TMR Research in Insulating Granular Magnetic Materials [1 ed.] 9781613247662, 9781611228670

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Copyright © 2011. Nova Science Publishers, Incorporated. All rights reserved. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

Copyright © 2011. Nova Science Publishers, Incorporated. All rights reserved. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

MATERIALS SCIENCE AND TECHNOLOGIES

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TMR RESEARCH IN INSULATING GRANULAR MAGNETIC MATERIALS No part of this digital document may be reproduced, stored in a retrieval system or transmitted in any form or by any means. The publisher has taken reasonable care in the preparation of this digital document, but makes no expressed or implied warranty of any kind and assumes no responsibility for any errors or omissions. No liability is assumed for incidental or consequential damages in connection with or arising out of information contained herein. This digital document is sold with the clear understanding that the publisher is not engaged in rendering legal, medical or any other professional services.

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MATERIALS SCIENCE AND TECHNOLOGIES

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TMR RESEARCH IN INSULATING GRANULAR MAGNETIC MATERIALS SHIGEO HONDA, NOBUYUKI HAYASHI, ISAO SAKAMOTO AND

TAMOTSU TORIYAMA EDITORS

Nova Science Publishers, Inc. New York

TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

Copyright © 2012 by Nova Science Publishers, Inc. All rights reserved. No part of this book may be reproduced, stored in a retrieval system or transmitted in any form or by any means: electronic, electrostatic, magnetic, tape, mechanical photocopying, recording or otherwise without the written permission of the Publisher. For permission to use material from this book please contact us: Telephone 631-231-7269; Fax 631-231-8175 Web Site: http://www.novapublishers.com NOTICE TO THE READER The Publisher has taken reasonable care in the preparation of this book, but makes no expressed or implied warranty of any kind and assumes no responsibility for any errors or omissions. No liability is assumed for incidental or consequential damages in connection with or arising out of information contained in this book. The Publisher shall not be liable for any special, consequential, or exemplary damages resulting, in whole or in part, from the readers’ use of, or reliance upon, this material. Any parts of this book based on government reports are so indicated and copyright is claimed for those parts to the extent applicable to compilations of such works.

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Independent verification should be sought for any data, advice or recommendations contained in this book. In addition, no responsibility is assumed by the publisher for any injury and/or damage to persons or property arising from any methods, products, instructions, ideas or otherwise contained in this publication. This publication is designed to provide accurate and authoritative information with regard to the subject matter covered herein. It is sold with the clear understanding that the Publisher is not engaged in rendering legal or any other professional services. If legal or any other expert assistance is required, the services of a competent person should be sought. FROM A DECLARATION OF PARTICIPANTS JOINTLY ADOPTED BY A COMMITTEE OF THE AMERICAN BAR ASSOCIATION AND A COMMITTEE OF PUBLISHERS. Additional color graphics may be available in the e-book version of this book. Library of Congress Cataloging-in-Publication Data TMR research in insulating granular magnetic materials / authors, Shigeo Honda ... [et al.]. p. cm. Includes bibliographical references and index. ISBN 978-1-61324-766-2 (eBook) 1. Electric insulators and insulation--Materials. 2. Granular materials--Electric properties. 3. Magnetic films. 4. Magnetoresistance. I. Honda, Shigeo, 1943TK3421.T58 2010 620'.43--dc22 2010043919 Published by Nova Science Publishers, Inc. † New York

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CONTENTS Preface Chapter 1

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Chapter 2

vii Giant Magnetoresistance in Magnetic Multilayers and Magnetic Granular Alloys S. Honda Tunneling Giant Magnetoresistance in Heterogeneous Insulating Granular Films and Magnetic Granular Alloys S. Honda and I. Sakamoto

Chapter 3

Formation of Granular Layers by Ion Implantation N. Hayashi, I. Sakamoto and T. Toriyama

Chapter 4

TMR Effect in Granular Layers by Ion Implantation N. Hayashi, T. Toriyama and I. Sakamoto

Index

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33 59 101 135

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PREFACE Artificially structured materials granular magnetic solids, as well as nanosized multilayer films are of great interest. The granules consist of nanometer sized metal particles embedded in an immiscible medium, in most cases an insulator. The nano-structured composites are known to exhibit peculiar magnetic properties, offering valuable applications in material science. A number of MLs or compositionally modulated magnetic films have been actively investigated for the purpose of developing new materials with superior characteristics or new functions such as a perpendicular magnetization, soft magnetism, or spin-flopping at low field. This book presents and reviews research in the study of insulating granular magnetic materials. Chapter 1 – Since 1980 the magnetic metal multilayers (MLs) or the compositionally modulated magnetic films have been investigated energetically for the purpose of developing new materials having good characteristics or new functions such as a perpendicular magnetization, soft magnetism, or spin-flopping at low field. Chapter 2 – As described in Chp.1, there has been extensive interest in mangetoresistance in the magnetic multilayers or the magnetic heterogeneous granular films, after the discovery of giant magnetoresistance (GMR) in Fe/Cr multilayers. A similar magnetoresistance (MR) was observed early in granular metals embedded in insulators, such as Ni-SiO2, and Co-SiO2, and also in artificial layered structures of [ferromagnetic metal/ insulator/ferromagnetic metal] junctions. Subsequently, a quite large MR ratio has been observed in tri-layer junctions such as Fe/Al2O3/Fe, and CoFe/Al2O3/Co(or NiFe), and in Co-Al-O granular magnetic films. These phenomena are caused from spindependent tunneling similar to GMR in Chp.1, and they are called tunneling giant magnetoresistance (TMR).

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viii

Shigeo Honda, Nobuyuki Hayashi, Isao Sakamoto et al.

In the films of the granular metals embedded in insulators, the electric resistivity depends strongly on a fraction of metal species. In the large metal fraction the metal grains are connected each other electrically, causing metallic conductance with low resistivity. With decreasing metal fraction, the resistivity increases and the sign of temperature coefficient of resistivity (TCR) changes from positive to negative value suggesting the nonmetallic nature below a threshold. Below the threshold the metal grains are isolated each other, as shown in Figure 17 in Chp.1 (where B has to be replaced by an insulator), and then the tunneling conductance and TMR appear. In the next section, the authors describe experimental procedures for the preparation of the heterogeneous insulating granular (Fe, Co)-SiO2 films, and for the measurements of the physical properties in the films. The electrical transportation properties will be described briefly in the section 3, based on the theory presented by Abeles and co-workers and using the experimental data obtained by the authors’ group. In the section 4, the authors will summarize the TMR effect in heterogeneous insulating granular system. Chapter 3 – As described in Chp.1, there has been extensive interest in mangetoresistance in the magnetic multilayers or the magnetic heterogeneous granular films, after the discovery of giant magnetoresistance (GMR) in Fe/Cr multilayers. A similar magnetoresistance (MR) was observed early in granular metals embedded in insulators, such as Ni-SiO2, and Co-SiO2, and also in artificial layered structures of [ferromagnetic metal/ insulator/ferromagnetic metal] junctions. Subsequently, a quite large MR ratio has been observed in tri-layer junctions such as Fe/Al2O3/Fe, and CoFe/Al2O3/Co(or NiFe), and in Co-Al-O granular magnetic films. These phenomena are caused from spindependent tunneling similar to GMR in Chp.1, and they are called tunneling giant magnetoresistance (TMR). Chapter 4 – The nanocomposites synthesized with transition elements are important for their magnetic properties, as well as the optical applications of composite layers on the basis of dielectrics with metal nanoparticles. As a matter of fact, when materials possessing long-range magnetic order are reduced in size, the magnetic order can be replaced by some other magnetic states such as superparamagnetism. Especially, nanocomposite of magnetic granules dispersed in insulating matrixes have been very fascinating because of their unique properties associated with quantum-size effects and the possibility of applications as enhanced magnetic refrigerants and high density magnetic recording media.

