and observed
clear anisotropy of the MNPs in this study. As magnetic
properties certainly depend on the core size of MNPs 4),
the effect of core size distribution was also studied withusing the oriented MNPs.
The energy of an MNP placed in an external magnetic
field consists of the anisotropy energy and the energy
associated with the external magnetic field, as given by
the Stoner–Wohlfarth model:

E  K uVM sin 2   μ0 MH cos    .

(1)where Ku, VM, , μ0, M, H, and  are the magnetic
anisotropy constant, the volume of the primary particle,
the angle between the easy axis and the magnetic field,
the permeability of free space, the magnetization of the
MNP, the intensity of the external field, and the angle
between the easy axis and the magnetization,
respectively10). Magnetic moments tend to align withthe direction of the magnetic easy axis and the magnetic
field to minimize the anisotropy energy and the energy
associated with the external magnetic field,
respectively. The anisotropy energy depends on the sizeof the MNP. In a superparamagnetic nanoparticle, theorientation of the magnetic moment is randomized by
the thermal energy, and the magnetic properties show
no coercivity or remanence because the anisotropy
energy is smaller than the thermal energy.
When an AC field is applied to the MNPs, magnetic
relaxation occurs owing to the delay of magnetization in
the AC field. Néel relaxation time N derived from the

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

55

INDEX

rotation of the magnetic moment is given by thefollowing equation:

K V 
τ N  τ 0 exp u M  ,
 k BT (2)where 0, kB, and Tare the attempt time, the Boltzmann
constant, and the temperature, respectively11,12).
2 Materials and methods
2.1 Materials and samples
Water-dispersed -Fe2O3 nanoparticles supplied from
Meito Sangyo Co. Ltd. with core diameters (dc) of 4 nm
were used as superparamagnetic nanoparticles. They
were coated with carboxymethyl-diethylaminoethyl
dextran. In addition, Fe3O4 nanoparticles purchasedfrom Nanostructured and Amorphous Materials Inc.
with dc of 20–30 nm were used as ferromagnetic
nanoparticles. They were coated with polyethylenimine
to achieve better dispersion in the solution.
Two types of solid samples were prepared for either
type of MNPs. One solid sample contained MNPs fixedwith epoxy bond in the absence of magnetic fields, andthe other solid sample contained MNPs fixed in a DC
magnetic field applied with an electromagnet for 8 h.
Thus, the easy axes of MNPs in the first sample wereoriented randomly. In contrast, in the second sample,
the magnetic easy axes of the MNPs were aligned. The
intensity of the DC field during the preparation of thesample with aligned easy axes was 575 kA/m. The
concentration of MNPs in both types of samples was
adjusted to 2 mg-Fe/ml.
We also prepared samples with the aligned easy axes
using Resovist®, a contrast agent generally used in MRI.
Resovist® is a commercially available MNP thatindicates superparamagnetic behavior5) and isgenerally used for experiments on MPI2) andhyperthermia13). It has been reported that Resovist®
consists of multi core particles with their certain size
distribution14). The magnetic property of Resovist® wascompared with those of the -Fe2O3 (4 nm) samples
which exhibited smaller core size distribution5).
2.2 Magnetization measurements
DC magnetization curves were measured with avibrating sample magnetometer (VSM), and AC
magnetization curves were measured with a pickup coilat a frequency of 1–100 kHz under the amplitudes of theapplied field of 4 kA/m and 16 kA/m15). For the solid
sample with aligned easy axes, DC and AC hysteresis
measurements were performed by applying magnetic
fields parallel and perpendicular to the easy axis, which
we defined as the easy axis sample and the hard axis
sample, respectively. All measurements were taken at
298 K. The saturation magnetization of the sample wasestimated by fitting the measured DC magnetization
curve at the field intensity of 800 kA/m to calculate themagnetization curve from the Langevin function.

