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Materials Processing In Magnetic Fields

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Processing in magnetic fields is a rapidly expanding research area with a wide range of promising applications in materials science, development and design. Industry now offers a variety of superconducting magnets specifically designed for this purpose and equipped with cryocoolers that eliminate the need for cryogenic fluids. Numerous research centers dedicated to materials research and processing in magnetic fields have been created around the world. This book is the result of an international by-invitation-only workshop that has been organized to review the most recent activities in this field. Over 50 scientists participated and 39 papers were selected for inclusion in this book.
سال:
2005
ناشر کتب:
World Scientific Publishing Company
زبان:
english
صفحات:
387
ISBN 10:
9812563725
ISBN 13:
9789812701800
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PDF, 20.46 MB

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Materials Science and Engineering: Forging Stronger Links to Users

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2

Materials Processing Handbook

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Materials Processing I n
Magnetic Fields

This page intentionally left blank

Materials Processing In
Magnetic Fields
Proceedings of the International Workshop on Materials Analysis
and Processing in Magnetic Fields
17 - 19 March 2004

Tallahassee, Florida

editors

Hans J Schneider-Muntau
Florida State University, USA

Hitoshi Wada
National institute for Materials Science, Japan

N E W JERSEY * L O N D O N

v
-

World Scientific

SINGAPORE * B E l J l N G * S H A N G H A I * HONG KONG

-

TAIPEI * CHENNAI

Published by

World Scientific Publishing Co. Pte. Ltd.
5 Toh Tuck Link, Singapore 596224

USA ofice: 27 Warren Street, Suite 401-402, Hackensack, NJ 07601
U K ofice: 57 Shelton Street, Covent Garden, London WC2H 9HE

British Library Cataloguing-in-PublicationData
A catalogue record for this book is available from the British Library.

MATERIALS PROCESSING IN MAGNETIC FIELDS
Proceedings of the International Workshop
copyright 0 2005 by World Scientific Publishing Co. Re. Ltd.

All righis reserved. This book, or parts ihereof; may not be reproduced in any form or by any means,
elecironic or mechanical, includingphotocopying, recording or any information storage and retrieval
sysiem now known or to be invented, wiihoui written permission from the Publisher.

Forphotocopying of material in this volume, please pay acopying fee through the Copyright Clearance
Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, USA. In this case permission to photocopy
is not required from the publisher.

ISBN 981-256-372-5

Printed in Singapore by Mainland Press

Editors’ Preface
Processing in magnetic fields is a rapidly expanding research area with a wide range of
promising applications in materials science, development and design. Industry now offers a
variety of superconducting magnets specifically designed for this purpose and equipped
with cryocoolers that eliminate the need for cryogenic fluids. Numerous research centers
dedicated to materials research and processing in magnetic fields have been created around
th; e world. This book is the result of an international by-invitation-only workshop that has
been organized to review the most recent activities in this field. Over 50 scientists
participated and 39 papers were selected for inclusion in this book.
Magnetic fields are at the origin of many effects in materials, of which the following
are a few examples. The high fields available now allow us to levitate diamagnetic matter
with the advantage of containerless processing of materials, or control gravity on earth
from several -g to +g. The magnetization induced in dia- and paramagnetic matter is strong
enough to change the structure and characteristics of materials, typically as a collective
phenomenon. Magnetic anisotropy can be used for aligning fibers, polymers, and
carbon nanotubes, resulting in matrix systems with superior quality. The magnetic field has
an impact on texturing of materials during a phase transition, in both liquid-to-solid and
solid-to-solid state transitions. Grain boundary migration and mobility changes have been
demonstrated in Bi and Zn crystals, giving us a perspective for texture development in
metals. The damping effect of magnetic fields on conductive liquids is exploited for improved crystal growth quality. Further research will investigate the use of static fields for
convection and texture control, possibly at reduced gravity levels. Field geometry and the
application of rotational fields are additional optimization parameters. Of growing interest
are the effects of magnetic fields in biology and their beneficial applications. Magnetic
fields can help manipulate cells and cellular processes, such as cell divisions. Magnetic
microspheres can be guided within the body for drug delivery and tumor treatment. Another
future application is to mix nanomagnetic particles with biological blood components for
treatment, blood cell separation or as a marker. Other applications of magnetic field
processing are magnetic separation, and processing of crystalline fibers through alignment.
Additional research areas were presented and discussed during the workshop and are
contained in this book.
The International Workshop on Materials Analysis and Processing in Magnetic
Fields has been jointly organized and sponsored by the Tsukuba Magnet Laboratory of the
National Institute for Materials Science in Tsukuba, Japan, and the National High Magnetic
Field Laboratory of the Florida State University in Tallahassee, Florida. It was held at the
NHMFL in Tallahassee on March 17-19,2004.

Hans J. Schneider-Muntau
Hitoshi Wada
V

This page intentionally left blank

Table of Contents
v

EDITORS’ PREFACE

Texturing and Phase Transitions
APPLICATION OF HIGH MAGNETIC FIELDS IN MATERIALS
PROCESSING
S. ASAI

3

HIGH MAGNETIC FIELDS EFFECTS ON SOLID STATE
TRANSFORMATIONSAT HIGH TEMPERATURE
E. BEAUGNON, E GAUCHERAND

11

ENHANCEMENT OF MATERIAL PROPERTIES BY MAGNETIC
FIELD ASSISTED PHASE TRANSFORMATION
B.Z. CUI, K. HAN, H. GARMESTANI,
H.J. SCHNEIDER-MUNTAU, J.H. SU, J.R LIU

19

EXPERIMENTAL INVESTIGATIONOF THE CRYSTALLIZATION
OF BHF IN HIGH MAGNETIC FIELDS
W ERTEL-INGRISCH, K. HARTMANN, X. WANG,
D. HULSENBERG

29

VARIATION OF PHASE TRANSFORMATIONTEMPERATURE
IN FE-C ALLOYS IN A HIGH MAGNETIC FIELD
X.J. HAO, H. OHTSUKA, H. WADA

41

MARTENSITIC TRANSFORMATIONIN SOME FERROUS
ALLOYS UNDER MAGNETIC FIELD
T. KAKESHITA

48

EXPLORING ULTRA-HIGH MAGNETIC FIELD PROCESSING
OF MATERIALS FOR DEVELOPING CUSTOMIZED
MICROSTRUCTURESAND ENHANCED PERFORMANCE
G.M. LUDTKA, R.A. JARAMILLO, R.A. KISNER, J.B. WILGEN,
G. MACKIEWICZ-LUDTKA, D.M. NICHOLSON,
T.R. WATKINS, I? KALU, R.D. ENGLAND

55

vii

viii

FUNDAMENTALSAND APPLICATIONS OF GRAIN BOUNDARY
DYNAMICS IN HIGH MAGNETIC FIELDS
D.A. MOLODOV

66

ENHANCEMENT OF TEXTURE AND CRITICAL CURRENT
DENSITY IN Bi,Sr2CalCu208SUPERCONDUCTING TAPES
THROUGH MAGNETIC FIELD PROCESSING
P.Vl?S.S. SASTRZ U.l? TROCIEWIlZ, H. MAEDA, J. SCHWAR'IZ

80

APPLICATION OF HIGH MAGNETIC FIELD TO TEXTURE
MODIFICATION IN ZINC ALLOY
A.D. SHEIKH-ALI, D.A. MOLODOY H. GARMESTANI

91

TEXTURING FROM LIQUID TO SOLID STATE BY ALIGNING
ANISOTROPIC MAGNETIC NUCLEI IN HIGH FIELDS
R.E TOURNIER

102

Chemical and Physical Processes
REFRACTIVE INDICES OF WATER AND AQUEOUS ELECTROLYTE 115
SOLUTIONS UNDER HIGH MAGNETIC FIELDS
H. HOSODA, H. MORI, N. SOGOSHI, S. NAKABAYASHI
SYNTHESIS OF CARBON MATERIALS BY THE IMPOSITION
OF A HIGH MAGNETIC FIELD
M.-G. SUNG, K. SASSA, A. GEDANKEN, K. IWAI, S. ASAI

124

HIGHLY EXCITED MOLECULES IN MAGNETIC FIELDS
K. TAKAZAWA

136

APPLICATION OF HIGH MAGNETIC FIELD TO CHEMICAL
AND PHYSICAL PROCESSES
I: TANIMOTO, W DUAN

141

MAGNETO-CHEMICAL SYSTEMS UNDER STRONG MAGNETIC
FIELDS: FUNDAMENTALSAND APPLICATIONS
M. YAMAGUCHI, I. YAMAMOTO

147

ix

Control of Liquids
APPLICATIONS OF AC AND DC MAGNETIC FIELDS IN
METALLURGICALAND CRYSTAL GROWTH PROCESSES
A. CRAMER, S. ECKERT, V GALINDO, J. PRIEDE, G . GERBETH

157

ELECTROMAGNETIC PROCESSING OF MATERIALS: FROM
THE CONCEPTS TO INDUSTRIALAPPLICATIONS
I:DELANNOY

169

SEMICONDUCTOR CRYSTAL GROWTH IN STATIC AND
ROTATING MAGNETIC FIELDS
M.I?VOLZ

178

Magnetic Separation
REMOVAL SYSTEM OF ARSENIC FROM GEOTHERMAL
WATER BY MAGNETIC SEPARATION TECHNOLOGY WITH
A SUPERCONDUCTINGMAGNET
H. OKADA, K. MITSUHASHI, 7: OHARA, H. WADA,
I: KUDOH, H. NAKAZAWA, A. CHIBA

197

MAGNETICALLY ENHANCED SOLID-LIQUID SEPARATION
C.M. REX K. KELLER, B. FUCHS

206

NEW APPLICATIONS OF MAGNETIC SEPARATION USING
SUPERCONDUCTINGMAGNETS AND COLLOID CHEMICAL
PROCESSES
S. TAKEDA, S.-J. YU, A. NAKAHIRA, I: IZUMI,
S. NISHIJIMA, 7: WATANABE

220

Biological Applications
NANOMAGNETICS IN BIOTECHNOLOGY
C.-J. CHEN, I: HAIK, J. CHATTERJEE

229

X

STRONG MAGNETIC FTELD INDUCED CHANGES OF GENE
EXPRESSION IN ARABIDOPSIS
A.-L. PAUL, R.J. FERL, B. KLINGENBERG, J.S. BROOKS,
A.N. MORGAN, J. YOWAK, M.W MEISEL

238

NEW APPLICATIONS OF MAGNETIC FIELD TO HUMAN
FRIENDLY MATERIALS AND HUMAN SUPPORTIVE SYSTEMS
S.TAKEDA, U. HAFELI, M.TONOIKE, I:IZUMI, K. EMA,
S.NISHIJMA

243

MAGNETIC ORIENTATION IN BIOLOGY VIRUS STRUCTURE
- BLOOD CLOT ASSEMBLY - CELL GUIDANCE
J. TORBET

249

MANIPULATING CELLS WITH STATIC MAGNETIC FIELDS
J.M. VALLES, JR., K. GUEVORKIAN

257

Diamagnetic Effects
EFFECTS OF MAGNETIC FIELDS ON FEEBLE MAGNETIC
MATERIALS
N. HIROTA, H. UETAKE, T TAKAYAMA, H. NAKAMURA,
M. KURASHIGE, S.HARA, Z SAITO, I:IKEZOE, T ANDO,
H. WADA, K. KITAZAWA

269

APPLICATION OF MAGNETIC LEVITATION TO PROCESSING
OF DIAMAGNETIC MATERIALS
I. MOGI, K. TAKAHASHI, S.AWAJI, K. WATANABE,
M. MOTOKAWA

278

PROTEIN CRYSTAL GROWTH IN LOW GRAVITY PROVIDED
BY A NEW TYPE OF SUPERCONDUCTINGMAGNET
N.I. WAKAYAMA,D.C. YIN, Z TANIMOTO, M. FUJIWARA,
K. HARATA. H. WADA

285

xi

Magnetic Anisotropy and Alignment
ALIGNMENT OF SINGLE WALL CARBON NANOTUBES UNDER
HIGH MAGNETIC FIELDS UTILIZING A SELFiiORGANIZING
OF EPOXY MATRIX
M.S. AL-HAIK, H. GARMESTANI, D.S. LI, M.Y:HUSSAINI,
K. DAHMEN, R. TANNENBAUM

295

ENHANCEMENT OF NANO-MECHANICALPROPERTIES OF AN
EPOXY PROCESSED UNDER HIGH MAGNETIC FIELDS
M.S. AL-HAIK, H. GARMESTANI, D. LI, M.I:HUSSAINI,
R. TANNENBAUM, K. DAHMEN

303

PROCESSING OF POLYMERS USING MAGNETIC FIELDS
E.P DOUGLAS

310

PROCESSING OF POLYMERIC MATERIALS UNDER MAGNETIC
FIELDS
Z KIMURA,M. YAMATO

32 1

MAGNETIC FIELD CONTROL OF STRUCTURESAND PROPERTIES 330
OF DIAMAGNETIC MOLECULAR ASSEMBLIES
I. OTSUKA, Z TAKAHASHI, K YAGUCHI, H. ABE, S. OZEKl
MAGNETIC ALIGNMENT AND CRYSTALLIZATIONBEHAVIOR
OF ISOTACTIC POLYSTYREm
M. YAMATO, Z KIMURA

337

Other Topics
INDUSTRIAL APPLICATIONS OF MAGNETIC RESONANCE
M.J. HENNESSY

347

GENERATION OF UNIFORM MAGNETIC FORCE FIELDS
0. OZAKI, S. MATSUMOTO, Z KIYOSHI, H. WADA

352

xii

EFFECT OF MAGNETIC FIELDS ON EXPLOSIVE WELDING
OF METALS AND EXPLOSIVE COMPACTION OF POWDERS
G.A. SHVETSOV VI. MALI, YU.L. BASHKATOK
A.G. ANISIMOV A.D. MATROSOV T.S. TESLENKO

360

AUTHOR INDEX

37 1

Texturing and Phase Transitions

This page intentionally left blank

APPLICATION OF HIGH MAGNETIC FIELDS IN MATERIALS
PROCESSING
S . ASAI
Dept. of Materials Processing Engineering, Graduate School of Engineering,
Nagoya University, Furo-cho, Chikusa-ku,Nagoya, 464-8603. Japan
The history of Electromagnetic Processing of Materials (EPM)is described. The
application of a high magnetic field in EPM is classified and then the two topics of
quantitative evaluation of phase transformation and texture alignment of ceramics are
introduced, which are our recent endeavors relating EPM to a high magnetic field.