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In: TMR Research in Insulating Granular … ISBN 978-1-61122-867-0 Editors: Sh. Honda, N. Hayashi et al. © 2012 Nova Science Publishers, Inc.

Chapter 1

GIANT MAGNETORESISTANCE IN MAGNETIC MULTILAYERS AND MAGNETIC GRANULAR ALLOYS S. Honda Copyright © 2011. Nova Science Publishers, Incorporated. All rights reserved.

Shimane University, Hiroshima, Japan

1A. METALLIC MAGNETIC MULTILAYERS 1. Introduction 1.1. Magnetic Metal Multilayers Since 1980 the magnetic metal multilayers (MLs) or the compositionally modulated magnetic films have been investigated energetically for the purpose of developing new materials having good characteristics or new functions such as a perpendicular magnetization, soft magnetism, or spin-flopping at low field. The MLs or the compositionally modulated films can be prepared by sputtering or evaporation process by depositing two different A and B layers alternately as drawn schematically in Figure 1. In sputtered rare earth-transition metal (RE-TM) films which have been used as magneto-optical recording materials, the film composition can be easily changed by applying a negative bias voltage (VB) to the substrate during sputtering; the RE concentration will decrease due to resputtering effect at negative VB [1]. By modulating VB it is then possible to produce compositionally modulated RE-TM films. Figure 2 shows a cross-sectional

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transmission electron microscopy (TEM) image for a compositionally modulated TbCo film prepared by alternately applying a bias voltage of VB = 250 V and 0V. Here, the clear contrast indicating the compositional modulation can be observed directly. It can also be seen that the films grow with a columnar microstructure perpendicularly to the substrate [2]. Co/Pd [3] and Co/Pt [4] multilayers show high perpendicular magnetic anisotropy. The coercivity of these MLs is strongly affected by the film structure, which varies with the conditions of the substrate- or buffer layer. Figure 3 shows a typical result where the perpendicular hysteresis loops in Co/Pt MLs are affected strongly by the substrate material [5]. Here, the very high coercivity is obtained in the film prepared on a Pt buffer layer. These perpendicular magnetic MLs with high coercivity have been exploited as perpendicular magnetic recording or magneto-optical recording materials.

 

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A-

B-

A-

Figure 1. Model for the multilayer consisting of A and B layers.

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40 nm

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Figure 2. Cross-sectional TEM image for the compositionally modulated TbCo film prepared by applying VB = -250 V and 0V alternately.

In the rare-earth/transition-metal (RE/TM) MLs, the magnetic moments of both layers of RE and TM couple antiferromagnetically with each other, with an interaction that is directly perpendicular to the film plane at zero field. The Gd(0.6nm)/Co(0.6nm) ML shows a peculiar magnetization curve for the perpendicular field, which is the easy magnetization axis, as indicated in Figure 4(a). Here, a two-step magnetization process can be observed; in the plateau state the magnetic moments aligned as shown in Figure 4(b)-, and in the saturation state the magnetic moments aligned ferromagnetically as shown in (b)-. Thus, the step loop can be explained by the spin-flopping for the antiferromagnetic coupling [6]. On the other hand, the usual loop is obtained for in-plane field which is a hard-magnetization direction. In this case, the magnetic moments aligned as shown in (b)- similarly to the antiferromagnet. In the Fe/Cr or Co/Cu MLs, the antiferromagnetic interaction occurs between adjacent Fe or Co layers across the non-magnetic Cr or Cu layer as shown schematically in Figure 5(a). Grünberg et al.[7] confirmed the antiferromagnetic coupling in Fe-Cr-Fe trilayer by means of light scattering. As described in the next sub-section, the giant magnetoresistance was observed by Baibichi et al. [8] in Fe/Cr MLs exhibiting antiferromagnetic coupling.

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- 0.5

0.5 H (T)

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Figure 3. Hysteresis loops for perpendicular fields in [Co(0.4nm)/Pt (1.1nm)]10 MLs sputter-deposited on a glass and a polyimide substrate and on a Pt underlayer coated on a polyimide substrate.

1.2. Giant Magnetoresistance In 1988 the giant magnetoresistance (GMR) was discovered by Baibich et al.[8] in (001) Fe/(001) Cr bcc MLs prepared on (001) GaAs substrates, in which the magnetic moments between adjacent Fe layers are coupled antiferromagnetically at zero field as described in the last sub-section. In zero field, the magnetic moment align antiferromagnetically as shown in Figure 5(a). This induces a high electric resistivity owing to the spin-dependent scattering at the interface. When the external field increases, the magnetic moments originally pointing against the field rotate gradually to the field direction, and finally align ferromagnetically as shown in Figure 5(b). Following the alignment of the moments, the resistivity decreases gradually because of reduction of the spin-dependent scattering. This behavior is indicated in Figure 5(c), which is a data obtained by our group in Co/Cu MLs. Baibich et al.[8] observed a very large change in resistance in Fe/Cr MLs; the resistance decreased during magnetizing by a factor of 2 at 4.2 K. They called the huge value of the magnetoresistance (MR) the giant magnetoresistance (GMR). Since the discovery of the GMR, various aspects of magnetic multilayers (MLs) have been investigated by many researchers. In Fe/Cr MLs, Parkin et al.[9] investigated the effect of the Cr layer thickness on the exchange interaction between the adjacent Fe layers and on the MR ratio, and found that the exchange interaction and MR ratio change periodically with the Cr layer thickness. A similar periodic change was observed in Co/Cu MLs by Parkin et al.[10] and Mosca et al.[11] The periodic change in the exchange interaction is interpreted in terms of RKKY model.[10]

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Giant Magnetoresistance in Magnetic Multilayers …

M (emu/cm3)

 

② 

①  Gd(0.6nm)/Co(0.6nm)]110

③  - 0.2

- 0.1

0.1

0.2 H (T)

(a)

Co

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Co Gd

H

Co Gd

Gd Co Gd

① 

② 

H Co Gd

③ 

Co Gd

(b) Figure 4. Magnetization curves in a Gd(0.6nm)/Co(0.6nm) ML measured at room temperature. A peculiar loop is obtained in the easy direction of the magnetization.

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Fe(Co) Cr(Cu) Fe(Co) Cr(Cu) Fe(Co)

(a) H = 0

(b) H > Hs

high resistivity

low resistivity

(c)

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(a)

-1

- 0.5

0 H (T)

0.5

1

Figure 5. Schematic directions for the magnetization in the multilayer consisting of alternative magnetic and nonmagnetic layers; (a) shows the antiferromagnetic alignment and high resistivity at zero-field, and (b) shows the ferromagnetic alignment and low resistivity for high field. The curve (c) shows the relation between the magnetoresistance and the field for a [Co(1nm)/Cu(1nm)]30 ML prepared at VB = 0 V and PAr = 5 mTorr.