56

2.3 Numerical simulation
In the magnetization measurement, the results
indicate the total magnetization of MNPs in the sample.
Therefore, the magnetic states and magnetization
distribution of individual MNPs were evaluated bynumerical simulations. Numerical simulations were
performed using the Landau–Lifshitz–Gilbert (LLG)
equation7, 16):dmm  H  m  H  ,dt1  2(3)where m, , , and H are the magnetization,
gyromagnetic ratio, damping parameter ( = 0.1), andeffective field, respectively. The magnetization was
calculated as the value normalized by the saturated
magnetization according to Eq. (3) using the Runge–
Kutta algorithm. The gyromagnetic ratio is estimated
by





 0 M sV M 1   2,
2 N k BT

(4)H  H ex  H ani  H th

(5)



where Ms is the saturated volume magnetization of 351
kA/m. The effective field is given by
In this equation,

H ani K u (m  n)n,
μ0 M s(6)where Hex, Hani, Hth, and n are the excited field, themagnetic anisotropy field, the fluctuating field due tothermal noise, and a unit vector along the easy axis,
respectively. The thermal fluctuating field is satisfied
by the following equations:H th,i (t )  0
H th,i (t ) H th, j (t ' ) 

(7)

k BT
2
δ ij δ(t  t ' ) .(8)2
1   μ0 M sVM

In these equations, i and j are Cartesian indices ofdifferent particles. ij is the Kronecker delta function,and  is the Dirac delta function. The AC field of 16
kA/m at 100 kHz was applied parallel andperpendicular to the easy axis as in the experiments. In
the numerical simulations, 28,672 particles were set
with core diameters of 5±2 nm (mean±SD) and magnetic
anisotropy constants of 20±1 kJ/m3, which followed theGaussian distribution. There was no relation between
the core diameters and the magnetic anisotropy
constants. The easy axes of MNPs were completelyoriented to same direction. Distributions of the numberof particles were evaluated according to theirdependence on the orientation of magnetization.
3 Results and discussion
3.1 DC magnetization curves
Figure 1 and Figure 2 show the DC magnetization
curves of both samples at field intensities of 800 kA/m
and 4 kA/m, respectively. The magnetization was

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INDEX

normalized bysaturation
magnetization
andrepresented in the unit of M/Ms. At the intensity of 575
kA/m, which corresponds to the intensity of the DC field
during the preparation of the sample with aligned easy
axes, magnetization greater than 85% of the saturation
magnetization is observed in the sample with randomly
oriented MNPs. For -Fe2O3 (4 nm) nanoparticles (Fig.
2 (a)), the coercivity of the DC magnetization curves inall samples is marginal. This is because the anisotropy
energy is smaller than the thermal scattering energy at
298 K owing to the small volume of the measured Fe2O3 (4 nm) nanoparticles. Therefore, the magnetic
moments are randomly oriented, and the totalmagnetization at zero field is zero, which ischaracteristic of superparamagnetic nanoparticles. The
magnetizations in the easy axis sample and hard axis
sample were larger and smaller than that in the
random sample, respectively. Clear anisotropy of thesuperparamagnetic nanoparticles at 298 K is observed.
The distribution of the magnetic moments is shown in
Fig. 3. When the magnetic field is applied to the MNPs,
the magnetic moments tend to align with the easy axis
owing to the energy associated with the applied field,
and then the easy axis and the hard axis indicate thedirections of large and small magnetization,
respectively. In the absence of the magnetic field, themagnetic moments tend to align with the easy axis, but

0.5

(a)
Oriented (easy axis)
Oriented (hard axis)
Random0(b)Magnetization [M/Ms]

1
0.5

DC
800 kA/m
-Fe2O3 (4 nm)

-500
0
500
Magnetic field [kA/m]

-1
-1000

0.1

Oriented (easy axis)
Oriented (hard axis)
Random

0

-0.3
-20

(b)
Oriented (easy axis)
Oriented (hard axis)
RandomDC
800 kA/m
Fe3O4 (20-30 nm)

-500
0
500
Magnetic field [kA/m]

1000

Fig. 1 DC magnetization curves of (a) -Fe2O3 (4nm) and (b) Fe3O4 (20–30 nm) at field
intensity of 800 kA/m.