1. Introduction
In the metal industry, electrical energy has long been used as heat energy due to
its cleanliness, high controllability and high energy density. Technologies using
electric energy were developed rather early and progressed without a
background of sufficient scientific understanding. Good examples of this- are
electromagnetic levitation and electromagnetic mixing that were invented in
1923 and 1932, respectively. To bridge the gap between technology and
scientific understanding, magnetohydrodynamics, which had been established
by Alfven in 1942, was first intioduced in 1982 at the IUTAM Conference titled
“The Application of Magnetohydrodynamics to Metallurgy”, held in Cambridge,
England. The conference may have introduced many people to the field of
Electromagnetic Processing of Materials (EPM), though the term “EPM” was
first formally used at the initial Symposium of EPM held in Nagoya, Japan in
1994. EPM research has been hitherto devoted to the economics of mass
production and nanotechnology in relation to high quality materials. Today,
EPM involves both Lorentz and magnetic forces relating to high magnetic fields.
Here, the application of a high magnetic field in EPM is classified and the
two topics of quantitative evaluation of phase transformation and texture
alignment of ceramics are introduced. These are our recent endeavors relating
EPM to a high magnetic field.
2.

Application of a High Magnetic Field in EPM

The technology relating to crystal orientation, structure alignment and spin
chemistry has emerged in EPM thanks to the development of superconducting
technology. Now, a high magnetic field utilizing a large space is available even
3

4
in small-scale laboratories. Table 1 indicates the utilization of a high magnetic
field in EPM.
A high magnetic field allows induction of crystal orientation, i.e. structural
alignment, even in non-magnetic materials. There are four necessary conditions
for crystal orientation under the imposition of a magnetic field. The first is .that
unit crystal cells of the material have a magnetic anisotropy. The second is that
the magnetization energy should be larger than the thermal energy. The third is
that materials should exist in the weak constraint medium in which a particle can
rotate under a feeble magnetization force. The fourth is that each particle should
be dispersed in the medium as a single crystal.
Table 1 Utilization of a high static magnetic field in EPM
Lorentz Force-----------Appearance of small electric current effect

-

MassTransport- - - - - - - - - - -Eliminatbn of inclusionsand surface defect

( x M B - V)B

Magnetizatior

I

(structualalignrnent phase transfomtion

Spin Chemistry----------Intermolecular cross-linking reaction

The possibility of magnetic transportation and magnetic rotation was
examined under several processes, such as solidification [ 1-41,
electro-deposition [ 5 ] , vapor-deposition [6-81, and solid phase reaction [9]. The
application of a high magnetic field has now been proven to be a promising
method in EPM.
Strengthening carbon fibers by imposing a high magnetic field is an
example of spin chemistry. Carbon fibers produced from PAN
(polyacrylonitrile) as a precursor are generally subjected to the three heat
treatment processes of stabilization and carbonization, followed by
graphitization. The carbon fibers produced from stabilized fibers in a magnetic
field showed higher tensile strength than those produced without a magnetic
field [9]. The fibers processed in a magnetic field have a larger crystallite size
than those treated in no magnetic field. An intermolecular cross-linking reaction
model [ 101 describes the crystallite size increase due to the imposition of a high

5

magnetic field. Regarding the spin chemistry [ll-161 upon which the
background of this model stems, one can say that this study is the first attempt to
link EPM with spin chemistry. The development of materials processing based
on the spin chemistry is considered to be a promising area in EPM.

3. Quantitative Evaluation of Phase Transformation
A new apparatus has been developed that can continuously measure the
magnetic force during phase transformation. The magnetic susceptibility is
calculated from the magnetic force obtained by using the apparatus, then the
transitional solid fraction during the solidifying and melting processes is
evaluated from the magnetic susceptibility. The magnetic susceptibility was
measured using the Gouy method [17,18] and is evaluated by measuring a
magnetic force F,.

Figure 1. Calculation of solid fraction.

The solid fraction in a solid-liquid mixed phase can be calculated from the
observed magnetic susceptibility, as follows. From Figure 1, which displays the
relationship between magnetic susceptibility and temperature, the magnetic
susceptibilities of solid and liquid phases can both be expressed with good
approximation by linear functions of the temperature around the melting point.
Specifically, the magnetic susceptibilities in the single solid and liquid phases
are given by Equations ( 2 ) and (3).

6

xms=C,lT+Cs2

(3)

The magnetic susceptibility of a liquid and solid mixture is given by Equation

(4).

xm=

f lxmlf

f sxms

(4)

In addition, Equation (5) holds that

f1ffs=l

(5)

Then

f, = x m
xm

-xld

-Xd

Once the magnetic susceptibility
and temperature of a mixture are
measured, the solid fraction f, can be derived from Equations (2), (3) and (6).
The relation between the solid fraction and temperature in the cooling process is
shown in Figure 2. The solid phase of about 20mass%precipitated until the
moment that the re-coalescence had finished and the temperature had recovered
to the melting point. The liquid phase of about 20mass%remained even after the
temperature descended below the melting point, i.e. about 2Omass% of melt was
supercooled.
1
Q9
Q8

3
e Q7
.IQ6
z

E

Q5
Q4

z* a3

Q2
Q1
0

2 4 0 2 5 0 m 2 7 0 2 8 0 2 9 0
Temperature&)
( Sample:Bi, under Ar atmosphere )

Figure 2. The relation between temperature and solid fraction (cooling).

7

The method developed here can be applied to the direct observation of
various phase transformation phenomena in solid, liquid and gas phases,
therefore we hope it will lead to a better understanding of various phase
transformations and reactions.

Rotatingof a crucible

Figure 3. Schematic view of the experimental apparatus used for rotating a crucible under a magnetic
field.

4. A Novel Process to Fabricate Highly Textured Ceramics in a High
Magnetic Field
A novel process, where a specimen is rotated during a slip casting under a high
magnetic field, has been proposed in order to fabricate highly textured ceramics.
The usefulness of the newly proposed process has been confirmed in Si3N4
ceramics. Figure 3 shows the schematic of the experimental apparatus using
rotation. In order to examine the effect of rotation, green samples were prepared
in a magnetic field and without rotation. For comparison, a sample was also
prepared with no magnetic field. After drying, the green samples were
embedded in a 60wt%Si3N4 +40 wt% BN powder bed in a graphite crucible and
heated to 1800°C for 1.5 hours in an N2 atmosphere, with no magnetic field.
Figure 4 schematically shows the functions of the magnetic field and the rotation.
In the substance in which the magnetic susceptibility in the a, b axis is higher
than that in the c-axis, xc c X,,b, a one-directional crystal orientation can not be
obtained in a slip casting under a high magnetic field, as the free choice of
crystal orientation exists in the a, b axis. When the magnetic field is imposed on
the suspension, the c-axis of the particles can align perpendicular to the
magnetic field. The condition where the specimen is rotated in the magnetic
field is equivalent to the case where the specimen is fixed and the magnetic field

8

Rotating of a
crucible

Figure 4. Schematic view showing the functions of the magnetic field and rotation of a crucible in a
magnetic field.

Figure 5. SEM micrographs of specimens made of a-SijN4 powder with p-SijN4 seeds: (a), (b)
without magnetic field; (c), (d) with magnetic field under crucible rotation.

is rotated. In this case, the c-axis of the particles will be perpendicular to the
plane in which the magnetic field rotates. Thus, the c-axis of the particles aligns
parallel to the direction of gravity. Figure 5 shows the SEM micrograph of the
polished surfaces of the specimen. Also, 0- Si3N4rod grains appear randomly
distributed in the specimen prepared without exposure to a magnetic field
(Figure 5. a and b). When the specimens are prepared by rotation under a

9

magnetic field, a highly textured material can be obtained as shown in Figure
5(c) and (d).

5.

Conclusion

The history of Electromagnetic Processing of Materials (EPM) has been
described. The application of a high magnetic field in EPM has been classified
and listed, and the two topics of quantitative evaluation of phase transformation
and texture alignment of ceramics have been introduced, which are our recent
endeavors relating EPM to a high magnetic field.

Acknowledgement
This work was supported in part by the 21st Century COE Program
“Nature-Guided Materials Processing” of the Ministry of Education, Culture,
Sports, Science and Technology.

References
1. Morikawa, H., Sassa, K., and Asai, S., Muter. Trans., JZM, 39, (1998) pp.
814-818.
2. Yasuda, H., Tokieda, K., and Ohnaka, I., Muter. Trans., JIM, 41, (2000) pp.
1005-1021.
3. Legrand, B.A., Chateigner, D., Pemer de la Bathie, R., and Tournier, R.,
Journal of Magnetism and Magnetic Materials, 173, (1997) pp. 20-28.
4. Noudem, J.G., Beille, J., Bourgault, D., Chateigner, D., and Tournier, R.,
Physica C, 264, (1996) pp. 325-330.
5. Taniguchi, T., Sassa, K. and Asai, S., Muter. Trans., JIM, 41, (2000) pp.
981-984.
6. Mitani, S., Bai, H.L., Wang, Z.J., Fujimori, H., and Motokawa, M., The 3rd
International Symposium on Electromagnetic Processing of Materials, Japan,
ISIJ, (2000) pp. 630-634.
7. Tahashi, M., Sassa, K., Hirabayashi, I., and Asai, S., Muter. Trans., JIM, 41,
(2000) pp. 985-990.
8. Awaji, S., Watanabe, K., Ma, Y., and Motokawa, M., Physica B, 294-295,
(2001) pp. 482-485.
9. Ito, M., Sassa, K, Doyama, M., Yamada, S., and Asai, S., TANSO, 191,
(2000), pp. 37-41.
lO.Sung, M.G., Sassa, K., Ogawa, H., Tanimoto, Y., and Asai, S., Muter. Trans.
43 (2001), pp. 2087-2091.
ll.Tanimoto, Y., Hayashi, H., Nagakura, S., Sakuragi, H., and Tokumaru, K.,
Chem. Phys. Lett. 41 (1976) p. 267.
12.Hatta, N., Chem. (1976) p. 547.

10

13.Shulte1-1, K., Staerk, H., Weller, A., Werner, H.J., and Nickel, B., Z. Physik.
Chem. NF 101( 1976) p. 37 1.
14.Michel-Beyerle, M.E., Haberkorn, R., Bube, W., Steffens, E., Schroder, H.,
Neusser, H.J., Schlag, E.W., and Seidlitz, H., Chem. Phys. 17 (1976) p. 139.
15.Ihara, I., Kato, M., Kanamori, I., Nakamura, K., Shimada, E., and Watanabe,
T., Symposium on New Magneto-science 2002, Proceedings of the 6"
Meeting, Nov. 2002, Tsukuba, Japan, pp. 298-302.
16.Asai, S.,Koumoto, K., Matsushita, Y., Yashima, E., Morinaga, M., Takeda,
K., Iritani, E., Tagawa, T., Tanahashi, M., Miyazawa, K., Science and
Technology of Advanced Materials, 4 (2003) to be published.
17.Iguchi, Y., Experimental chemical course 9, Maruzen Ltd. (1991) pp.
439-450.
18.Suzuki, N., Metal data book, Japan Metal Institute, Maruzen Ltd. (1974) pp.
10, 18.