Many theoretical [12,13] and experimental [14] studies have suggested that the spin-dependent interfacial scattering plays an important role for the MR ratio. Furthermore, the GMR effect is associated with the long-range antiferromagnetic (AF) exchange interaction between the adjacent magnetic layers across the nonmagnetic metal layer, which is strongly affected by the interface structure and the crystalline orientation. Thus, the AF coupling and also the MR ratio show the strong dependence on the film structure. In this sub-chapter, we survey the relation between the exchange coupling, or the MR

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ratio, and the film morphology on the basis of our experiments for the sputtered Co/Cu multilayers [15-17]. Regarding that the morphology of sputtered films depends significantly on the sputtering conditions such as pressure [18] or dc bias voltage (VB) applied to the substrate [19], we prepared several Co/Cu MLs by changing the sputtering conditions, and we then subsequently investigated the AF coupling constant and MR ratio as a function of the sputtering condition. Furthermore, we made clear the annealing effect on the AF coupling and MR ratio [20,21].

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2. Experimental Co/Cu MLs were prepared by rf magnetron sputtering onto water-cooled glass substrates with application of a dc bias voltage VB = 0 or -30 V at an Ar gas pressure (PAr) of 3 – 130 mTorr after evacuating the chamber below 5×10-7 Torr. Sputtering was carried out using separate targets of Co and Cu whose diameters were 100 mm. Input RF power was 200 W for the Co target and 50 W for the Cu one. The deposition rates of Co and Cu were 0.1 – 0.27 nm/s and 0.1 – 0.21 nm/s, depending on the PAr and VB. The Co layer thickness (dCo) was held constant at 0.8 nm and the Cu layer thickness (dCu) was varied from 0.7 – 5.2 nm. Thirty bilayers of the Co and Cu were deposited for each sample. The film morphology and the crystalline structure were studied by crosssectional TEM and X-ray diffraction (XRD) using Cu-Kα radiation. Magnetic properties were investigated with a vibrating sample magnetometer (VSM), and magnetoresistance (MR) was measured at room temperature (RT) by dc four-terminal configuration with a transverse field of up to 0.9 T applied in the film plane as shown in Figure 6. The MR ratio was defined as ΔR/Rs where Rs is the resistance at the field of 0.9 T and ΔR is the difference of the maximum resistance from Rs. Annealing of the sample was done in vacuum of pressure lower than 1×10-6 Torr. The samples were heated with a rising rate of 40˚C/min up to 200˚C, then held there for several minutes, and finally cooled naturally to room temperature, where magnetic properties and XRD were measured, together with the temperature dependence of MR was measured in the range of 100 – 300 K. The same procedure was repeated several times, and thus the variations in the physical properties were plotted as a function of the annealing time.

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10 mm

 

H 5 mm

I V 20 mm Figure 6. Pattern of the sample for the measurement of magnetoresistance.

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3. Survey of Results 3.1. As-deposited Multilayers As reported by Parkin et al.[10] and Mosca et al.[11], the MR ratio changes periodically with the thickness of the nonmagnetic Cu layer (dCu). The results in our experiments are shown in Figure 7 for Co/Cu MLs sputtered at various PAr with VB = 0 V. When the MLs were prepared at lower pressure than 11 mTorr, well defined oscillation appears with a period of 1.0 – 1.3 nm. Particularly, the MLs prepared at 5 mTorr exhibit the maximum MR ratio of 44.5% at dCu = 1.0 nm (the 1st peak). Typical MR and magnetization curves for the MLs prepared at 5 mTorr are shown in Figure 8. For the first peak (Figure 8a), the MR curve shows only little hysteresis and the magnetization curve saturates first until about 0.4 T, which indicates the strong AF coupling. The AF coupling energy J is estimated to be 0.04 erg/cm2 from the relation

J=

1 H s M Co d Co , 4

(1)

where MCo = 1350 emu/cm3 and dCo are the saturation magnetization and the thickness of the Co layers, respectively. This magnetization curve, however, shows higher remanence value at zero field, in spite of the high J-value. This is probably due to Co atoms that have diffused into the Cu layers through the pinhole, which favors the ferromagnetic coupling between the adjacent Co layers through the bridge effect.22) We will discuss this high remanence value

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in the next sub-section. Figure 8(b) shows the curves for the 2nd peak of the MR oscillation (dCu = 2.0 nm). The MR curve shows hysteresis having with two peaks at the coercive fields and an MR ratio of about 24%. The magnetization curve indicates smaller saturation field than 0.05 T, which corresponds to J = 0.008 erg/cm2. On the other hand, for PAr≧20 mTorr, the MR ratio increases monotonically with dCu without the oscillation as shown in Figure 7, indicating that the RKKY like exchange interaction is weak. The MR curve for the MLs prepared higher than 20 mTorr displays sharp peaks at coercive fields, where the magnetic domains in adjacent Co layers are randomly arranged with each other because of the interlayer “decoupling” [22], and the GMR occurs owing to the “spin valve” effect [23]. The curve of monotonic increase for PAr≧ 20 mTorr coincides with the minimum points of the MR-oscillation for low PAr. The MR value of the valley corresponds to the GMR originating from the spin valve effect of decoupling parts. Figure 9 shows the relation between the MR ratio and dCo for the MLs prepared with a negative bias voltage of -30 V (-30V-MLs). Contrary to the 0V-MLs, the 2nd peak intensity increases with PAr going up to 11% at 64 mTorr. The well defined 2nd, 3rd and 4th oscillation peaks can be clearly observed for the -30V-MLs prepared at 64 mTorr, even though the 1st peak is not observed. In this case, the value of the valleys in the oscillation curve are smaller than those for the 0V-MLs. This suggests that the GMR mainly arises from the AF coupling of RKKY-like exchange interaction, as confirmed from the MR and magnetization curves in Figure 10, which is for the 2nd peak of the MR oscillation. The magnetization increases linearly up to the saturation field Hs = 0.03 T, and the remanent magnetization at zero field is very small, which implies that the adjacent Co layers are well coupled antiferromagnetically with a coupling constant that is estimated to be J = 0.008 erg/cm2. Corresponding to the magnetization curve, the resistivity decreases monotonically with magnetic field from the maximum value at nearly zero field. Thus, the MR characteristics depend strongly on the sputtering conditions, which cause the change in the film morphology. Typical cross-sectional TEM images are shown in Figure 11. For the 0V-ML prepared at 64 mTorr in Figure 11(a), a layer structure cannot be observed, while a porous structure consisting of columnar grains are seen. In this case, Co atoms diffuse easily into the Cu layer through the column boundaries and form a “bridge” with ferromagnetic (F) coupling [22], which diminishes the MR oscillation. On the other hand, a clear layer structure is observed to exist in the 0V-ML prepared at 5 mTorr (b) which showed the highest MR value in the 1st peak, and also in the -30V-ML

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prepared at 46 mTorr (c) which exhibited a large 2nd peak and a small remanence value in the magnetization curve.

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Figure 7. MR ratio versus Cu layer thickness for [Co(0.8nm)/Cu(dCu)]30 MLs prepared by sputtering at various Ar gas pressure PAr with VB = 0V.

-1.0

-0.5

0

0.5

1.0

H (T)

-1.0

-0.5

0

H (T)

(a) 

0.5

1.0

(b) 

Figure 8. Magnetoresistance and magnetization curves for typical sputter-deposited Co/Cu MLs of PAr = 5.1 mTorr and VB = 0V for dCu = 1.0 nm (a) and 2.0 nm (b). TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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Figure 9. MR ratio versus Cu layer thickness for [Co(0.8nm)/Cu(dCu)]30 MLs prepared by sputtering at various Ar gas pressure PAr with VB = -30V.