DC
16 kA/m
-Fe2O3 (4 nm)

-0.2

1000

0

-0.5

0.2

-0.1

-0.5
-1
-1000

0.3

Magnetization [M/Ms]

Magnetization [M/Ms]

1

3.2 AC magnetization curves
Figure 4 and Figure 5 show the DC magnetization
curves and AC magnetization curves of Fe3O4 (20–30nm) nanoparticles at field frequencies of 1–100 kHz andintensities of 4 kA/m and 16 kA/m, respectively. The
coercivities of the DC and AC magnetization curves arevery similar at the field intensity of 4 kA/m regardlessof the easy axis direction. It can be concluded that themagnetic moment does not exhibit detectable rotation
at such small field intensities as 4 kA/m owing to thelarge diameters of the particles. In contrast, at large
field intensities such as 16 kA/m, the coercivities of the
DC and AC magnetization curves are very similar for
the field applied along the hard axis. However, thecoercivity of the DC magnetization curve is smaller

Magnetization [M/Ms](a)because they are randomly oriented owing to thethermal energy, the total magnetization is zero at zero
field.
On the other hand, Fe3O4 (20–30 nm) nanoparticles
exhibit large magnetization and coercivity when themagnetic field is applied parallel to the easy axis andsmall magnetization and coercivity when the magnetic
field is applied along the hard axis, as shown in Fig. 2(b). This is because the anisotropy energy barrier islarge for the easy axis sample and small for the hard
axis sample.

0.5

0.25

-10
0
10
Magnetic field [kA/m]

20

Oriented (easy axis)
Oriented (hard axis)
Random

0

-0.25
-0.5
-20

DC
16 kA/m
Fe3O4 (20-30 nm)

-10
0
10
Magnetic field [kA/m]

20

Fig. 2 DC magnetization curves of (a) -Fe2O3 (4nm) and (b) Fe3O4 (20–30 nm) at field
intensity of 16 kA/m.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

57

INDEX(a)(b)

and random sample at a frequency of 10 kHz andintensity of 16 kA/m. As the DC magnetization, the easy
axis sample shows the highest value, whereas the hard
axis sample shows the lowest. The harmonic signal in(c)Magnetic moment

MPI can be increased by applying an AC field along theeasy axis due to larger magnetization of the MNP 9).

easy axis

MNP

Fig. 3 Distribution of magnetic moments whenmagnetic field is applied to (a) easy axis and(b) hard axis. Distribution of magnetic
moments at zero field (c).

than that of the AC magnetization curve for the field
applied along the easy axis. We can conclude thatmagnetization reversal occurs at 16 kA/m. The
coercivity remains constant for the frequency range of
1–100 kHz. This is because the peak frequency of Néel
relaxation for the Fe3O4 (20–30 nm) nanoparticles
calculated from Equation (2) is less than 0.01 Hz (Ku =23 kJ/m3)17) and lies almost outside of the measured
frequency range.
Figure 6 shows a comparison between the AC
magnetization curves for the samples of -Fe2O3 (4 nm)
nanoparticles with easy axis sample, hard axis sample,

0.1

(a)
DC
1 kHz
10 kHz
100 kHz

-0.2
-5

(b)
Magnetization [M/Ms]

0.2
0.1

Oriented (easy axis)

-2.5
0
2.5
Magnetic field [kA/m]

-0.2
-5

5

(b)
DC
1 kHz
10 kHz
100 kHz

-0.5
-20
0.5

16 kA/m
Fe3O4 (20-30 nm)
Oriented (easy axis)

-10
0
10
Magnetic field [kA/m]

Oriented (hard axis)

5

20

DC
1 kHz
10 kHz
100 kHz0-0.25

4 kA/m
Fe3O4 (20-30 nm)

Fig. 4 DC and AC magnetization curves of Fe3O4(20–30 nm) nanoparticles at field intensityof 4 kA/m. Direction of applied field wasalong (a) easy axis and (b) hard axis.