HIGH MAGNETIC FIELDS EFFECTS ON SOLID STATE
TRANSFORMATIONS AT HIGH TEMPERATURE
E. BEAUGNON, F. GAUCHERAND
CNRSKRETA-LdC, Grenoble, France
Annealing in a high magnetic field of asquenched Co-B near-eutectic alloy promotes
solid-state anisotropic growth of ferromagnetic Co particles along the magnetic field.
Competition between surface energy and demagnetizing energy or effect of the magnetic
torque and creeping of the matrix can both qualitatively explain the observed alignment.
However, both mechanisms are in contradiction with the non-saturation of the
phenomenon in fields up to 16 Tesla. The effect of the magnetic force on the diffusion
near a ferromagneticparticle is discussed.

1. Introduction
In recent years, high magnetic fields have been widely used in material
processing experiments, particularly in magnetic texturing from a solidifying
melt where the residual magnetic anisotropy of solidification nuclei allows their
magnetic alignment in the liquid phase [1,2]. The aim of this work is to study
strong magnetic field effect on texturing, but during solid-state transformations.
In this case, no fluid phase can allow the free rotation of anisotropic particles.
However, it is expected that slow diffusion processes can also lead to
anisotropic textures.
Magnetic field effect on diffusivity has been tested in a few experiments.
Youdelis et al. [3], from experimental results in the Al-Cu system, discussed the
inhibition of diffusion perpendicular to an applied field due to Lorentz forces on
diffusion-transported ions and electrons. In the A1-Cu system, the Cu diffusion
was reduced by 25 % in a magnetic field of 3 Tesla. In contrast, Nakajima et al.
[4] could not observe any significant effect on the diffusion of Ni in Ti, in fields
up to 4 Tesla. In the austenite to ferrite transformations in steels, Xu et al. [5]
measured a magnetic-field-dependent parabolic growth rate, but, depending on
the temperature, the magnetic field effect could be either an inhibition or an
enhancement of the diffusivity.
Strong texturing effects have also already been observed in clean Bi
bicrystals, where the growth of the grain with the axis of larger magnetic
susceptibility is promoted at the expense of the other grain [6]. In steels,
Ohtsuka et a1 observed a shape alignment of large ferrite grains after the
austenite to ferrite transformations in high magnetic field.
11

12
In this study, the magnetic texturing of Co precipitates in a Co-B eutectic is
studied. The high Curie temperature of the Co phase allows annealing
temperatures close to the melting point to promote diffusion, but still in the
ferromagnetic state to enhance the magnetic field effects on the precipitates.
2.

Experiments

Cylindrical samples of Co-B near-eutectic alloys, containing 18.5 B at% [7],
were prepared by induction melting in a cold crucible and cast in a cold copper
mold. As seen by SEM images, the as-quenched alloys exhibit a fine, submicron, lamellar eutectic structure. Parallel lamellas are organized in domains,
but no net orientation of the many domains could be observed.
Samples were then annealed at 900 "C for 65 h in 0 Tesla, 7 Tesla and 16
Tesla. Polished surfaces cut in a plane along the field direction were observed by
optical microscopy. After annealing, the microstructure strongly differed from
that of the as-quenched alloy: Cobalt ferromagnetic particles coalesced to form
large (several microns) globular particles with random shapes dispersed in an
homogeneous matrix which is assumed to be the paramagnetic Co2B phase.
Images were analyzed using a PC version of the free software NlHimage
(Psion Image from Psion Corp.). In this analysis, each Co particle is fitted as an
ellipsoid; the ellipsoid size and orientation versus applied field is then measured
on several thousands of particles for statistical analysis of the magnetic
orientation.

3. Results on Magnetic Alignment
The particle orientation distribution after annealing is presented in Figure la, b
and c. The number of particles is plotted versus the angle between the particle
long axis and the magnetic field (when applied) or the vertical direction (at 900).
In zero magnetic field, although no orientation could be observed on the ascast lamellar eutectic, a large anisotropy exists in the annealed samples where
more particles are aligned near the horizontal plane. This result suggests that
some texture already existed prior to the heat' treatment: the strong radial
temperature gradient induced in the cylindrical sample during the rapid
solidification in the cold mold promoted a radial growth of the eutectic.
In 7 Tesla, the radial structure still remains but a large alignment parallel to
the field is observed. The alignment effect is even more obvious in the sample
submitted to 16 Tesla, where almost all of the radial alignment is erased and
only the magnetic alignment at 90" prevails.

13

0

30

60

90

120

150

180

Argyle
Figure la. Particle counts versus angle for a zero magnetic field annealing.

0

30

60

90

120

Angle

150

la0

0

30

60

90

120

150

1BO

Angle

Figure Ib (left) and Ic (right). Magnetic field effect on the particle angle distribution after annealing.

An order parameter has been defined as the number of particles whose long
axis is closer to the field direction divided by the number of particles whose long
axis is closer to the perpendicular direction. A random distribution would give 1
and, for an axial distribution along the field direction, a value larger than 1. In
zero field, the value is 0.7, revealing the radial texture due to fast solidification
of the cast alloy. In 7 Tesla, it reaches 1.2 and rises to 1.95 with 16 Tesla: no
saturation of the phenomenon is observed above a few Tesla. Several models
can qualitatively account for the magnetic texturing in the solid state and are
discussed in view of this non-saturation.
4.

Competition Between Surface Energy and Demagnetizing Energy

The minimization of the interfacial energy of the Co particle in the surrounding
matrix, obtained for a spherical shape, competes with the demagnetizing energy,
which is a minimum for an elongated shape along the field. For an ellipsoid
along the field direction, the equilibrium shape is defined by the minimization of

14
the total energy E as a function of cL=c/a, the shape factor of the ellipsoid, for a
constant volume V.
The total energy E is given by: E = oS+'/z~nM2Vwith S the particle
surface, M its magnetization, (T the surface energy and n the demagnetizing
factor. For a given average radius R, and an elongated ellipsoid, n and S are
given by:
1
a*-1

J

argch(a) - 1
JZ

The demagnetizing factor n varies from 1 for a flat shape to 0 for a needle.
Below the saturation, the magnetization M in an applied field Ba is governed by
the demagnetizing field with hM=Ba/n. h M is about 1 Tesla at high
temperature, so that the saturation is at least reached from Ba = 1 T, whatever
the n value. The surface energy is unknown, but typical values are in the range
of 0.1 - 1 J/m2 for metals. The observed average radius of the particle is in the
range of 4 - 10 pm. Several examples of the total energy are given Figure 2.
Jhm3

4 p m 0.1 J/m2

Jhi?

4 p m 0.5 Jjm'

6MoW

\

625000
M)oo00

575000

y.Il_.
525000

2

4

6

8

i

O

8 p r n 0.5 J/mz

Jhi?
31GQW
330000
320000
310000

moo
290000

I

2-

6

8

10

Figure 2: sum of the demagnetizing energy and total surface energy for different average radius R
and surface energy o.

15

In each case, it is found that the Co particle equilibrium shape is noticeably
elongated along the field. However, the saturation magnetization is already
reached at 1 Tesla, so that no more effect should have been observed in 16 Tesla
as compared to 7 Tesla. In addition, this simple model does not account for
particles that are already elongated in random directions as observed in the
samples.
5. Magnetic Torque
As most of the particles have an anisotropic shape, a magnetic torque is
developed in the magnetic field. It can then be expected that, for a long
experiment time, the slow creeping of the surrounding matrix will allow the
rotation of the anisotropic Co particles.
In a low external field, Ba, below a threshold field, Ba*, where the
magnetization saturates, M is governed by the demagnetizing field, so that the
inner field (Ba - k n M ) is zero, with n the demagnetizing factor along the long
axis of the particle. Let I$ be the angle between the applied field and the major
axis of the particle, then the magnetization, M, is set at a constant angle 0 with
Ha so that 0 < 4, with:

Ba < Ba* = p,Msl,/a
d
p,,M = B

and tan(6) = tan(cp)E is constant

When M saturates to Ms, the demagnetizing field can no longer cancel the
external field. The magnetization is no longer locked at a constant angle but
rotates to align with Ha, so that 0 is now given by:

1-3n
Ms
sin(cp - 0) = -sin(26)4
Ha
As Ba is increased, and because 0 goes to zero, the magnetic torque

r = (MAB)V

has a finite limiting value r(Ba=-):

r(Ba = -) = poM s 2 Vysin(2cp)
For any n values and for different angles I$, it is then found that the limiting
torque value is rapidly approached below a few Tesla, thus the experimental
high field dependence of the magnetic orientation is not taken into account in
this torque model.

16

6. Curvature Effects and Local Magnetic Forces
The preceding mechanisms involve a particle transformation (shape factor or
angle versus field) at a constant volume. In the actual annealing experiment, the
particles are growing from a sub-micron lamellar eutectic to a globular
distribution with a typical particle size of a few microns. The mechanism by
which particles coalesce is the Ostwald ripening where large particles are
growing at the expense of the smaller, which dissolve. The driving force of the
diffusion from small to large particles is the shift of the concentration
equilibrium in the matrix at the interface with the particle, due to curvature
effects. For a mean curvature (Ur), this concentration shift is given by the
Gibbs-Thomson equation:

with C the concentration and AC the concentration shift of the Co ions in the
surrounding matrix at the interface with the ferromagnetic Co particles. y is the
interfacial energy and 52 the molar volume.
Near a ferromagnetic particle, the local magnetic field is distorted with a
maximum value at the vertical poles (along the external field direction) and a
minimum value at the equatorial plane. Strong local magnetic gradients exert
forces on the diffusing Co ions in the surrounding matrix (see Figure 3), but the
thermal energy kT almost counterbalances the effect. The equilibrium
concentration profile then obeys a Boltzman distribution and is proportional to
x , ~is the magnetic susceptibility per single Co
e x p ( ~ , ~ B ~ / 2 k kwhere
T)
paramagnetic ion.

.-...
-.-...< \ * , * , , , , , , - - .
< . . 9 , , , , . , , , . - -

C

--..,,.l.I.1\......~ _ _ , , , , l l l l L . \ , , . . - -

.-*.,,,,...,..,..--Figure 3: schematic magram of the magnetic forces near a sphencal particle (left) and local effect on
the Co ion concentration in the surroundmgmatnx (right).

A first order development leads to the concentration variation AC that then
follows:

17

-_AC C

(x.1

Ba)AB
kT

1 0

- ( x d Ba)AB
1 0

RT

xmol

with Ba the applied field,
the molar susceptibility and AB the local field
variation.
Near the magnetic poles of the particles, a small radius rtipcan be stabilized
by the magnetic force that increases the local concentration in the matrix and
prevents dissolution. On the other hand, near the equatorial plane, the magnetic
force expels ions except in the case of a very elongated particle where the local
distortion field is nearly zero.
A stable configuration, considering both Gibbs-Thomson and magnetic
effect, is then a particle elongated along the field direction, with a radius rtip at
the poles for which both effects are equal. With the S.I. susceptibility of the
ions in the matrix and Ms the magnetization of the particle, the radius r,ipis then
given by:

x

x

For y = 0.1 J/m2, Ba = 16 Tesla, p&ls = 1.1 Tesla, and = 7.5
(a value
extrapolated from the susceptibility measured in the liquid state [8]), one finds
rtip= 1.5 pm, a quite realistic value for the particles actually observed.
7.

Conclusion

The application of a high static magnetic field (up to 16 Tesla) on the high
temperature solid-state annealing of a Co-B quenched eutectic alloy leads to an
anisotropic microstructure where ferromagnetic Co particles are statistically
aligned along the field direction.
Two models were proposed to explain this result:
- the particles elongate along the field to minimize their demagnetizing field;
- the particles rotate to align their longer axis along the field.
In both cases, transformations occur by solid-state diffusion, and the magnetic
field effect should saturate with the maximum magnetization of the
ferromagnetic Co particles. It is experimentally demonstrated that, up to 16
Tesla, no saturation is measured, in contradiction with both models.
A new model is proposed, based on the strong local magnetic forces around
a ferromagnetic particle that polarize the diffusion and modify the Ostwald
ripening classical scheme. In this model, growth near the magnetic poles is
promoted and local magnetic attraction prevents the dissolution of the tip of the
particle so that a final, stable elongated shape along the field can be obtained.

18
Further investigations are required to develop this new model, but the first
estimations of the magnetic force effects are consistent with experimental
results.

References
1.
2.
3.
4.
5.
6.
7.
8.