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- 1.0

- 0.5

0

0.5

1.0

0.5

1.0

H (T)

- 1.0

- 0.5

H ( )

Figure 10. Magnetoresistance and magnetization curves for typical sputter-deposited Co/Cu MLs of PAr = 64 mTorr and VB = -30V for dCu = 2.3 nm. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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VB=‐30V   PAr=46mTorr 

Figure 11. Cross-sectional TEM images for the [Co(0.8nm)/Cu(2.4nm)]30 MLs prepared at 64 mTorr with VB = 0V (a), 5.1 mTorr with 0V (b) and 46 mTorr with 30V.

These structure variations due to the Ar gas pressure PAr and the bias voltage VB can be explained by the kinetic energy of the sputtered atoms reaching the substrate surface and the ion assist effect as found by Yasugi et al.[2] When sputtering was done in an atmosphere of high PAr, the sputtered atoms will frequently collide with Ar gas atoms during the flight to the substrate, and the kinetic energy of the atom reaching the substrate will become too small for the atoms to move on the grown film surface. Thus, the columnar structure is formed by the shadowing effect in high PAr as reported by ref. [2], and this also induces the rough film surface, which interfere with formation of a clear and smooth layer structure. With decreasing PAr, the

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kinetic energy of the atoms reaching the substrate will increase permitting the atoms to move on the grown film surface thus leading to formation of a smooth surface and the fine film structure as reported in ref. [2]. Thus, the clear layer structure is established in the 0V-ML prepared at 5 mTorr. At very low pressure, however, the sputtered atoms with very high energy will bombard the grown film surface and re-sputter the atoms on the film surface, which induces the damage in the layer structure. By applying the negative bias voltage, the film structure becomes fine and the smooth layer structure is established owing to the ion assist effect even at high pressure, as found in Figure 11(c). However, the diminishing of the 1st peak in -30V-ML prepared at 46 mTorr cannot be explained at present. With decreasing PAr, strong ion bombardment induces the resputtering of the film surface, resulting in the damage in the layer structure.

Figure 12. MR ratio versus AF coupling constant J for the 1st and 2nd oscillatory peaks for Co/Cu MLs. ○,●: 0V, 3 mTorr, △,▲: 0V, 5.1 mTorr, □,■: 0V, 11 mTorr, ▼: 30V. The open marks represent the 1st peak and the solid marks represent the 2nd peak. The broken lines show the values obtained from refs. 4 and 5.

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As seen in Figs.7 and 9, the MR ratios for the 1st and 2nd peaks changes with the sputtering conditions. The values are plotted in Figure 12 as a function of the AF coupling constant J. The data are well fitted on each line for the 1st and 2nd peaks, which indicate a linear relation between the MR ratio and the AF coupling. In this figure the data for 1st peak at 4.2 K obtained by Parkin et al.[10] and Mosca et al.[11] are also shown. Our data show the larger gradient, which indicates that the large MR ratio is obtained at low field compared with MLs in refs. [10] and [11].

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- 1.0

- 0.5

0

0.5

1.0

H (T)

Figure 13. Magnetoresistance curves for the as-deposited and annealed [Co(0.8nm)/Cu(0.9nm)]30 ML prepared at 5.1 mTorr with 0V. The annealing was carried out at 200˚C for 120 min.

3.2. Annealing Effect Figure 13 shows the MR curves for the [Co(0.79nm)/Cu(0.94nm)]30 ML prepared at 0V and 5 mTorr, which is the sample near the 1st peak of the GMR oscillation [21]. Here, both the states of as-deposited and subsequently annealed ML at 200˚C for 120 min are shown. The MR ratio steeply decreased from 30.5% to 20% by annealing for about 10 min, and then gradually to about 9 % after 120 min annealing. The decrease of the MR ratio is associated with the change in the shape of the magnetization curves as shown in Figure 14, where the easy magnetization axis lies in the film plane [21]. Upon annealing the saturation field Hs for the easy direction significantly to below 0.1 T from about 0.6 T, and the remanent magnetization at zero field increased near the saturation magnetization.

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⊥ 

∥ 

∥ 

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- 1.5

- 1.0 - 0.5

0

⊥ 

0.5

1.0

1.5

H (T) Figure 14. Magnetization curves for the same MLs as Figure 13.

The film structure was investigated by X-ray diffraction (XRD) during annealing. Figure 15 [21] shows the small-angle (SA) and high-angle (HA) XRD for the same sample as Figs.13 and 14. The HA-XRD shows a diffraction peak near the fcc-Co(111) and Cu(111) peaks, and the SA-XRD coming from the multilayer periodicity shows only the first-order peak corresponding to the periodicity of 1.73 nm (= dCo + dCu). Figure 15 does not show the significant change in both the HA- and SA-XRD by annealing, although the MR ratio and the magnetization curves vary markedly as found from Figures 13 and 14. For the 1st peak (dCu = 1.0 – 1.1 nm) in the other samples prepared at PAr = 3 – 11 mTorr with 0V, we examined also the annealing effect on the MR ratio and the XRD structure as a function of the annealing time, giving the results shown in Figure 16 [20]. Figure 16(a) indicates that the drastic decrease in the MR ratio at the early stage occurs for all the samples investigated. The annealing time dependence of the intensity I1 in the 1st peak of SA-XRD is shown in Figure 16(b). This figure shows that the decrease in I1 is very small

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for the annealing times used, even though the data display are some scatters. Thus, the long-range interface structure seems to be retained as shown by XRD investigation. However, the AF exchange interaction varies drastically by annealing.

Figure 15. Small-angle (SA) X-ray diffraction patterns for the same ML as Figure 13.

The decrease in the MR ratio due to annealing cannot be explained by the XRD data. Saito et al.[21] examined the interface structure for both the asdeposited and annealed films using nuclear magnetic resonance (NMR) spectra, which is sensitive to the interface structure. The NMR spectra indicated that both the as-deposited and annealed films have pinholes across neighboring layers and a number of Co atoms diffuse through the pinholes. These Co atoms may lead to ferromagnetic “bridges” across neighboring layers causing the masking [24,25] of the AF coupling, which results in the larger remanence magnetization as found in Figs.8 and 14. The NMR spectra clearly demonstratethat the number of Co atoms in the Cu layer near the boundary increases with annealing. Thus, the change in atomic structure in the layer boundaries causes the drastic decrease in AF coupling constant and also in the MR ratio.

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  (a) 

(b) 

Figure 16. Variations in the MR ratio (a) and the 1st SA-XRD intensity (b) for the 0V Co/Cu MLs of the 1st peaks of the oscillatory GMR as a function of annealing time at 200˚C.

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Magnetic metal

Non-magntic metal

A (Fe, Co, Ni)

B (Ag, Cu, etc.)

Figure 17. Model of a granular film.

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i j

(b) H < Hs

High resistivity

(a) H > Hs

Low resistivity

i

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(c)

θi

= 2 = [M(H)/Ms]2

φij j 

Figure 18. Schematics for the magnetic moments in the granules. (a) is for a field that is lower than the saturation field Hs, and (b) is for a field that is higher than Hs. (c) is the directions of the moments i and j.

4. Summary In metallic multilayers consisting of alternate ferromagnetic and nonmagnetic layers, the giant magnetoresistance can be observed when the antiferromagnetic (AF) exchange interaction occurs between the neighboring ferromagnetic layers across the nonmagnetic layer. The AF coupling is strongly affected by the film morphology, especially on the atomic structure in the layer boundaries.