58

0

0.25

-2.5
0
2.5
Magnetic field [kA/m]

DC
1 kHz
10 kHz
100 kHz-0.25

4 kA/m
Fe3O4 (20-30 nm)

0

-0.1

0.5

0.25

0

-0.1

Magnetization [M/Ms]

Magnetization [M/Ms]

0.2

Magnetization [M/Ms](a)Although the coercivities of the DC magnetization
curves are marginal in all samples, the easy axis sample
shows the highest coercivity, whereas the hard axis
sample shows the lowest. This behavior indicates that
there is a delay between the magnetization and theapplied field associated with the angle between theapplied field and the easy axis, which implies a
difference in the anisotropy energy barrier. Similar
results for superparamagnetic nanoparticles andferromagnetic nanoparticles have been obtained
previously from both numerical simulations andexperiments18–20).
Figure 7 shows the distribution of the number ofparticles for the 5±2 nm particles obtained fromnumerical simulations at a frequency of 100 kHz andintensity of 16 kA/m. The horizontal and vertical axes
indicate the magnetization of MNP and the percentage
distribution of particles, respectively. The maximum
value of the distribution, which describes thepercentage of particles with reversed magnetization, is

-0.5
-20

16 kA/m
Fe3O4 (20-30 nm)
Oriented (hard axis)

-10
0
10
Magnetic field [kA/m]

20

Fig. 5 DC and AC magnetization curves of Fe3O4(20–30 nm) nanoparticles at field intensityof 16 kA/m. Direction of applied field wasalong (a) easy axis and (b) hard axis.

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX(a)0.1
0

10 kHz
16 kA/m
-Fe2O3 (4 nm)

-0.1
-0.2
-0.3
-20

0.01
0
-0.01
0

0.4

-10
0
10
Magnetic field [kA/m]

0.8

Fig. 6 AC magnetization curves of -Fe2O3 (4 nm)
nanoparticles at frequency of 10 kHz.
Intensity of field was 16 kA/m.
greater in the easy axis sample than in the randomly
oriented sample. However, there are some particles
whose magnetization is opposite to the applied field in
the easy axis sample owing to the strong binding force
associated with the large anisotropy energy barrier. In
the hard axis sample, the maximal value of the
distribution is observed at approximately ±0.2 M/MS.
This indicates that the magnetic moments are
predominantly aligned along the easy axis rather thanmagnetization reversal.
The DC and AC magnetization curves of the -Fe2O3(4 nm) nanoparticles at frequencies of 1–100 kHz andintensity of 16 kA/m are shown in Fig. 8. In Fig. 8 (a),the coercivity of the easy axis sample clearly increases
with increasing frequency. In contrast, the coercivity of
the hard axis sample is marginal up to 100 kHz, as
shown in Fig. 8 (b). According to the results of numericalsimulations (Fig. 7), magnetic relaxation occurs with alarge delay relative to the applied AC magnetic field in
the easy axis sample. In contrast, magnetic reversal
does not occur in the AC field with the oscillation of themagnetic moment on the easy axis in the hard axis
sample. Thus, hysteresis is not observed in themagnetization curves of the hard axis sample even athigh frequencies. This manifests an experimental
observation of the fast Néel relaxation process.
The magnetization (M/MS) under the applied field of
16 kA/m during the magnetization reversal at 100 kHz
was calculated by summation of distributed
magnetization shown in Fig. 7. The magnetizations ofrandom, easy axis, and hard axis samples were derived
as 0.21, 0.30, and 0.11 M/MS from Figs. 7(a), (b), and (c),respectively. The calculated values of magnetizations ofeasy axis and hard axis samples were compared withthe measured magnetization shown in Fig. 8. As
indicated by the arrows in the figure, the measured
magnetization under the applied field of 16 kA/m at 100
kHz were 0.25 and 0.20 M/MS for the easy axis and hard
axis samples, respectively. The difference between 0.25

7
6
5
4

(ｂ)

16 kA/m-16 kA/m

0.21 M/MS

2
1

15
12

-0.5
0
0.5
Magnetization [M/MS]
Oriented (easy axis)
100 kHz5±2 nm1

16 kA/m-16 kA/m

0.30 M/MS

9
6
3
0
-1

(ｃ)