P. de Rango et al, Nature, (1991), p. 349.
E. Beaugnon et al, Journal de Physique I, 3 , 2 (1993).
Youdelis et al., Canadian Journal ofphysics, 48, (1970) and 42 (1964).
Nakajima et al, Trans. JIM, 26, 1 (1985).
Xu et al., Trans. MRS Japan, 25,2 (2000).
Molodov et al., Sriptu Muterialla, 37, 8 (1997).
S. Omori et al., Trans. JIM, 17 (1976).
F. Gaucherand, PhD thesis, Joseph Fourier University (2001).

ENHANCEMENT OF MATERIAL PROPERTIES BY
MAGNETIC FIELD ASSISTED PHASE TRANSFORMATION
B.Z. CUI

'**, K. HAN ',H. GARMESTANI

H.J. SCHNEIDER-MUNTAU
J.H. SU J.P.LIU *

'National High Magnetic Field Laboratory, Florida State University,
Tallahassee, FL 32310
'Department of Physics, Universig of Texas at Arlington, Arlington, TX 76019
3School of Materials Science and Engineering, Georgia Institute of Technology,
Atlanta, GA 30332

Using magnetic field assisted phase transformation, an enhancement of the exchange
coupling and hard magnetic properties of melt-spun Nd~FellBla-Fe-typenanocomposites
was achieved by optimizing their nanosmcture and morphology. Compared with the
Nddr5.6DylFeuMolBs sample annealed without a magnetic field, the magnetic
annealing results in a noticeable improvement in the coercivity i&, the remanence 4nM,
and energy product (BH),.
for Nd2.&'rs.~DylFe~MolB6
alloys. (BH)mu at 50 K was
enhanced by 43.7% after magnetic annealing in a 19 T field. The kink in the
demagnetization curve disappeared and, additionally, a much better squareness of the
demagnetizationcurves was observed in the magnetically annealed samples.

1. Introduction

Nanocomposite magnets have drawn extensive attention due to their great
commercial potential and broad applications in nano-electromechanical systems,
automatic control engineering, micro-mechanical devices, magnetic imaging,
magnetic fluids, biomagnetic sensors, nanomedicine, catalysts and other
applications. Nanocomposite magnets have extremely high theoretical energy
products (BH),, of up to 100 MGOe. However, there is still a big discrepancy
between the theoretical and experimental (BH),,
[ 1-51. Some challenges still
remain in the improvement of the naonostructure morphology and crystal
texture, such as refinement of grain size, homogeneity of the grain size
distribution, optimization of grain configuration, and hard nanograin alignment.
A breakthrough in (BH),, could be achieved by preparing homogenized and
textured nanostructures to take advantage of the magnetic anisotropy of hard
phases and exchange coupling between the soft and hard nanograins [4,5].The
present work reports an approach to achieve nanostructural optimization and
magnetic field-induced crystal texture through magnetic field assisted or
controlled phase transformation of Nd2Fe14B/cl-Fe-type melt-spun
Nd2.J'r5.,JIy1Fes4MolB6 nanocomposites.
19

20

2. Experimental Procedure
Nd2.~r5,6Dy1Feg4MolB6
ribbons were prepared by melt spinning with a
molybdenum wheel speed of 35 d s . The ribbons were annealed at 690" C for
20 min with and without an in-plane field of 19 T.
The phase components, nanostructured morphology and crystallization
behavior of the samples were studied by x-ray diffraction (XRD) using Cu Ka
radiation, a JEOL-2010 transmission electron microscopy (TEM) and a PerkinElmer DSC7 differential scanning calorimeter (DSC) at a heating rate of 20"
C h i n . The average grain sizes of a-Fe were deduced from Scherrer's method of
XRD. The magnetic properties were measured by a Quantum Design SQUID
magnetometer in fields up to 6.5 T. The ribbon plane was placed parallel to the
magnetic field direction. No demagnetization correction was done for these
ribbon samples.

3. Results and Discussions
As indicated by XRD investigations, the as-spun Nd2.4Pr5.~y1Feg4Mo
sample is a composite in which a small amount of a-Fe and
(Nd,Pr,Dy)2(Fe,Mo)l4B (2: 14:1) nanocrystallines are embedded in an
amorphous matrix. DSC study of the as-spun Nd2.4Pr5.6Dy1Fe84MolB6
alloy
without a magnetic field shows that nucleation temperatures T,, of the a-Fe and
2:14:1 phases are 525" C and 590" C, respectively. Annealing at 690" C for 20
min leads to the formation of a mixture of nanostructured a-Fe and 2:14:1
phases. Figure 1 shows XRD patterns of the Nd2.4Pr5.6Dy1Feg4M01B6 ribbons
annealed at 690°C for 20 min without and with a 19 T in-plane field. Compared
with the sample annealed with a 19 T field, it can be seen that there are fewer
XRD peaks of the 2: 14: 1 phase in the sample annealed without a field i.e., some
peaks with relatively low intensity even disappear. The average grain sizes of aFe, calculated from the a-Fe (110) diffraction peak using the Scherrer formula
are 17 nm and 20 nm, respectively, for the samples annealed with and without a
19 T in-plane field. The error bars for the average grain sizes are 10%. Figure 2
shows the TEM bright field images and corresponding selected area diffraction
of the two samples shown in Figure 1. It can be seen that magnetic annealing
introduces somewhat finer, less angular (fewer sharp edges) and more
homogeneously distributed soft and hard nanograins.

21

I

25

- I1

1199 TT

0

30

35

40

45

0-2:14:1
t - a-Fe

50

2 8 (degree)
Figure 1. XRD patterns of N ~ ~ ~ P ~ s ~ D ~samples
~ F ~ &
annealed
o I B at~ 690" C for 20 min without
and with a 19 T in-plane field.

Figure 2(a). E M bright field image and corresponding selected area diffraction of
Nd~Srs.6DylFe&lolB6sample annealed at 690" C for 20 min without a magnetic field.

Figure 2(b). TEM bright field image and Corresponding selected area diffraction of
Ndz S r 5 6DylFes&O& sample annealed at 690" C for 20 min with a 19 T in-plane field.

22

Crystallization of amorphous solids is generally considered as a nucleation
and growth process of the crystalline phases. In the nucleation theory the steady
state homogeneous nucleation rate I,, is given by [6]

I,, = b*exp(-Q/RT)*exp(-AGJRT)

(1)

where b is a pre-exponential factor, Q is the activation energy for the transfer of
atoms across the surface of the nucleus, which is approximately equal to the
diffusion activation energy, and AG, is the free energy required to form a
nucleus of the critical size. AG, can be written as

AG=
~ A?/(AG:).

(2)

where AG, is the Gibbs free energy difference between the crystal and the
matrix amorphous phase, ie. AG, = Gcrysa- G m o r p ~cu s0, A is a coefficient and
y is the interfacial energy of the crystaVamorphous interface. In the present
work, Gcvsa can be the Gibbs free energy G,-F~or G2.4:. for the a-Fe and 2:14:1
phases, respectively. The amorphous matrix is magnetically soft with Curie
temperature T, of 310" C and coercivity of 0.6 kOe at room temperature, which
is paramagnetic during the crystallization process. During the crystallization
process and subsequent isothermal anneal at 690"C, the newly formed
nanocrystallines are a ferromagnetic a-Fe phase (T, = 771" C) and a
paramagnetic 2:14:1 phase (T, = 325" C). When a magnetic field is applied to
the system, the Gibbs free energy Ga-~e
of the a-Fe nanocrystallines decreases
due to the addition of the magnetostatic energy. As the new phase a-Fe is still a
ferromagnetic phase and the matrix is paramagnetic, the extra Gibbs free energy
difference AGHintroduced by a magnetic field H is [7]

AG" = -@M(T).H

- 1/2*~*H'+ E , * ( & ~ ~ H ) * H * B

(3)

The first term, -kAM(T)*H, represents the energy due to the magnetostatic
effect, where AM(T) is the difference in magnetization between the a-qe and
amorphous matrix at certain temperature. The terms - 1 / 2 * ~ * H and
&,*(dw/dH)*H*B represent the energies due to the high field susceptibility and
forced volume magnetostriction effects, respectively, where x is the high field
susceptibility in the matrix, E, the volume change associated with the phase
transformation, o the force volume magnetostriction, and B the matrix bulk
modulus.
It follows that the presence of H increases the absolute value of AG,,
resulting in an increase of the driving force for crystallization and further the
nucleation rate I,, of a-Fe. The above analysis also indicates that the higher the

23

magnetic field, the greater is the I,, for a-Fe. Therefore, magnetic annealing
promotes crystallization of the amorphous matrix and allows more nucleation
centers to form and grow the soft a-Fe phase during the crystallization process.
It was reported that the nucleation rate of ferromagnetic ferrite was remarkably
accelerated by three times under the influence of a 10 T field in the Fe-C alloy.
However, the growth rate is nearly the same with and without a 10 T field, so
the magnetic field has little effect on the atomic jump frequency and the
activation energy for the crystal growth of ferrite [S]. Compared to annealing
without the magnetic field, magnetic annealing leads to reduced a-Fe grain sizes
and more uniform distribution of the grains for both the hard and soft phases, as
is shown by the above-mentioned experimental results. Similar results were also
observed in the melt-spun Nd2Fe14B/Fe3B-typenanocomposites [9].
Alternately, the Gibbs free energy G2:.k1 of the 2:14:1 phase should also
decrease slightly, so the nucleation rate I,, of 2:14:1 phase should increase
marginally in the presence of a high magnetic field. But these changes are much
smaller than those of a-Fe as the 2:14:1 phase is in a paramagnetic state during
the crystallization process. However, from the shape and numbers of the XRD
peaks of the 2:14:1 phase (Fig. I), it is clear that magnetic annealing also
promotes crystallization of the amorphous matrix to form the 2: 14:1 phase as the
magnetic field is as high as 19 T.
Figure 3 gives the dependence of the intrinsic coercivity iH,, the maximum
magnetic energy product (BH),, and the remanence 4nMr on temperatures for
Nd2.$r5.&~1Fe84Mo$6 annealed at 690" C for 20 min without and with a field
of 19 T. Figure 4 shows the representative room temperature demagnetization
curves of the two samples of Figure 3. It is observed from Figures 3 and 4 that,
compared with the sample annealed without a magnetic field, there is a
noticeable improvement in i&, (BH),
and 4nM, for magnetically annealed
Nd2.$r5.&y1Fe84MolB6 alloys. Especially, (BH)max at 50 K and 300 K is
enhanced by 43.7% and 35.7%, respectively, after magnetic annealing in a field
of 19 T. The kink in the demagnetization curve disappears and, in addition, a
much better squareness of the demagnetization curves is observed with magnetic
annealing.

24

--e no

field

-19T
9.0
50

LOO

150

200

250

300

T 6)

Figure 3. Dependence of iK, (BH),,, and 4nM, on temperatures for Nd2.4Pr5.6DylFe~Mo~B6
annealed at 690" C for 20 min without and with a 19 T field.

-0-

690°C, 20 min, 19 T

+69OoC,20 min, no fielc

10
n
v1

8
E 5

e

d

0

-6

-4

-2

0

2

4

H (kOe)
Figure 4. Room temperature demagnetization curves of the two samples of Figure 3.

25
The exchange coupling interactions between soft and hard phases were
evaluated by using 6M (Henkel) plots, in which 6M = w(H)- [l - 2m(H)1 [lo];
where a(€€)
is M,(H)/M,, the reduced isothermal remanence, and m(H) is
&(H)/M,, the reduced'demagnetization remanence. Figure 5 shows 6M plots as
a function of the applied field for the two samples of Figure 3. An initially
positive 6M is observed in both annealed samples with and without a 19 T field,
confirming the existence of ferromagnetic exchange coupling interactions
between the 2:14:1 and a-Fe nanograins. The magnitude of the 6M peak of the
magnetically annealed sample is higher than that of the sample annealed without
a magnetic field, demonstrating stronger exchange coupling interactions
between the soft and hard nanograins in the magnetically annealed sample. We
further estimate the exchange coupling in the two samples of Figure 3 by spinreorientation temperature changes of the 2:14:1 phase.
Figure 6 shows the thermomagnetic curves for the same two samples as
shown in Figure 3. As can be seen in Figure 6, the spin-reorientation
temperatures T,, of the hard phase 2:14:1 in the samples annealed with and
without a field of 19 T are 55 K and 61 K, respectively. The inter-gain
exchange coupling is known to lead to a rotation of the magnetization in
neighboring grains away from the local easy direction and to hinder spin
reorientation. The stronger the inter-grains exchange coupling interactions
between the nanograins, the lower is T,, [11,12]. In this experiment, a lower Tsr
in the sample annealed with a 19 T field convincingly demonstrates that
magnetic annealing leads to a stronger inter-grain exchange coupling. The
improvement of the hard magnetic properties of the magnetically annealed
sample results mainly from the magnetic-field-induced enhanced exchange
coupling, which is due to optimization of the nanostructured morphology by
magnetic annealing, such as reduced grain sizes of a-Fe and the more uniform
distribution of the grains for both the hard and soft phases. Additionally,
magnetic annealing improves the crystal quality of the hard 2:14:1 phase (Figure
l), which is also favorable for the magnetic hardening for nanocomposite
magnets.