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1B. GRANULAR METALLIC FILMS

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1. Introduction When two different metals A and B, which are mutually insoluble in the phase diagram such as Co-Cu, Fe-Cu, Co-Ag and Fe-Ag, are co-deposited by evaporation or sputtering with small fraction of A, phase separation occurs, and small particles or granules of A embedded in the matrix B are formed as shown in Figure 17. Such a heterogeneous alloy is called a granular system. If the metal A is ferromagnetic and B is a nonmagnetic metal such as CoCu and Fe-Ag, the small single-domain ferromagnetic granules are formed. The size and density of the granules are dictated by the volume fraction of A. In system with a very small fraction of A, the nano-sized granules remain isolated magnetically in the nonmagnetic matrix B, and then the magnetic moment of each granule directs randomly at low field as shown in Figure 18(a). In this case, the electric resistance is high, because the spin-dependent scattering takes place at the interface separating magnetic and nonmagnetic phase or inside the magnetic granules. With increasing external field, the magnetic moment aligns gradually to the field direction as shown in Figure 18(b), which reduces the electric resistance. This phenomenon is similar to the GMR in multilayered systems. Thus, a large MR appears in the granular systems such as Co-Cu [26,27], Fe-Cu [27], Co-Ag [28,29] and Fe-Ag [30]. In a granular system, the GMR is the extra resistance due to spindependent scattering from nonaligned ferromagnetic moments and the ratio of GMR is proportional to the nonaligned degree , where φ ij is the angle between the moments i and j as shown in Figure 18(c). When the moments of the granules are uncorrelated, the relation = 2 = [M(H)/M s]2 is held, where θ i is the angle between the moment i and the external field H as shown in Figure 18(c), and M(H) and Ms are the magnetization in field H and the saturation magnetization, respectively. Thus, the resistivity in field H, ρ(H), is proportional to (M/Ms)2 as reported by Xiao et al. [27] and can be expressed by

⎡ ⎛M ρ ( H ) = ρ s + ρ m ⎢1 − ⎜⎜ ⎢⎣ ⎝ M s

⎞ ⎟⎟ ⎠

2

⎤ ⎥, ⎥⎦

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(2)

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where ρ s and ρ m are constants independent of H. The random alignment of the moments occurs in a system where the magnetic granules are isolated magnetically or are in a superparamagnetic state. These situations occur for small volume fractions of magnetic atoms. Zhang [31] and Asano et al. [32] showed theoretically the effect of the granular size on the MR ratio based on a model in which the spin-dependent scattering occurs primarily at the surface of the magnetic granules, and indicated that the MR ratio is inversely proportional to the diameter of the magnetic granules for a fixed composition. The inverse relation has been observed experimentally during annealing processes by several authors [3335]. However, the GMR is very sensitive to some of the preparation parameters [36], and to the magnetic phase such as the ferromagnetism or superparamagnetism. When the sample is annealed, the granules grow in size and the magnetic phase of small granules changes from the superparamagnetism to the ferromagnetism, which causes the decrease in the MR ratio [37]. Therefore, the results in the annealing experiments is not always accepted as a proof for the theoretical prediction. In this sub-section, we survey the granular-type GMR and the magnetic properties in sputtered Co-Ag films, as a function of the film structure and both the size and density of magnetic Co granules based on our experimental results [37,38].

2. Experimental Heterogeneous Co-Ag films were prepared on glass substrates by rf sputtering at 5 – 20 mTorr Ar gas pressure (PAr) from a composite target of Ag plates placed asymmetrically on a Co disk. The asymmetry in the target composition induced a gradient of the composition in the films prepared on the glass substrates, and thus we obtained samples having a range of compositions in a single sputtering run. In order to change the grain size in the as-prepared films, we applied a dc bias voltage (VB) to substrates from +50 to -150 V. The deposition rate was 0.25 – 0.7 nm/s dependent on PAr and VB, and the total film thickness ranged from 150 to 420 nm. Film structure was examined by XRD, and the composition was analyzed by Rutherford backscattering (RBS) method. Magnetic properties were investigated using a vibrating sample magnetometer (VSM), and the MR ratio was measured with in-plane field up to 1.5 T at room temperature. Annealing of the sample was carried out in a vacuum lower than 1×10-6 Torr.

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3. Survey of the Results 3.1. As-deposited State The magnetization curves for the Co-Ag films are shown in Figure 19 as a function of Co content (x). The films of x = 45 – 80 at% exhibit similar loops to the magnetic bubble materials for perpendicular fields [Figs.19(a) and (b)], suggesting a significant perpendicular magnetic anisotropy and the existence of stripe magnetic domains at zero field. With decreasing x, the magnetic anisotropy decreases gradually and the magnetization curves exhibit the superparamagnetic nature of nonhysteresis and unsaturation for x < 25 at% [Figure 19(d)], because the blocking temperature TB averaged over the film decreases to much lower than room temperature.

- 1.5

- 1.0

1.0

1.5

- 1.5

- 1.0

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H (T)

- 1.5

- 1.0

1.0

1.5

1.5

1.0

H (T)

- 1.5

- 1.0

H (T)

1.0

1.5

H (T)

Figure 19. Magnetization curves for the Co-Ag films with various compositions prepared at 5 mTorr with 0 V.

In the films of x = 25 - 45 at%, the magnetization curves for both in-plane and perpendicular fields are characterized by the superposition of the ferromagnetic and superparamagnetic curves [Figure 19(c)].

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- 1.5 - 1.0

- 0.5

0

H (T)

0.5

1.0

1.5

Figure 20. Experimental magnetization curve (solid line) for the Co-Ag film prepared at 5 mTorr with 0 V having 44 at% Co composition, and calculated curves of Eq.(3), where FM (dotted line) and SP (broken line) are for the ferromagnetic and superparamagnetic components, respectively, and total (dash-dotted line) is the sum of FM plus SP. Here, the parameters used are MFMs = 125 emu/cm3, MFMr = 10.5 emu/cm3, Hc = 189 Oe, MSPs = 240 emu/cm3 and μ = 4600 μ B.

The magnetization curve such as Figure 19(c) can be split into ferromagnetic and superparamagnetic components using the formula reported by Stearns and Cheng [39]:

M(H) =

s 2MFM

π

r ⎧ ⎛ μH ⎞ ⎛ μH ⎞−1 ⎫⎪ ⎧⎪H ± Hc ⎛ πMFM ⎞⎫⎪ s ⎪ ⎟ ⎜ tan ⎨ tan⎜ s ⎟⎬ + MSP ⎨coth⎜ ⎟ − ⎜ ⎟ ⎬ , ⎪⎩ Hc ⎪⎩ ⎝ kT ⎠ ⎝ kT ⎠ ⎪⎭ ⎝ 2MFM ⎠⎪⎭ −1

(3) where MFMs, MFMr and Hc are the saturation magnetization, remanence and coercivity of the ferromagnetic component, and MSPs and μ are the superparamagnetic saturation magnetization and magnetic moment of a granule for the superparamagnetic phase, respectively. Figure 20 shows an example in the film of x = 44 at% deposited at 5 mTorr with 0 V: the dotted and broken lines indicate the first and second terms of Eq.(3), respectively, and the summed curve (dash-dotted line) coincides with the experimental curve (solid line). The parameters used are described in the figure caption.