Random100 kHz5±2 nm

3

0
-1

20

Distribution [%]0.2

Distribution [%]Oriented (easy axis)
Oriented (hard axis)
RandomDistribution [%]Magnetization [M/Ms]

0.3

-0.5
0
0.5
Magnetization [M/MS]

1

4
3
2
1
0
-1

0.11 M/MS
Oriented (hasy axis)
100 kHz5±2 nm

16 kA/m-16 kA/m

-0.5
0
0.5
Magnetization [M/MS]

1

Fig. 7 Calculated results for distribution of numberof particles when magnetic field was applied
to (a) random sample, (b) easy axis sample,and (c) hard axis sample.
and 0.20 M/MS is smaller than that between calculated
values of 0. 30 and 0.11 M/MS, which can be explainedby imperfection of particle orientation in the preparedsamples. In spite of this quantitative analysis, thesimulated magnetization agreed well qualitatively withthe measured AC magnetization curves in terms ofcharacterization of difference in the easy axis and hard
axis samples.
Figure 9 compares the AC magnetization curves of

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

59

INDEX

(a)
DC
1 kHz
10 kHz
100 kHz

0.2
0.1

0.25 M/MS

16 kA/m

0 -Fe O (4 nm)
2

-0.1

3

Oriented (easy axis)

-0.3
-20

Magnetization [M/Ms]

0.3

0

-0.1

-10
0
10
Magnetic field [kA/m]
DC
1 kHz
10 kHz
100 kHz

0.2
0.1

0.2

Resovist
-Fe2O3 (4 nm)

0

-0.2

-0.2

(b)

0.4

Magnetization [M/Ms]

Magnetization [M/Ms]

0.3

(b)

0.20 M/MS

16 kA/m
-Fe2O3 (4 nm)
Oriented (hard axis)

4 kA/m
500 kHz
Oriented (easy axis)

-2.5
0
2.5
Magnetic field [kA/m]

-10
0
10
Magnetic field [kA/m]

20

5

0.2
0.1

Resovist
-Fe2O3 (4 nm)

0

-0.1

-0.2
-0.3
-20

-0.4
-5

20

Magnetization [M/Ms]

(a)

-0.2
-5

4 kA/m
500 kHz
Oriented (hard axis)

-2.5
0
2.5
Magnetic field [kA/m]

5

Fig. 8 DC and AC magnetization curves of -Fe2O3(4 nm) at frequencies of 1–100 kHz andintensity of 16 kA/m with model of magnetic
moment distribution. Direction of appliedfield was along (a) easy axis and (b) hard
axis.

Fig. 9 AC magnetization curves of -Fe2O3 (4 nm)
nanoparticles and Resovist® at field intensityof 4 kA/m and frequency of 500 kHz. Directionof applied field was along (a) easy axis and (b)hard axis.

the -Fe2O3 (4 nm) nanoparticles with those of Resovist®
at a field intensity of 4 kA/m and frequency of 500 kHz.
The core size distribution of Resovist® has been
estimated by Yoshida et al., and the size distribution of
6.1–21.6 nm has been reported14). Thus, the particles oflarge diameters with longer Néel relaxation time thanthat of -Fe2O3 (4 nm) exhibit large coercivity in boththe easy axis and the hard axis samples.

easy axis and hard axis of the samples, respectively. The
increase of magnetization under an applied AC field along
the easy axis is expected to increase the harmonic signal
in MPI. In the case of -Fe2O3 (4 nm) nanoparticles, the

4 Conclusions
In this study, samples of immobilized particles of Fe2O3 (dc = 4 nm), Fe3O4 (dc = 20–30 nm), and Resovist®,
whose easy axes were aligned or randomly oriented,
were prepared. Their DC and AC magnetization
properties up to 500 kHz were analyzed, and the effectof easy axis orientation on superparamagnetic
nanoparticles and ferromagnetic nanoparticles wasevaluated. In both superparamagnetic -Fe2O3 (dc = 4nm) and ferromagnetic Fe3O4 (dc = 20–30 nm)
nanoparticles, the DC and AC magnetization curvesindicated larger and smaller magnetization due toanisotropy energy, when the field was applied to the