26
I

1

0.2

E

I-\

0.0

-

-.-0--8=.==
/o-

0

5

15

10

20

25

H (kOe)
Figure 5. 6M plots as a function of the applied field for the two samples of Figure 3.

0

50

100

150

200

250

300

T (K)
Figure 6. Thennomagnetic curves in an applied field of 0.1 T for the two samples of Figure 3.

It is known that a significant improvement in (BH),,
is expected from
successfully fabricated anisotropic nanocomposites with crystallographic
textures of hard phases [4,5]. Magnetic annealing in a high field is a promising
method for manufacturing textured naomagnetic materials. If the matrix phase is
a liquid, the new phase can move freely and may align along the magnetic field
due to its magnetocrystalline anisotropy. In this case, to align a paramagnetic or
ferromagnetic particle, the driving force (magnetic torque) due to
magnetocrystalline anisotropy of the new phases in a paramagnetic or

27
ferromagnetic state could still be larger at high magnetic fields than those due to
the shape anisotropy and the thermal disordering effects [ 131. In these cases, a
hard phase-textured nanostructure has lower energy than a random structure in a
magnetic field. The magnetically induced crystallographic alignment of the new
magnetic phases (such as 2:14:1 and fct FePt etc.) driven by a
magnetocrystalline anisotropy can, therefore, still be successfully achieved
during the magnetic annealing if the applied magnetic field strength is high
enough [ 131. If the matrix phase is a solid, the magnetic field must overcome the
interface energy term in order to align the new phase as well as the energy
associated with the shape anisotropy and the thermal disordering effects. The
alignment of the new phase can be achieved by changing the habit plane of the
matrix and the new phase in the presence of a high magnetic field. Such a
change of orientation is difficult but possible, especially if the new phase is still
in the ferromagnetic state during the magnetic field assisted solid-state phase
transformation. Further research is underway toward achieving a breakthrough
in (BH),,
by preparation of homogeneously nanostructured and, especially,
highly textured nanocomposites.

4. Conclusions
For Nd2Fe14Bla-Fe-type melt-spun Nd2.&5.&~~Fe8~Mo$6 nanocomposites,
magnetic field assisted phase transformation from amorphous matrix to a
composite of 2: 14: 1 and a-Fe nanocrystallines results in an optimization of the
nanostructure through promotion of the crystallization of the amorphous matrix
and the generation of more nucleation sites. The optimization of the
nanostructure morphology leads to an enhanced exchange coupling and a
noticeable improvement in the hard magnetic properties for the
Nd2,J’r~.6Dy1Fes4Mo~B6
samples.
Acknowledgments
This work was supported by DARPA through AEtO under grant DAAD19-03-10038. Dr. B. Z. Cui thanks Prof. N. Dalal at the Department of Chemistry at
Florida State University for the partial financial support.
References
1. Manaf, R.A. Buckley, H.A. Davies, J., Magn. Magn. Mater., 128, (1993) p.
302.
2. Cui, B.Z., Sun, X.K., Xiong, L.Y., Liu, W., Zhang, Z.D., Yang, Z.Q., Wang,
A.M. and Deng, J.N., J. Muter. Res. 16, (2001) p. 709.

28

3. Jurczyk, M., Collocott, S.J., Dunlop, J.B. and Gwan, P.B., J . Phys. D: Appl.
Phys. 29, (1996) p. 2284.
4. Skornski, R. and Coey, J.M.D., Phys. Rev. B, 48, (1993) p.15812.
5. Fischer, R., Schrefl, T., Kronrniiller, H. and Fiddler, J., J. Mugn. Mugn.
Muter. 150, (1995) p. 329.
6. Lu, K., Muter. Sci. & Eng. R, 16, (1996) p. 161.
7. T. Kakeshita, T. Saburi, and K. Shimizu, Muter. Sci. & Eng. A , 21 (1996) pp.
273-275.
8. Ohtsuka, H., Hao, X.J., and Wada, H., International Workshop on Materials
Analysis and Processing in Magnetic Fields Tallahassee, FL, March 2004, to
be published.
9. Zhao, T.M., Hao, Y.Y., Xu, X.R., Yang, Y.S. and Hu, Z.Q., J . A p p l . P h y s .
85, (1999) p.518.
10.Henke1, O., Phys. Stat. Sol., 7 (1964) p. 919.
ll.Kou, X.C., Dahlgren, M., Grossinger, R. and Wiesinger, G., Jour. Appl.
Phys. 81, (1997) p. 4428.
12.Cui, B.Z., Sun, X.K., Xiong, L.Y., Tang, S.T., Zhang, X.X., Liu, W., Geng,
D.Y. and Zhang, Z.D., J. Alloys and Compounds 340, (2002) p. 242.
13.Courtois, P., de la Bgthie, R.P., and Tournier, R., J. Mugn. Mugn. Muter.
153, (1996) p. 224.

EXPERIMENTAL INVESTIGATION OF THE
CRYSTALLIZATION OF BHF IN HIGH MAGNETIC FIELDS

w. ERTEL-INGRISCH’, K. HARTMAN”, x.WANG’,
D. WLSENBERG’
‘Junior Research Group “Electromagnetic Processing of Materials”
2Departmentfor Glass and Ceramic Technology
Technische Universitat of Ilmenau, Ilmenau, Germany
Ba-Hexafemte (BHF) powder is a hard magnetic material of wide t e c h c a l application.
High-end applications require a very homogeneous nanometer-scale single domain BHF
powder with optimum magnetic properties that are mainly controlled by the conditions
(temperature, time) during its formation. One method to synthesize BHF nanocrystalline
powder with satisfying magnetic properties is the glass crystallizationtechnique [ l ] - [3].
This method starts from melts prepared in the BaO-Fe~03-B~O3-system.which are
homogenized, and quenched to glass flakes applying a double-roller rapid quenching
technique. Obtained glass flakes are processed in a subsequent tempering process
crystallizing very homogeneous nanometer-size BHF powder. A wide variety of
analytical techniques is applied to monitor chemical composition (electron microprobe
analysis, XRD), structural and physical properties of intermediate products (flakes) and
obtained powders @EM, TEM, DTA, XRD, vibrating sample magnetometry). Precise
control of the parameters and optimization of the process leads to BHF powders with a
maximum coercivity of up to 400 W m .
Application of high magnetic fields during the crystallization process of BHF is
expected to result in a net increase of its magnetic properties. Therefore, experiments in a
cryogen-free magnet (CFM) of up to 5 tesla magnetic flux density equipped with a hightemperature oven to precisely control process parameters will be performed to investigate
the influence of a strong magnetic field on its crystallizationprocess. Of special interest
is the complex question whether magnetic fields can help control and improve the single
domain crystal structure by influencing nucleation, domain growth and orientation of
growing domains towards each other. Initial experimental results indicate a fundamental
change in the crystallizationprocess resulting in further optimized magnetic properties of
BHF powders when crystallizationis performed in a strong magnetic field.

1. Introduction
Barium-Hexaferrite (BHF) powder is one of the most important hard magnetic
materials used in industry. Its technical applications range from basic use as
permanent magnet in daily life to high-end technological applications as analog
and digital high-density magnetic recording media, in electronic computation
devices, and generally in electric drive technology.
Especially for high-end applications, very homogeneous sub-micrometerscale single domain BHF powder is required that can be produced by applying a
glass crystallization technique (GCT [ I]-[3]). However, starting from the
29

30

ternary system BaO-Fe~03-Bz03, or in technical respect, BaC03 and FezO3,
compositions with mole fractions of Fez03 higher than 20% tend to crystallize
spontaneously into numerous phases of vastly varying magnetic properties.
Therefore, precise knowledge of the influence of the various parameters
controlling the crystallization of BHF powders and their precise control during
production is essential to guarantee both right chemical composition and
structure as well as optimum magnetic properties of the final product.
Magnetic properties are, of course, controlled on one hand by the
corresponding chemical starting composition to obtain the intended chemical
phases and their magnetic properties. However, the right thermal treatment has a
decisive influence on the initiation of any crystallization due to the formation of
crystallization nuclei. Nucleation is crucial for all subsequent processes, and
decisive for the size and homogeneity of the final product. During an initial step,
as many seeds should be created as possible while growth and crystallization
must be avoided - a prerequisite which can be obtained only hypothetically.
During the subsequent tempering process, BHF crystals start to grow on top
of originally formed crystallization seeds. Since BHF belongs to the hexagonal
system, with its c-axis easily magnetized, honeycomb-shaped crystals grow. The
goal is to obtain crystals of small, uniform size corresponding to a single
magnetic domain to avoid losses in the net magnetic properties of the powder.
Optimum magnetic properties were reported [2] for BHF powders with crystal
sizes of less than 1 pm in diameter, a homogeneous size distribution, and a
narrow aspect ratio (diameter versus thickness) of approximately 3: 1. Compared
to theoretical considerations with the aim to increase both remanence and
coercitivity, the maximum energy density ((B€€)& obtainable from BHF is,
however, about twice as high as technically obtained [3]. Scientific
investigations for further improvement of BHF powders are urgently needed.
Furthermore, there are several scientifically interesting questions still
unanswered: Can crystallization processes like that of BHF be influenced, or
even controlled, applying strong magnetic fields? Can we support and control
the direction of crystallization by magnetically induced transport processes? Can
we control the formation of crystallization seeds using magnetic fields? Can we
use uncompensated spin moments of heavy 3d-metals (e.g., Fe, Co, Ni, Mn, Cr,
V) above the Curie temperature in connection with strong magnetic fields to
initiate nucleation and subsequent crystallization? Can nucleation in general be
controlled by magnetic fields?
We, therefore, chose to investigate crystallization behaviour of BHF under
the influence of a strong magnetic field (up to 5 T) inside a superconducting,

31

cryogen-free magnet (CFM) system, equipped with a precisely controllable
high-temperature furnace inside its warm bore. Experiments can be performed
starting from 0 to 5 T with the magnetic field facing upward or downward, while
the furnace allows maximum temperatures of up to 1600°C over extended
periods of time. The temperature profile over the area of maximum temperature
("hot spot") is very constant, with a maximum variation in temperature of set
point +/- 5°C. This results in a very homogeneous temperature regime over an
experimentally accessible range of 5 x 8 cm (diameter versus height). The
included ramping programs and gas-mixing equipment guarantees precise
control of all process parameters such as temperature, time and oxygen fugacity
(Fe2'/Fe3'-ratio).

2. Experiments Without Magnetic Field
2.1. Glass Crystallization Technique
The GCT traces back to investigations by Kubo et al. (1980) [4]and starts, for
our purposes, from melts prepared in the BaO-FezO3-B~O3-system as
investigated by [2], [5] and [6].Since melts prepared in this system contain Fe3+,
these melts are black, and tend to crystallize instantaneously above Fe203concentrations of more than 20 mole-%. Since crystallization should take place
under controlled conditions of temperature and time to enable us to control
chemical composition, crystallized phase, and simultaneously crystal size,
nucleation and subsequent crystal growth must be avoided during
homogenisation of the primary melt and subsequent quench to an amorphous
glass. To avoid crystallization during quench, quench rates larger than lo5 Ks-'
are necessary. These high quench rates can be obtained using a double-roller
rapid quenching apparatus as described below.
The GCT consists of basically two fusion processes: In the first step,
starting components are vigorously mixed together and homogenized on a roller
bench. This homogenized oxide mixture is then fused to obtain a primary melt,
which is quenched onto a steel plate internally cooled by refrigerated water
( 10°C). In a subsequent fusion process, obtained glass chips from the first fusion
step are added and fused inside a vertical Pt container equipped with a tiny hole
in the thin lower tip. Through this tiny hole, liquid drops of melt trickle directly
onto a double-roller rapid quenching apparatus made of 2 stainless-steel
cylinders of about 100 mm diameter, which are internally water cooled and
counter-rotating at speeds of 300 rpm or more. The gap between these two
cylinders is less than 0.1 mm and the melt drop is squeezed through this gap.