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Thus, we can obtain the saturation magnetization of each component for various compositions. The data of MFMs and MSPs for films prepared at 10 mTorr with 0 V are shown in Figure 21(a). This shows that the superparamagnetic component is dominant for x < 40 at%. The resistivity change ρ(0) – ρ(1.5) for applying the maximum field of 1.5 T is plotted in Figure 21(b). Figure 21 shows an analogy between the composition-dependences of s MSP and Δρ max. Therefore, in Figure 22 we plotted the values of Δρ max and MR-ratio as a function of the saturation magnetization of the superparamagnetic component, MSPs for the samples prepared at 10 mTorr and 0 V. This figure indicates that the Δρ max and MR-ratio are proportional to MSPs, and suggests that the the granular type GMR comes mainly from the superparamagnetic component and depends on the density of the superparamagnetic granules. In this figure, however, the solid and broken lines, obtained by the least square method, intersect not the origin but the abscissa at positive values. This reason will be discussed below.

ρ (0) - ρ (1.5) (mΩcm)

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Figure 21. (a) Saturation magnetization of ferro- (□) and superparamagnetic (△) components analyzed and (b) maximum resistivity change as a function of Co composition for Co-Ag films prepared at 10 mTorr and 0 V. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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S. Honda

As described in the introduction, Zhang [31] and Asano et al. [32] predicted theoretically that the MR-ratio is proportional to the inverse of the granular radius. Here, we will check the theoretical prediction by calculating the average radius (rg) from the average volume of the granules (Vg) by assuming spherical shape for the granules;

⎛ 3V g rg = ⎜⎜ ⎝ 4π

⎞ ⎟⎟ ⎠

1/ 3

.

(4)

The average volume Vg can be calculated by assuming that the granules have the fcc structure with a lattice constant of a0 = 0.354 nm and have the same magnetic moment as bulk Co atoms of 1.72 μ B:

max (μΩ -cm),

Δρ

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MR (%)

V g = N Co a 03 ,

(5)

16 14 12 10 8 6 4 2 0

10mTorr, 0V MR

Δρmax

0

100

200 s

300

3

M SP (emu/cm ) Figure 22. Maximum resistivity change Δρ max = ρ(0) – ρ(1.5), and the maximum MR ratio Δρ max/ρ(1.5) versus superparamagnetic saturation magnetization MSPs for Co-Ag films prepared at 10 mTorr and 0V.

where

N Co =

μ 1.72μ B

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(6)

Giant Magnetoresistance in Magnetic Multilayers …

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is the average number of the Co atoms in a superparamagnetic granule per volume, μ is the magnetic moment of the superparamagnetic granule obtained in Eq.(3), and μ B is the Bohr magneton. Then, we can calculate the radius (rg) using Eq.(3) for all the samples, and we can plot the relation between the MR ratio and the inverse of radius (1/rg). However, the MR-ratio is proportional to rg1/3 as discussed below rather than 1/rg, which is contrary to the theoretical prediction proposed by Zhang [31] and Asano et al [32]. The density of the superparamagnetic granules (Ng) can be estimated using the superparamagnetic saturation magnetization (Msps); s N g = M SP /μ

(7)

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We therefore plotted in Figure 23 the superparamagnetic component of the resistivity-change (Δρ sp) for the samples prepared at 5 mTorr as a function of the product

Figure 23. Δρ SP versus VgNg for Co-Ag films prepared at 5 mTorr and various bias voltage. Here, the data after annealing are plotted using solid circles.

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Figure 24. Annealing effect for the Co-Ag film with x = 37% prepared at 5 mTorr and 0 V. The column (a) shows the magnetization curves, and (b) shows the MR curves. s M SP a03 μ s Vg ⋅ N g = a ⋅ = M SP . 1.72μ B μ 1.72μ B 3 0

(8)

Here, the value of Δρ SP is obtained by removing the hysteretic component from the total resistivity-change Δρ. Figure 23 indicates again that Δρ SP is proportional to the product of VgNg, namely the superparamagnetic saturation magnetization MSPs, as shown in Figure 22, independent of the preparation condition, even though the grain size and film structure are severely affected by those conditions. In this figure, the data after annealed are plotted also by the solid circles, and they fit well on the linear line. The annealing effect will be described at the next sub-section. In Figure 23 the line of least square fitting intersects not the origin but the abscissa at positive values similarly to Figure 22. This is presumed to be a result of the fact that the probability of the spin-dependent scattering becomes very small, resulting in small value of GMR, when Ng and/or Vg are very small, because of the decrease in the collision probability between the conduction electron and the granule.

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3.2. Annealing Effect In the previous experiments for annealing effects on MR [33-35], the linear relation was reported between the MR ratio and the inverse of the granular radius (1/rg) for a constant Co volume fraction. We also examined the annealing effects on MR for the samples prepared at 5 mTorr and 0 V. As shown in the column (a) in Figure 24, the magnetization curve changes from an isotropic nature, characterizing superparamagnetism, to an anisotropic nature originating from ferromagnetism by annealing. This indicates that the granules grow in size during annealing and the size of some grains overcomes the threshold for phase transition of magnetism. On the other hand, the MR ratio becomes smaller with annealing as indicated in the column (b) in Figure 24. Using the procedure described at the last sub-section, we split the magnetization curve into the ferromagnetic and superparamagnetic components, and calculated the radius rg of the superparamagnetic granules. The variations in rg and Δρ SP due to annealing are plotted in Figure 25 for several compositions x. As found in Figure 25, the radius rg increases with annealing, and the value of Δρ SP decreases. As plotted in Figure 26, Δρ SP is proportional to 1/rg with the relation of Δρ SP = C1/rg + C2 for each composition x, where the constant C1 decreases slightly with x. The proportional relation is consistent with the previous reports [33-35]. However, we have to pay attention to the fact that the density of the superparamagnetic granule is reduced by annealing owing to the phase transition from superparamagnetism to ferromagnetism with grain growth as described above. Therefore, it cannot always be certified that the decrease in ΔρSP is due to the reduction of the granular surface on annealing. The data for the annealed sample are plotted also in Figure 23 by the solid points, and they fit well on the line. This indicates strongly that the MR ratio in granular type GMR depends on the superparamagnetic saturation magnetization or the density of the superparamagnetic granule.

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Figure 25. Variations in the radius rg and Δρ SP for the superparamagnetic component of Co-Ag films by annealing. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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Figure 26. Relation between Δρ SP and 1/rg for annealed Co-Ag films.

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CONCLUSION When small magnetic granules are embedded in a non-magnetic metal, the GMR occurs by the mechanism of spin-dependent scattering, that is proportional to the degree of correlation of the magnetic moments of neighboring granules. The high MR ratio can be obtained in the system, in which the magnetic moments align randomly at zero field. This situation is performed in the superparamagnetic granular film, and it is exhibited that the MR ratio is proportional to the saturation magnetization of the superparamagnetic phase.