60superparamagnetic feature was clearly observed withrandom orientation of the easy axis, regardless of thealignment of immobilized particles. In the AC
magnetization curves of the -Fe2O3 (4 nm)
nanoparticles, there was an increase in coercivity withincreasing field frequency applied to the easy axis
direction. In contrast, when the field was applied to thehard axis, the coercivity was almost marginal up to 100
kHz so that magnetization could follow the appliedmagnetic field even at high frequency. Thus, the fastmagnetization process of Néel relaxation was
experimentally observed. The AC magnetization curveof Resovist®, measured at a frequency of 500 kHz,indicated larger coercivity than that of the -Fe2O3 (4nm) nanoparticles. This is because Resovist® has alonger Néel relaxation time owing to the larger core sizeof the particles compared to -Fe2O3 (4 nm). The
frequency characteristics and effect of core size
distribution of MNPs both in the easy axis and hard axis

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

INDEX

directions were clearly revealed.
In addition, numeric simulation using the LLG
equation was performed to understand magnetization
distribution of individual MNPs, and the results agreed
well qualitatively with the experimental results. The
calculated results successfully demonstrated that alarge proportion of magnetization was reversed by theapplication of an AC field to the easy axis and that there
remained particles whose magnetization direction was
opposite to the applied field owing to the strong binding
force associated with the large anisotropy energy
barrier. When the AC field was applied to the hard axis,
the magnetization, which was predominantly aligned tothe easy axis, was oscillated by the applied field.
Acknowledgments The authors wish to acknowledge
Dr. Takashi Yoshida at Kyushu University for his
advice in numerical simulation. This work was partially
supported by the JSPS KAKENHI Grant Number
15H05764, 17H03275, and 17K14693.
References
1) Q. A. Pankhurst, N. K. T. Thanh, S. K. Jones, and J. Dobson:
J. Phys. D Appl. Phys., 42, 224001–224015 (2009).
2) B. Gleich and J. Weizenecker: Nature, 435, 1214–1217(2005).
3) E. Y. Yu, M. Bishop, B. Zheng, R. M. Ferguson, A. P.
Khandhar, S. J. Kemp, K. M. Krishnan, P. W. Goodwill, andS. M. Conolly: Nano Letters, 17, 1648–1654 (2017).
4) T. Yoshida, T. Sasayama, and K. Enpuku: Int. J. Magn. Part.
Imaging, 3, 1703002 (2017).

5) S. Ota, R. Takeda, T. Yamada, I. Kato, S. Nohara, and Y.
Takemura: Int. J. Magn. Part. Imaging, 3, 1703003 (2017).
6) H. Mamiya and B. Jeyadevan: Sci. Rep., 1, 157 (2011).
7) T. Yoshida, S. Bai, A. Hirokawa, K. Tanabe, and K. Enpuku:
J. Magn. Magn. Mater., 380, 105–110 (2015).
8) S. Ota and Y. Takemura: Appl. Phys. Express, 10, 085001(2017).
9) T. Yoshida, Y. Matsugi, N. Tsujimura, T. Sasayama, K.
Enpuku, T. Viereck, M. Schilling, and F. Ludwig: J. Magn.
Magn. Mater., 427, 162 (2017).
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London. Ser. A, 240, 599–642 (1948).
11) L. Néel: Ann. Geophys., 5, 99 (1949).
12) W. F. Brown, Jr.: Phys. Rev., 130, 1677 (1963).
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Bay, H. M. Fan, and J. Ding: Adv. Mater., 27, 1939–1944(2015).
14) T. Yoshida, N. B. Othman, and K. Enpuku: J. Appl. Phys.,
114, 173908 (2013).
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Jeun, S. Bae, and Y. Takemura: J. Appl. Phys., 107, 09B322(2010).
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Received Oct. 13, 2017; Accepted Jan. 18, 2018

Journal of the Magnetics Society of Japan Vol.42, No.2, 2018

61

INDEX

Editorial Committee Members・Paper Committee Members
K. Kobayashi and T. Ono (Director), 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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