32

The heat is transferred onto the steel surfaces and the melt drop is quenched
instantaneously to a black, glassy flake. The double-roller rapid quenching
technique supplies quench rates of up to lo5 Ks-’. Consequently, all glass flakes
showed no indication of any crystallization. This finding was confirmed by
XRD measurements on powdered glass flakes.
Starting from a pristine, amorphous glassy state, these glass flakes
withstood a tempering process under temperature conditions optimum for the
growth of BHF crystals. Experiments were performed between 600 and 900°C
for a period of 2 hours and longer. At 800°C and for run durations of 2 hours,
homogeneous nanometer-size BHF crystals were obtained. Crystals are
separated from the B203-matrix by chemical treatment with diluted acetic acid
and recovered by centrifugation. With this technique, BHF powders of far less
than 500 nm diameter, a narrow crystal size distribution, a crystal aspect ratio of
3:1, and a maximum coercivity of up to 400 kA/m ([2,5,6]) were obtained.
Examples of BHF crystals are shown in Figure 1.
Powders produced by the method described above are used as reference
materials (“zero point”) for all subsequent investigations performed inside
magnetic fields. By direct comparison of results obtained under conditions
without any magnetic field present with results obtained by crystallization inside
the CFM, the potential influence of a magnetic field on the crystallization of
BHF powders can be directly studied and demonstrated.

Figure 1: REM image of BHF crystals obtained from a starting composition of 0.40 Ba0-0.27
Fe203-0.33 BzO,, tempered at 800°C for 2 hours.

33

2.2. Analytical Investigations
Presently, six compositions of the system Ba0-FezO3-BzO3 have been
investigated, and BHF powders were produced using the GCT. Since at least
Bz03represents a volatile species in our system, melt and product compositions
were checked throughout the entire GCT process by electron microprobe
analysis (EMP). This allows precise knowledge and monitoring of potential
changes in chemistry during the entire process.
Each step of the sample preparation was accompanied by extensive studies
of chemical and physical properties. Composition of the glass chips obtained
after initial fusion of oxide mixture was checked by electron microprobe
analysis for their major element composition. DTA investigations were
performed to determine optimum conditions for the formation of BHF in
successive tempering experiments. Flakes obtained from the second fusion step
and rapid quenching were checked for their amorphous state by XRD and REM
measurements prior to any further use in tempering experiments. After
tempering, powdered flakes were analysed for their crystalline phases by XRD.
A wide variety of analytical investigations such as density determinations,
REM, optical microscopy, and magnetic property measurements were routinely
performed on all powders. Based on these data and the precise record of the
experimental conditions existing during their formation, parameters controlling
the BHF crystallization can be evaluated, and optimum conditions for the
formation of high-end grade BHF powder can be determined.

2.3. Results

EMP confirmed the major composition of the mixtures. XRD investigations
showed patterns identical with pure and perfectly shaped BHF crystals. BHF
powders obtained and analysed by REM revealed in consequence as well
hexagonal platelets of single-crystalline and single-domain BHF of far less than
500 nm diameter, with a narrow aspect ratio (diameter versus thickness) of
about 3:l. This is in perfect agreement with results obtained by [1,2,5,6,7,8,9].
DTA investigations showed that BHF crystals can be obtained when annealing
(tempering) flakes above 650°C. However, perfect crystal morphology can be
obtained between 780 to 800°C. Higher temperatures lead to degradation of the
hexagonal shape of the BHF crystals due to solid-body reactions and beginning
fusion.
Increasing crystal perfection with increasing annealing temperature up to a
maximum temperature is visible in the determined magnetic properties. Results
are shown in Figure 2 for the determination of the coercivity Hc. Measurements

34

were performed at the IPHT (Jena) applying a vibrating sample magnetometer
(static measurement), measuring the B-H behavior of powdered samples,
assuming spherical particle size and Stoner-Wohlfahrt behavior without
performing a demagnetization correction. The coercivity & increases from
600°C up to 850 or 900°C depending on the chemical starting composition of

400
n

EI

-!!

xu

z

3s
u
a2

3

-.-

+35 BaO, 35 Fe a03,30 B )03
40 BaO, 25 Fe *O,,35 BIO,

I

I

I

I

-

350 - - 40 BaO, 33 Fe ,O,, 27 B aO,
300 - - 0 - - 44 BaO, 15 Fe,O,, 41 BaO,
250
200
150

-

100

-

50
01

I

I

I

I

1

I

I

I

550 600 650 700 750 800 850 900 950

Temperature [“C]
Figure 2. Coercivity (in kA/m) of BHF powders versus annealing temperature (in “C)after 2 hours
of annealing. The grey triangle represents a repetition of the same composition with an extended
tempering duration of 48 hours. Chemical composition of the starting mixtures (in mole-%) as
indicated.

the oxide mixture from which the BHF powder was obtained. The “saddle”
shaped dependence at lower temperatures (< 750°C) is most likely due to the
change of the major phase of Ba-ferrite present (monoferrite more stable at
lower temperatures, hexaferrite formed above 650°C). Similar behaviour was
observed in [7] and explained as change in size dependence of the particles from
superparamagnetic to stable single domain ones [4,10]. Chemical composition
of the initial oxide mixture has a decisive influence on the coercivity &.
Compositions with too low Fez03 content of 15 mole-% (open diamonds) do not
supply the necessary amount of Fe to form the active magnetic phases. In
consequence, coercivity & stays low and nearly constant. A starting oxide
mixture with 25 mole-% Fez03 results in an already excellent coecivity &,
while higher amounts (33 or 35 mole-%) lead to generally higher coercivities
below 800°C. A maximum coercivity of about 350 kA/m is obtained for the

35
25 mole-% composition for annealing temperatures of 850°C while
compositions with higher initial Fe20s-content have not yet reached their
maximum. Increasing the annealing time from 2 hours to 48 hours for the same
composition and annealing temperature increases the coercivity from 186 W m
to 260 kA/m (compare Figure 2: grey triangle) most likely due to an increased
crystal perfection and orientation of the magnetic domains.
The determination of the remanence MR showed nearly identical behavior in
respect to the annealing temperature. The remanence MR increases between
600°C and 900°C from about 30 mT to a maximum value of 230 mT obtained at
around 800°C with no significant dependence on chemical starting composition.
While performing aspect ratio determinations, BHF powders were deposited
on glass slides held above a BHF solution kept inside a beaker in an ultrasonic
bath. Consequently, particles were deposited such that aspect ratio
measurements could be more easily performed. A REM picture of such a
deposited BHF sample is given in Figure 3. As a consequence of this procedure,
BHF crystals start to rearrange themselves in a chain-like configuration along
their crystallographic c-axis. This facilitates the determination of aspect ratios
with a higher precision than possible from images as shown in Figure 1. On the
other hand, this image is a good example of what crystallization experiments
inside strong magnetic fields are hoped to yield: the control of seed formation
and crystal growth. After hitiation, crystals growth should result in the
formation of single domain crystals of nearly identical size and shape. Crystal
perfection should be high, and magnetic domains should be oriented parallel to
each other by the magnetic field, resulting in improved magnetic material
properties both on the nanometer as well as on the macroscopic scale.
Materials and results obtained under magnetic field-free conditions will be
used as magnetic field-free (MFF) reference material for all future materials
synthesized inside magnetic fields. Comparison of these materials will allow the
precise evaluation of the influence of magnetic fields on the crystallization
process.

36

Figure 3: REM image of BHF powder deposited applying ultrasonic techniques. BHF powder was
obtained from a starting oxide mixture with 0.40 Ba0-0.27 Fez03-0.33€3203, tempered at 800°C for
2 hours.

3. Experiments In High Magnetic Fields
3.1. Ctyogen-Free Magnet (CFM) System

Observed macroscopic magnetic properties of BHF are not exclusively
controlled by the magnetic properties of the crystals and their single domains but
as well by the relative orientation towards each other, and some other physical
properties (e.g., density or packing). A high magnetic field applied during the
entire crystallization process (initiation and formation of crystallization seeds,
crystallization with altered crystal growth and transport properties) should result
in a net increase of the magnetic properties (e.g., coercivity, remanence, energy
density).
Therefore, experiments in a cryogen-free magnet (CFM) of up to 5 T
magnetic flux density equipped with a high-temperature oven to control
precisely process parameters like temperature and tempering conditions will be
performed to investigate whether and how strong magnetic field can help to
increase the magnetic potential of BHF nano-crystalline powders especially in
respect to an optimization of magnetic properties and their industrial application.
As mentioned earlier, an increase of (BH),, of a factor of 2 should theoretically
be possible [3].
For our experiments we use a CFM system built by CRYOGENICS
(London, UK) as shown in Figure 4a and b. It can supply a maximum field
strength of 5 T in its warm bore of 300 mm diameter. The CFM is mounted on a
tiltable stand, and can be rotated into a vertical orientation of the warm bore axis

37
suitable for the installation of a high temperature furnace (HTF). Determination
of the supplied magnetic flux density between 1 and 5 T leads to a 3-D
calibration of the magnetic field inside the warm bore, and indicated that it is
constant within 0.05 T over the dimensions of the Pt crucibles (4 cm diameter
and 6 cm height) used for future crystallization experiments. These dimensions
allow for about 200 g of melt in a single experiment. This mass is more than
sufficient to guarantee an as wide range of analytical investigations of the
obtained crystallized glass samples as possible.
3.2. The High-Temperature-Furnace (HTF)
Based on this “3-D map” of the magnetic field distribution inside the warm bore
of the CFM, an HTF was designed and ordered by XENON Advanced Heating
(Freiberg, Germany). It is a specially designed, vertical muffle tube furnace that
runs at peak temperatures of 1600°C over extended periods of time, while
cooled at its exterior to guarantee the 40°C maximum temperature of the inner
warm bore surface of the CFM.
The built-in heating system guarantees a flat maximum temperature profile
inside the furnace (“hot zone”), and a nearly constant temperature condition over
the experimental charge kept in the Pt crucible of 6 cm height and 4 cm width.

Figure 4a. CFM (cryogen-free magnet) system on tilt stand power supply, temperature readout
system, and computer for automatic control in background.

38

Figure 4b. CFM rotated to vertical position and HTF inserted into warm bore.

Both hot zone and area of maximum, constant magnetic field strength are
designed to overlap in the center of the experimental region. The HTF uses a
EUROTHERM controller that allows high accuracy and reproducibility of
experimental conditions for experiments with ramp-up and ramp-down
procedures as well as for those at constant tempering temperatures. The installed
gas mixing setups will allow future experiments at constant oxygen fugacities to
additionally investigate the influence of the Fe2+/Fe3+ratio on the crystallization
behavior and involved phases of the melts investigated.
3.3. Crystallization Experiments at High Temperatures Inside High Magnetic
Fields
Serious problems with the reliability of the heating system forced us to cancel
the initiation of crystallization experiments under controlled temperature
conditions and to remodel the entire built-in heating system. The remodeling is
still in progress.
We, therefore, decided to start some qualitative experiments. Since
controlled temperature conditions are not available, we fused a BHF starting
mixture inside a standard GERO high temperature furnace at 1400°C. We then
transferred the melt contained in a Pt crucible with lid after 2 hours of fusion
into our CFM standing at 1 T. Outside of the magnetic field, the melt convected
vigorously. The damping effect of the magnetic field on the convective flow

39

while transferring the melt into the magnetic field was impressive. During this
cooling process, temperature conditions inside the sample were monitored over
1 hour, decreasing from 1400°C to 200°C. Parallel to this experiment a batch of
BHF powder of identical composition was allowed to cool naturally outside the
magnetic field under identical temperature conditions.
After 1 hour, both samples, still at roughly 200"C, were quenched to room
temperature by running tap water over the Pt crucible walls. Both samples were
then inserted into the CFM and tested for their magnetic behavior. BHF that was
crystallized outside the magnetic field showed only small magnetic properties
and could easily be moved in and out of the CFM standing at 1 T. The sample
crystallized inside this field, however, showed a strong paramagnetic behavior
when inserted into the warm bore of the CFM. Fast movements were damped
significantly in direct comparison with the sample crystallized outside the
magnetic field. Attractive forces of the magnetic field were significantly larger
in comparison to the sample crystallized without magnetic field. Unfortunately,
further quantitative investigations cannot be presumed until the HTF is delivered
and experiments are repeated under precisely known temperature conditions.
However, results of these preliminary experiments have already demonstrated
the potential of crystallization processes performed inside magnetic fields and
the effect on the magnetic properties of materials created in this way.
Acknowledgements
This study was financially supported by the Thiiringer Ministerium fiir
Wissenschaft, Forschung und Kunst (TMWFK). Many thanks go to the
Bayerisches Geoinstitut (University of Bayreuth) for performing the electron
microprobe analyses, and to Dr. Muller from the Institut fiir Physikalische
Hochtechnologie (IPHT, Jena, Germany) for helping with the vibrating sample
magnetometer measurements and their interpretations.
References
1. Watanabe, K., Hoshi, K., Crystallisation kinetics of fine barium hexaferrite,
BaFe12019,particles in a glass matrix. Phys. Chem. Glasses 40(2), (1999),
pp. 75-78.
2. Knauf, O., Nutzung groOer Abkiihlungsgeschwindigkeiten zum
Amorphisieren spontan kristallisierender oxidischer Schmelzen, dargelegt
am System BaO-Fe203-B203. Dissertation, Technische Universitat of
Ilmenau, 1988.
3. Heck, C., Magnetische Werkstoffe und ihre technische Anwendung. A.
Hitthig Verlag, Heidelberg (1975), 280 pages.