REFERENCES [1]

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[2] [3] [4] [5] [6] [7] [8]

[9] [10] [11] [12] [13] [14]

S.Honda, M.Ohkoshi and T.Kusuda, IEEE Trans. Magn. MAG-22 (1986) 1221. S.Yasugi, S.Honda, M.Ohkoshi and T.Kusuda, J. Appl. Phys. 52 (1981) 2298. P.F.Carcia, A.P.Meinhaldt and A.Suna, Appl. Phys. Lett. 47 (1985) 178. P.F.Carcia, J. Appl. Phys. 63 (1988) 5066. S.Honda, H.Tanimoto and T.Kusuda, IEEE Trans. Magn. 26 (1990) 2730. M.Nawate, H.Kriake, K.Doi, S.Honda and T.Kusuda, J. Magn. Magn. Mater. 104-107 (1992) 1861. P.Grünberg, R.Schreiber, Y.Pang, M.B.Brodsky and H.Sower, Phys. Rev. Lett. 57 (1986) 2442. M.N.Baibichi, J.M.Broto, A.Fert, F.Nguyen van Dau, F.Petroff, P.Etienne, G.Creuset, A.Friderich and J.Chazelas, Phys. Rev. Lett. 62 (1988) 2472. S.S.P.Parkin, Z.G.Li and D.J.Smith, Appl. Phys. Lett. 58 (1991) 2472. S.S.P.Parkin, R.Bhadra and K.P.Roche, Phys. Rev. Lett. 66 (1991) 2152. D.H.Mosca, F.Petroff, A.Fert, P.A.Schroeder, W.P.Pratt Jr. and R.Laloee, J. Magn. Magn. Mater. 94 (1991) L1. P.M.Levy, S.Zhang and A.Fert, Phys. Rev. Lett. 65 (1990) 1643. J.Inoue, A.Oguri and S.Maekawa, J. Phys. Soc. Jpn. 60 (1991) 376. Y.Saito, K.Inomata, A.Goto and H.Yasuoka, J. Phys. Soc. Jpn. 62 (1993) 1450.

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Giant Magnetoresistance in Magnetic Multilayers … [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]

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[25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37]

31

S.Honda, T.Mimura, S.Ohmoto and M.Nawate, IEEE Trans. Magn. 28 (1992) 2745. S.Honda, S.Ohmoto, R.Imada and M.Nawate, J. Magn. Magn. Mater. 126 (1993)419. S.Honda and M.Nawate, Proc. 3rd Intl. Symp. Magn. Mater. Process & Devices (1994) 189. Z.G.Li and P.F.Carcia, J. Appl.Phys.71 (1992) 842. S.Honda, H.Tanimoto, J.Ago, M.Nawate and T.Kusuda, J. Appl. Phys. 70 (1991) 6047. M.Nawate, K.Inage, R.Imada, M.Itogawa and S.Honda, Jpn. J. Appl. Phys. 34 (1995) 3082. Y.Saito, K.Inomata, M>Nawate, S.Honda, A.Goto and H.Yasuoka, Jpn. J. Appl. Phys. 34 (1995) 3088. M.Vohl, J.A.Wolf, P.Grünberg, K.Spörl, D.Weller and B.Zeper, J. Magn. Magn. Mater. 93 (1991) 403. B.Dieny, V.S.Speriosu, S.S.P.Parkin, D.A.Gurney, D.R.Wilhoit and D.Mauri, Phys. Rev. B43 (1991) 1297. U.Gradmann, H.J.Elmers and J.Kohlhepp, Mater. Res. Symp. Proc. 313 (1993) 107. S.S.P.Parkin, R.F.Marks, R.F.C.Farrow, G.R.Harp, Q.H.Lam and R.J.Savoy, Phys. Rev. B46 (1992) 9262. A.E.Berkowitz, M.J.Carey, J.R.Michell, A.P.Young, S.Zhang, F.E.Spada, F.T.Parker, A.Huttenn, Phys. Rev. Lett. 68 (1992)3745. J.Q.Xiao, J.S.Jiang and C.L.chien, Phys. Rev. Lett. 68 (1992) 3749. J.A.Bamard, A.Waknis, M.Tan, E.Halftek, M.R.Parker and M.L.Watson, J. Magn. Magn. Mater. 114 (1992) L230. J.Q.Xiao, J.S.Jiang and C.L.Chien, Phys. Rev. B46 (1992) 9266. J.Q.Wang, P.Xiong and G.Xiao, Phys. Rev. B47 (1993) 8341. S.Zhang, Appl. Phys. Lett. 61 (1992) 1855. Y.Asano, A.Oguri, J.Inoue and S.Maekawa, Phys. Rev. B49 (1994) 12831. P.Xiong, G.Xiao, J.Q.Wang, J.Q.Xiao, J.S.Jiang and C.L.Chien, Phys. Rev. Lett. 69 (1992) 3220. J.Q.Wang and G.Xiao, Phys. Rev. B49 (1994) 3982. P.Auric, S.R.Tiexeria, B.Dieny, A.Chamberod and O.Redon, J. Magn. Magn. Mater. 146 (1995) 153. J.Q.Wagn, E.Price and G.Xiao, J. Appl. Phys. 75 (1994) 6903. S.Honda, M.Nawate, M.Tanaka and T.Okada, J. Appl. Phys. 82 (1997) 764.

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S.Honda, T.Okada and M.Nawate, J. Magne. Magn. Mater. 165 (1997) 326. M.B.Stearns and Y.Cheng, J. Appl. Phys. 75 (1994) 6894.

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[39]

S. Honda

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In: TMR Research in Insulating Granular … ISBN 978-1-61122-867-0 Editors: Sh. Honda, N. Hayashi et al. © 2012 Nova Science Publishers, Inc.

Chapter 2

TUNNELING GIANT MAGNETORESISTANCE IN HETEROGENEOUS INSULATING GRANULAR FILMS AND MAGNETIC GRANULAR ALLOYS S. Honda1 and I. Sakamoto2 Copyright © 2011. Nova Science Publishers, Incorporated. All rights reserved.

1

Shimane University, Hiroshima, Japan 2 Hosei University, Tokoyo, Japan

1. INTRODUCTION As described in Chp.1, there has been extensive interest in mangetoresistance in the magnetic multilayers or the magnetic heterogeneous granular films, after the discovery of giant magnetoresistance (GMR) in Fe/Cr multilayers [1]. A similar magnetoresistance (MR) was observed early in granular metals embedded in insulators, such as Ni-SiO2 [2.3], and CoSiO2 [4], and also in artificial layered structures of [ferromagnetic metal/ insulator/ferromagnetic metal] junctions [5]. Subsequently, a quite large MR ratio has been observed in tri-layer junctions such as Fe/Al2O3/Fe [6], and CoFe/Al2O3/Co(or NiFe) [7], and in Co-Al-O granular magnetic films [8]. These phenomena are caused from spin-dependent tunneling similar to GMR in Chp.1, and they are called tunneling giant magnetoresistance (TMR). In the films of the granular metals embedded in insulators, the electric resistivity depends strongly on a fraction of metal species. In the large metal

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S. Honda and I. Sakamoto

fraction the metal grains are connected each other electrically, causing metallic conductance with low resistivity. With decreasing metal fraction, the resistivity increases and the sign of temperature coefficient of resistivity (TCR) changes from positive to negative value suggesting the nonmetallic nature below a threshold. Below the threshold the metal grains are isolated each other, as shown in Figure 17 in Chp.1 (where B has to be replaced by an insulator), and then the tunneling conductance and TMR appear. In the next section, we describe experimental procedures for the preparation of the heterogeneous insulating granular (Fe, Co)-SiO2 films [9], and for the measurements of the physical properties in the films. The electrical transportation properties will be described briefly in the section 3, based on the theory presented by Abeles and co-workers [3,10.11] and using the experimental data obtained by our group. In the section 4, we will summarize the TMR effect in heterogeneous insulating granular system.

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2. EXPERIMENTAL Heterogeneous granular Fe-, Co- and/or FeCo-SiO2 films were prepared on glass substrates by rf magnetron sputtering at 10 mTorr Ar gas pressure from composite targets consisting of Fe, Co or FeCo plates placed asymmetrically on a SiO2 disk, which caused the gradient in the film composition and then brought us many samples of various compositions with one single sputtering run. The atomic composition of Fe, Co and Si in the sample was measured by Rutherford backscattering spectroscopy (RBS), and then the volume fraction x of Fe, Co or FeCo was calculated using bulk densities of Fe, Co and SiO2. The crystalline structure of the granule was examined by X-ray diffraction (XRD) method with Cu-K radiation and the film structure was observed by transmission electron microscopy (TEM). The magnetic properties were measured at room temperature (RT) with a vibrating sample magnetometer. The magnetoresistance (MR) was obtained using a dc method with in-plane magnetic field.