40

4.

5.
6.
7.

8.
9.

10.

Kubo, O., Ido, T., Inomato, T., Yohoyoma, H., German patent # DE 304
1960, Tokyo Shibaura Denki, 1980/81.
Hiilsenberg, D., Knauf, 0.. Hamann, B., Glass Crystallization Technique
for Ultrafine Ceramic Powders. AGM Meeting of the German Ceramic
Society, Weimar (1993), Germany.
Hulsenberg, D., Knauf, O., Hamann, B., Glass crystallization Technique for
Ultrafine Ceramic Powders. DKG 71(11-12), (1994), pp. 707-711.
Gornert, P., Sinn, E., Schiippel, W., Pfeiffer, H., Rosler, M., Schubert, Th.,
Jurisch, M., Sellger, R., Structural and Magnetic Properties of BaFelz.
&oXTixOl9 Powders Prepared by the Glass Crystallization Method, ZEEE
Transactions on Magnetics, 26 (January 1990), pp. 12-14.
Rosler, M., Gornert, P., Sinn, E., Structural and Magnetic Properties of BaFerrite Fine Particles Grown by Glass Crystallization. Z.Phys. D - Atoms,
Molecules and Clusters, 19, (1991), pp. 279-281.
Gornert, P., Schiippel, W., Sinn, E., Schumacher, F., Hempel, K.A., Turilli,
G., Paoluzi, A., Rosler, M., Comparative measurements of the effective
anisotropy field Ha for Barium Ferrites. J. Magnet. Magnet. Mat. 114,
(1992), pp. 193-201.
Shirk, B.T., Ruessem, W.R., Magnetic Properties of Barium Ferrite Formed
by Crystallization of a Glass, J. Am. Cerum. SOC. 53(4), (1970), pp. 192196.

VARIATION OF PHASE TRANSFORMATION TEMPERATURE
IN FE-C ALLOYS IN A HIGH MAGNETIC FIELD
X.J. HAO, H. OHTSUKA, H. WADA
Tsukuba Magnet Laboratory, National Institute for Materials Science,
Tsukuba, Ibaraki, 305-0003 Japan
A magnetic field can affect the transformation temperature and microstructure if a
transformed phase has a different susceptibility from the parent phase. Fe-C alloy is an
ideal system to show the magnetic field effect since, in this system, austenite (FCC
stmcture) is a paramagnetic phase and ferrite (bcc structure) is a ferromagnetic phase
below 770 "C. In this paper, phase transformation temperature in Fe-C alloys in a
magnetic field was measured from a cooling curve. It was found that the transformation
temperature for pure Fe from austenite to ferrite has a linear relationship with magnetic
field strength, increasing about 0.8 "C per tesla. For eutectoid transformation in Fe-0.8C
alloy, similar relationship exists; the transformation temperature increases about 1.5 "C
per tesla. The measured interaction energy between magnetic field and ferrite is larger
than that calculated from molecular field theory. An elongated and aligned microstructure
by ferrite transformation in a high magnetic field was found in a Fe-0.4C alloy, but was
not found in pure Fe and Fe-0.8C alloy.

1. Introduction
The expectation is that an external high magnetic field affects solidsolid phase
transformation behaviors and transformed structures, and possibly improves the
mechanical and magnetic properties of materials. In fact, the structural
alignment in solidsolid transformations in high magnetic fields has been
reported for ferrite transformation [ 1-51 and reverse transformation [6,7] by
continuous cooling or isothermal holding in austenite and ferrite dual phase
region. Thermodynamic calculation OR an equilibrium phase diagram of Fe-C
binary system proposed that a high magnetic field increases the
austenite(y)/ferrite(a) equilibrium temperature, carbon solubility in c1 phase and
eutectoid carbon content [8,9]. Kakeshita et a1 [lo] have investigated the effects
of magnetic field on martensite transformation temperature, but there are few
experimental data to confirm the effect of magnetic field on ferrite
transformation temperature. In this paper, we report on our experiments on the
effect of a magnetic field on phase transformation temperature and
microstructure in Fe-C alloys and compare them with the results of theoretical
calculations.

41

42

2. Experiments
The alloys used in this study were Fe-0.8C alloy prepared by vacuum induction
melting and high purity Fe (99.99%). After hot rolling, specimens were
machined to 5mmx 5mm xlmm and then set in a vacuum furnace, which was
installed in a helium-free-type superconducting magnet with a bore size of
@ 100 mm. A magnetic field perpendicular to the 5mmx5mm specimen's surface
was increased to 10 T in 27 minutes before austenitization and kept constant
during austenitization and subsequent cooling, then decreased to 0 T. Specimens
were fixed at the center of the magnetic field and the magnetic force on the
specimen is negligible. Specimens were austenitized at 1000 "C for 15 minutes
and cooled to 600 "Cat a cooling rate of 10 "Umin. The specimen temperature
was measured by a thermocouple contacted with the specimen and recorded by a
digital recorder. Since the temperature controller has some delay to control the
sample temperature to programmed set value if the sample temperature changes
suddenly, it is possible to find a peak in the cooling curve. This peak determined
the phase transformation temperature. Microstructure observation was
performed on the plane parallel to the direction of magnetic field by optical
microscope after polishing and 3% Nital etching.
10-

(b)
8\

Time

T(0)=906.Z°C

-

0

2

4
6
8
1
Magnetic field strength, HIT

0

Figure 1. Cooling curves (a) of pure Fe in magnetic field after austenitizationat 1000 "C for 15 min,
and the femte transformation temperature increasing with magnetic field strength (b).

43

Time

Figure 2. Cooling curves (a) of Fe-0.8C alloy in magnetic field after austenitization at 1000 "C for
15min, and the pearlite transformationtemperature increasing with magnetic field strength (b).

3. Results and Discussions
Figure l(a) shows the cooling curve segments of pure Fe during transformation
in magnetic fields. Without a magnetic field, the transformation temperature of
pure Fe from austenite to ferrite is about 906.2 T , which is about 6 "C lower
than the equilibrium temperature (912 "C). This difference is known as
supercooling and it provides the chemical driving force for transformation. The
transformation temperature increases gradually with increasing magnetic field
strength. With a maximum magnetic field of 10 T, the transformation
temperature is about 914.9 'C, which is 8.7 "C higher than with no magnetic
field. The increased temperature AT, (=T(H)-T(O), the transformation
temperature difference between field and no field) was plotted against magnetic
field strength in Figure l(b). It shows that the transformation temperature
increases linearly with magnetic field strength. The transformation temperature
increases about 0.8 "C for an increasing magnetic field strength of 1 T.
Figure 2(a) shows the cooling curve segments of Fe-0.8C alloy during
transformation in magnetic fields. The measured eutectoid transformation
temperature is 704.9 "C without applying the magnetic field. The equilibrium
eutectoid transformation temperature is 727 "C. A larger supercooling (about
22 "C) results compared to ferrite transformation in pure Fe. Eutectoid
transformation needs element redistribution in ferrite and cementite, and the
transformation temperature is relatively lower. With magnetic field, the
transformation temperature increases. With a magnetic field of 10 T, the
temperature increase to about 720.1 "C, which is about 15 'C higher than that
without a magnetic field. Figure 2(b) shows the relationship between AT and
magnetic field strength. They have a linear relationship, same as the ferrite

44
transformation. The AT is about 1.5 "C per tesla of increasing magnetic field
strength.

I

1.0-

0.0
0

"

", ,

, I , ,

200

,,,...'I

I I , .

400

,, ,,, 0 , .

. ..

--.
,,.'.I,,

600

800

, , , lTTF

1000

Temperature. "C

Figure 3. Variation of magnetization of pure Fe with temperature in an external magnetic field
calculated from molecular field theory. & is the spontaneous magnetization of Fe at 0 K. The
square symbols are calculated from measuring the transformationtemperature change (see text).

It is surprising that AT for pure Fe is more than half of that for eutectoid
transformation if we remember that ferrite is in a paramagnetic state at the
transformation temperature, whereas ferrite is in a ferromagnetic state for
eutectoid transformation. Figure 3 shows the magnetization of pure Fe without
or with an external field of 10 T calculated by molecular field theory. Magnetic
moment exists even above Curie temperature and the field-induced
magnetization at ferrite transformation temperature in pure Fe, 915 "C, is about
0.07w(&=1.74 X lo6A h , saturation magnetization of pure Fe at 0 K). Below
T,, the magnetization is composed of spontaneous magnetization (M,f) in zero
field and field-induced magnetization (Mfi). At eutectoid transformation
temperature in Fe-0.8C alloy, 720"C, total magnetization was calculated to be
0.42&, which includes a field induced magnetization of 0.07 &. The change of
Gibbs free energy of one mole pure Fe in an external field H is,

where V, is the volume of one mole Fe and
is the permeability of vacuum.
The free energy of ferrite in a magnetic field of 10 T decreases about 4.3 J/mol
at 915 "C and 47.5 J/mol at 720 "C calculated from this equation using above

45

magnetization data. The free energy change in magnetic field at 915 "C is more
than one order smaller than that at 720 "C.
If we consider that the magnetic energy has a similar effect on phase
transformation as that of chemical driving force, it is possible to measure the
magnetic energy from the transformation temperature. Figure 4 shows the
chemical driving force changes with temperature in pure Fe and Fe-0.8C alloy
calculated by Thermo-Calc. The chemical driving force of ferrite transformation
for pure Fe is 4.8 J/mol at 906.2 "C and -2.6 J/mol at 914.9 "C. So the magnetic
field provides a free energy of 7.4 J/mol to compensate the chemical driving
force (4.8-(-2.6)=7.4 J/mol). Similarly, the magnetic energy for eutectoid
transformation was calculated to be about 71 J/mol by measuring AT. It should
be mentioned that only the ferrite phase is considered during calculation, though
pearlite consists of ferrite and cementite, in which cementite is paramagnetic and
has a volume fraction of about 13%.
aaa

1

"I

;

.$ 30

n
880

890

900

910

Temperature. "C

920

930

0
680

690

700

710

720

730

Temperature. "C

Figure 4. Chemical driving force changes with temperature for ferrite transformation in pure Fe (a)
and pearlite transformation in Fe-0.8C alloy (b). The magnetic energy of the transformed phases in a
magnetic field of 10 T is shown in the figure.

The calculated magnetic energy based on molecular field theory was found
to be smaller than the value calculated from transformation temperature. There
may be two reasons for this. First, it is possible that the magnetization in the
high magnetic field calculated by molecular field theory is lower than the actual
magnetization. The square symbol shown in Figure 3 is the magnetization
calculated from transformation temperature. In fact, some research has already
proved that the magnetization near Curie temperature, calculated by molecular
theory, is smaller than experimental data. Another possibility is that a high
magnetic field not only changes the thermodynamic properties of phase
transformation but also affects the kinetic properties, such as nucleation sites,

46

interfacial migration and atom diffusion. Some experiments already show this
kind of effect, however, there are presently no conclusive results.
Though the transformation temperature in pure Fe is about 145 "C higher
than the Curie temperature, a high magnetic field of 10 T can shift it as much as
8°C. Thermodynamic calculation results shown in Figure 4 indicate that the
driving force for ferrite transformation of pure Fe increases much more slowly
with decreasing transformation temperature than that of eutectoid
transformation, so a weak magnetization and then a small magnetic energy can
induce large transformation temperature shifting. In a system, it was suggested
that if the slopes of the Gibbs free energy curve of two phases is close, it is
possible to control the phase transformation by an external magnetic field even
if both phases are not in a ferromagnetic state.
An elongated and aligned structure was found in the Fe-0.4C alloy by ferrite
transformation during slow cooling or isothermal transformation in a high
magnetic field, and it was suggested that this was due to the demagnetization
field developed in ferrite [2,3]. However, this type of structure was not found
through microstructural observation in pure Fe and Fe-0.8C alloy. For pure Fe,
this can be attributed to the fact that the magnetization at transformation
temperature is too small, and the demagnetization field is too weak. For
eutectoid transformation in Fe-0.8C alloy, the lamellar morphology and
orientation relationship between pearlite and parent austenite phase possibly
account for the lack of the abovementioned structure.
4.

Conclusions

The effects of high magnetic field on the phase transformation temperature and
microstructure in Fe-C alloys were investigated. It was found that the
transformation temperature for pure Fe from austenite to ferrite has linear
relationship with magnetic field strength, increasing about 0.8 "C per tesla. For
eutectoid transformation in Fe-0.8C alloy, a similar relationship exists; the
transformation temperature increases about 1.5 "C per tesla. The measured
interaction energy between magnetic field and ferrite is larger than that
calculated from molecular field theory. An elongated and aligned microstructure
by ferrite transformation in a high magnetic field was found in Fe-0.4C alloy,
but was not found in pure Fe and Fe-0.8C alloy.

References
1. Ohtsuka, H., Xu, Ya, and Wada, H., Materials Transactions, JIM, 41, (2000)
pp. 907-910.

47

2. Hao, X.J., Ohtsuka, H., DE Rango, P., and Wada, H., Muter. Trans., 44,
(2003) pp. 21 1-213.
3. Hao, X.J., Ohtsuka, H., and Wada, H., Muter. Trans., 44,(2003) pp. 25322536.
4. DE Rango, P., Hao, X.J., Ohtsuka, H., and Wada, H., Trans. Muter. Res. SOC.
Japan, 28, (2003) pp. 225-226.
5. Shimotomai, M., Maruta, K., Mine, K., and Matsui, M., Actu Muter., 51,
(2003) pp. 2921-2932.
6. Hao, X.J., Ohtsuka, H., and Wada, H., Trans. Muter. Res. SOC. Japan, 28,
(2003) pp. 223-224.
7. Ohtsuka, H., Hao, X.J., and Wada, H., Muter. Trans., 44, (2003) pp. 25292531.
8. Choi, J-K., Ohtsuka, H., Xu, Y., and Choo, W-Y., Scriptu. Muter., 43, (2000)
pp. 221-226.
9. Enomoto, M., Guo, H., Tazuke, Y., Abe, Y.R., and Shimotomai, M., Metull.
Muter. Trans., 32A, (2001) pp. 445-453.
lO.Kakeshita, T., Kuroiwa, K., Shimizu, K., Ikeda, T., Yamagishi, A. and Date,
M., Muter. Trans. JIM, 34, (1993) pp. 423-428.

MARTENSITIC TRANSFORMATIONIN SOME FERROUS
ALLOYS UNDER MAGNETIC FIELD
T.KAKESHITA
Department of Materials Science and Engineering, Graduate School of Engineering,
Osaka University, 2-1, Yamada-oka,Suita, Osaka 565-0871, Japan
Martensitic transformations of ferromagnetic materials are extensively influenced by a
magnetic field. The transformationtemperature changes, because the free energy between
parent and martensite phases changes under magnetic field. In a special case, the
martensite phase is induced only while a magnetic field is applied, and is transformed
back to the parent phase when the magnetic field is removed. The martensite plate formed
under magnetic field tends to align to the field direction. Furthermore, in some
ferromagnetic shape memory alloys, a giant magnetic field-induced strain of more than
1 % appears due to rearrangement of martensite variants.

1. Introduction
Martensitic transformations are generally influenced by external fields [ 1,2], of
which the magnetic field is an example [3,4]. In fact, many researchers [3-91
have examined effects of magnetic fields on martensitic transformations. We
also examined them systematically [4-81 and found many interesting
phenomena.
In this paper, we discuss our studies on the effects of magnetic fields on
martensitic transformations in some ferrous alloys; (i) effect of magnetic field on
the martensitic transformation start temperature, M,,and the validity of the
equation proposed by our group to evaluate the relation between M, and
magnetic field; (ii) Magnetoelastic martensitic transformation (maretensites are
induced only while a magnetic field is applied and are transformed back to the
parent phase when the magnetic field is removed) in an ausaged Fe-Ni-Co-Ti
shape memory alloy; (iii) morphology of martensite formed under magnetic
field; (iv) a giant magnetic field-induced strain due to rearrangement of variants
in the martensite state of Fe-Pd and Fe,Pt ferromagnetic shape memory alloys
exhibiting a thermoelastic martensitic transformation.

2. Experiment
Specimens
used
were
Fe-3 1.7at.%Ni,
Fe-24at.%Pt
and
Fe-3 1.9Ni-9.8Co-4.1Ti(at.%) polycrystals, and Fe-3 1.6Ni, Fe-3 1.2at.%Pd, and
Fe3Pt single crystals. They were prepared using a high frequency induction

48

49
furnace or by arc melting. An ordering heat treatment was made in the
Fe-24at.%Pt and its degree of order was 0.8. Single crystals of Fe-31.2at.%Pd
and Fe3Pt were grown by a floating zone method. The Fe-31.2at.%Pd was
solution treated at 1373 K followed by quenching into iced water. Fe3Pt was
solution treated at 1373 K followed by ordering treatment at 873 K.
High field magnetization measurements were performed at Research Center
for Materials Science at Extreme Conditions, Osaka University, using a pulsed
magnetic field with a maximum strength of about 31 MA/m. Magnetic
field-induced strain was measured by a sensitive three terminal capacitance
method, where specimens were mounted in a parallel-plate capacitance cell.

3. Results and Discussion
3.1. Effect of Magnetic Field on Martensitic Transformation Temperature
The martensite phase is usually induced thermally by cooling a specimen. In
some alloys, the transformation temperature Ms changes to Ms’ under a
magnetic field Hc. The solid circles of Figure 2 shows AM, (= M,’- M,) as a
function of magnetic field for the Fe-31.7at.%Ni (a) and Fe-24at.%Pt (b)
examined by pulsed magnetic field experiments. It is known from the figures
that the shift of M, increases with increasing magnetic field. We proposed the
following equation [4] to estimate the relation between the critical magnetic
field and the transformation start temperature:

AG(Ms)-AG(Ms’) =-AM(Ms’). Hc-(ll2)*x& .Hc2 + €0 .(awl

Hc. B

(I)

where AG(Ms) and AG(Ms’) represent the difference in Gibbs chemical free
energy between the parent and martensite phases at M, and M,’ temperatures,
respectively, hM(Ms’) the difference in spontaneous magnetization between the
parent and martensitic states at Ms’,yhf
the high magnetic field susceptibility in
the parent phase, 6 the volume change associated with martensitic
transformation, w the parent forced volume magnetostriction and B the parent
bulk modulus. The first, second and third terms on the right-hand side of Eq. (1)
represent the energies due to the magnetostatic, high field susceptibility and
forced volume magnetostriction effects, respectively. Based on the equation, H,
vs. M,’relations have been thermodynamically calculated for the present alloys
by using the physical quantities involved in the equation, obtained by referring
to the previous studies and by measurment in the present study [4,5]. The
calculated results are shown in Figure 1, where the dotted lines indicated with
M.S.E., H.F.E., F.M.E. and (M.S.E.+H.F.E.+F.M.E.) represent the H, vs. M,’

50

relations calculated for the magnetostatic, high field susceptibility, forced
volume magnetostriction and their total effects, respectively. The calculated
relations (M.S.E.+H.F.E.+F.M.E.) agree well with the experimental ones for
both of the alloys.

60

Fe-31.7al%Ni

Magnetic Field. ti IMAm-‘

Figure 1. Calculated and measured shift of Ms as a function of magnetic field for Invar Fe-3 1.7at%
Ni(a) and Invar Fe-24.0at.%R.(b).

3.2. Magnetoelastic Martensitic Transformation
Alloys exhibiting a thermoelastic martensitic transformation exhibit
pseudoelastic behavior due to the stress-induced martensitic and its reverse
transformation upon loading and unloading. Analogous to this behavior we
expected that an ausaged Fe-31.9Ni-9.8Co-4.1Ti (at.%) shape memory
alloy [ 151 should show field-induced martensitic and its reverse transformation
upon applying and removing magnetic field. The M,,A, and Af of the alloy used
are 127, 60 and 159 K, respectively. The latent heat of the transformation is
about 334.4 J/mol. The difference in spontaneous magnetization between the
parent and martensite phases is about 0.3 pB/atom at M,.
We applied a pulsed high magnetic field to the specimen at a temperature
above Af, 163 K( d T (= T-M,)= 36 K,T > Af). As seen in the M(r)-H(r) curve of
Figure2(a), there is no hysteresis of the magnetization when the maximum
strength is 22.22 MNm, meaning that the martensite phase is not induced. When
a higher field is applied, the rate of increase of magnetization against magnetic
field changes at H, = 23.08 MNm, as indicated with an arrow. When the
magnetic field is removed, the increased magnetization returns to the initial

51

value at about Hf=5.56 MA/m indicated with another arrow. This means that
martensitic transformation is induced at H, and its reverse transformation is
completed at Hf. These observations show that the magnetoelastic martensitic
transformation is certainly realized in the ausaged Fe-Ni-Co-Ti alloy, and such
behavior is always realized at temperatures above Af.

5.56MAIm

lo

10

Tr163KfT > A f )

20

Magnetic Field (MA/m)

Figure 2. M-Hcurves of an ausaged Fe-Ni-Co-Ti alloy at 163 K, which is just above the reverse
transformation finish temperature Af.

3.3. Morphology of Magnetic Fieldinduced Martensite
When martensite is induced thermally from the parent phase, many
crystallographic domains (variants) form nearly equivalently. Alternatively,
when it is induced by magnetic field, specific variants tend to grow
preferentially. Figure 3 shows the microstructure of the magnetic field-induced
martensite of an Fe-31.6at.%Ni alloy single crystal. The field is applied along
[110] direction, which corresponds to the horizontal direction in Figure 3. It is
clearly seen in Figure 3 that several martensite plates grow nearly parallel to the
field direction and run through the single crystal from one end to its other. This
preferential growth in the field direction is also observed by applying the
magnetic field along [loo] and [ l l l ] . The reason for the lengthwise growth
under a magnetic field is not clear, but a shape magnetic anisotropy effect seems
to play an important role.

52

Figure 3. An optical micrograph of magnetic field-induced m e n s i t e in Fe-31.6at.%Nialloy single
crystal. The magnetic field direction is horizontal and is parallel to the [I101 direction.

3.4. Giant Magnetostriction in Ferromagnetic Shape Memory Alloys
The large strain appearing in shape memory alloys is caused by the
rearrangement of variants under external stress. In some ferromagnetic shape
memory alloys, the rearrangement of variants can be induced by a magnetic
field, resulting in a giant magnetic field-induced strain [ 12-14]. Typical
examples exhibiting such behavior are Ni-Mn-Ga, Fe-Pd and Fe3Pt. Each of
them transforms from a cubic structure to a tetragonal structure and its
tetragonality ( c h ) is slightly smaller than unity. As a result, there are three
lattice correspondence variants. The easy axis of magnetization is the a-axis for
Fe-Pd alloy, and is the c-axis for Ni-Mn-Ga and Fe3Pt. These easy axes
correspond to one of the <OOl>p directions of the parent phase. In the following,
we show some results of Fe-31.2at.%Pd and Fe3Pt [15-161.
I

.

I

I

I

I

I

I

3.0
h

E 2.0

0.0
0.0

1.o
1.5
Magnetic Field, H/(MAlrn)

0.5

Figure 4 Magnetic field-induced strain of Fe-31.2Pd alloy at 77 K. Measurement was made by
applying a magnetic field along [OOI] after cooling down without a magnetic field.

53
The martensitic transformation temperature of Fe-3 1.2at.%Pd is about
230 K. A single crystal of Fe-31.2Pd was cooled down to 77 K under zero
magnetic field, then a magnetic field was applied along [OOI] direction. In this
process a large field-induced strain (expansion) of about 3% appeared as shown
in Figure 4. We observed the rearrangement of variants through an optical
microscope under a magnetic field. In association with the rearrangement of
variants, the magnetization curve shows a large hysteresis whose area is nearly
the same as the energy dissipation evaluated by a stress-strain curve. The
uniaxial magnetocrystalline anisotropy constant K,, is about -350 W/m3 at 77 K.
The large value of lKul and small value of twinning shear will be the primary
reason for the twinning plane movement under magnetic field.
The martensitic transformation temperature of Fe,Pt (degree of order 0.8) is
about 85 K. A single crystal of Fe3Pt was cooled down to 4.2 K under zero
magnetic field, then a magnetic field was applied along [OOl] direction. A large
field-induced strain (contraction) of about 2.3% appeared in this process. A part
of the strain (about 0.6%) recovers in the field removing process and it
repeatedly appears in the subsequent field applying processes. This reversible
strain is nearly three times as large as that of TERFENOL-D (Tbl.xDyxFez)[17],
which is well known as a magnetostrictive material.
5.

Summary

We have shown that magnetic field influences extremely martensitic
transformations: the martensitic transformation temperature is