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35

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3. ELECTRICAL TRANSPORTAT IN HETEROGENEOUS INSULATING GRANULAR FILMS The resistivity at zero field ρ(0) for (Fe,Co)-SiO2 sputtered films are plotted as a function of volume fraction x in Figure 1, where the data for CoSiO2 films presented by Barzilai et al. [4] are shown by dotted line. The resistivity increases logarithmically with decreasing x, and abrupt variations are observed at around x = 0.44 for Fe-SiO2 and 0.60 for Co-SiO2. This suggests that the transportation-type changes at these thresholds. The reason, why the threshold depends on the kind of metal, is not clear at present. Probably, this is caused by the geometric structure of the granules due to the different crystalline structure of metal; Fe is bcc and Co is hcp structure. Figure 2 shows a temperature dependence of ρ(0) for Fe-SiO2 films. In x > 0.44 [(c) and (d)] the resistivity increases with temperature as observed commonly in metals, while it decreases with temperature for x ‫ أ‬0.44 [(a) and (b)] as observed in amorphous films such as GdCo sputtered films [12] and in semiconductors. In order to clarify the origin of the change in electrical phase at the threshold, we investigated the film structure using XRD and TEM. Figure 3 shows typical XRD patterns for Fe-SiO2 as a function of x. For larger x, bcc Fe (110) peak appears at 2ρ = 44.5˚. With decreasing x, the peak intensity weakens and the width broadens, indicating that the Fe grain size becomes smaller. Here, the grain diameter d, estimated using Scheller’s formula, is indicated. The diameter d decreases from about 9.5 nm at x = 0.64 to about 6.4 nm at x = 0.52. For x < 0.40, the peak intensity becomes too small to be detected. This indicates that the Fe grain becomes gradually small with decreasing x and the isolated Fe granules are formed for x < 0.40. Figure 4 shows the TEM image for the Fe-SiO2 film of x = 0.45. In this figure, we can see the chain of small Fe grains. It is suggested that the electric current is transported through small channel from the grain to grain. This conductance is referred to percolation. The resistivity increases with decreasing x because the channel becomes narrower. For x < 0.44, the grains are isolated with each other and then the transportation-type changes from metallic to tunneling or hoping conduction. The composition limit giving this conduction change is called to be a percolation threshold. The tunneling conductivity in the films of granular metals embedded in the insulator has been discussed by Abeles and co-workers [3,10,11]. Here, we describe roughly the discussion presented in refs.3, 10 and 11. In these papers,

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S. Honda and I. Sakamoto

the granular metal was assumed to be a sphere with diameter d and separated distance s from the neighbor as shown in Figure 5. In order to transfer an electron between two neutral grains, a charging energy is required. The charging energy Ec is approximately given by

e2 e2 s 2e 2 = F (ε , ) ≈ , C d d Kd

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Ec =

(1)

where C is the capacity of the system of two grains, e is the charge of an electron, ε is the dielectric constant of the insulator, F is the function of ε and s/d, and K is an effective dielectric constant of the system. The charging energy Ec is required to generate charge carriers, and then the carrier density n is proportional to the Boltzmann factor {exp[-Ec/kT]}1/2, where the square root comes from the same reason as for the formulation of the density of both positive and negative carriers formed in intrinsic semiconductors having the energy gap Eg, {exp[-Eg/kT]}1/2. When an electron tunnels through the barrier having the height φ and the thickness s, the translation probability D is proportional to exp[-2χs], where χ = [2m/ћ 2]1/2 with m denoting the electron mass, and ћ = h/2π is Planck’s constant [13]. Thus, the total conductivity σ is written by  

4)

Figure 1. Resistivity ρ(0) for (Fe, Co)-SiO2 films measured at room temperature as a function of metal volume-composition x. The broken line is data presented in ref.4). TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

Tunneling Giant Magnetoresistance …

Fe-SiO2

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Figure 2. Temperature dependences of resistivity ρ(0) at zero field for Fe-SiO2 films with various volume fractions x. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

37

38

S. Honda and I. Sakamoto

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Fe-SiO2

Figure 3. X-ray diffraction patterns for Fe-SiO2 films having various volume fraction x, measured with Cu-K radiation. The vertical solid line indicates the peak position corresponding to the (110) planes of bulk Fe. The value of D means the size of Fe granules evaluated from the (110) peak width.

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Tunneling Giant Magnetoresistance …

39

 

Fe 45 vol%

20 nm

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Figure 4. TEM image of Fe-SiO2 film of x = 0.45. ∞

∞

0

0

σ ∝ ∫ β ( s ) ⋅ n ⋅ Dds = ∫ β ( s ) exp[ −2 χs − =∫

∞

0

c β ( s ) exp[ 2 χs − ]ds 2 χskT

Ec ]ds 2kT

,

(2)

where

c ≡ χsEc

(3)

is a constant, and β(s) is the density of percolation paths. Finally, we can write the temperature dependence of the conductivity σ [11];

σ = σ 0 exp[ −2 (c / kT ) ] ,

(4)

where σ 0 is a constant independent of temperature. Here, we abbreviated the details for deriving this equation.

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S. Honda and I. Sakamoto

 

e

-

s

d

Figure 5. Schematic model assumed for calculation in refs. 11) and 12). Fe-SiO2

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Figure 6. Relationship between ρ(0) and 1/T1/2 for Fe-SiO2 films with various volume fractions x. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

Tunneling Giant Magnetoresistance …

41

  

Ni-SiO2

Fe-SiO2 12)

x

(vol.%)

Figure 7. Values of the constants 1/ρt (open circles) and 1/ρc (squares) as a function of x calculated using Eq.(5) for Fe-SiO2 films. The solid circles are the data for Ni-SiO2 films presented in ref.12).

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  Ni-SiO2

Fe-SiO2 12)

x

(vol.%)

Figure 8. Tunneling activation energy c as a function of x for (Fe, Co)-SiO2 films. Here, the data for Co-SiO2,4) and Au-Al2O3,11) are plotted also. The solid line indicates Eq. (6) with η = 1 eV presented in ref.11).

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S. Honda and I. Sakamoto

To validate Eq.(4), we re-plotted in Figure 6 logarithmically the resistivity ρ(0) in Figure 2 as a function of (1/T)1/2 for the region of x ‫أ‬0.45. Contrary to  Eq.(4), the data show a nonlinearity, indicating that the other effect is included in the conductance. We found that the data fit well on the relation:

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σ ( 0) =

⎧ 1 1 c ⎫ 1 . = exp ⎨− 2 ⎬+ ρ ( 0) ρ t kT ⎭ ρ c ⎩

(5)

The equation indicates the parallel circuit consisting of the tunneling conduction of Eq.(4) and the constant conductance independent of temperature,1/ρc , even though the origin of ρc is not clarified. Thus, we can obtain the constants 1/ρt, 1/ρc and c as a function of x. They are plotted in Figures 7 and 8. For low x, the value of 1/ρt is much larger than 1/ρc, indicating the tunneling conduction to be dominant. However, 1/ρc increases drastically with x, and becomes the same order as 1/ρt for x ~ 0.45, which suggests that 1/ρc comes from the current flowing through the narrow channel connecting between the adjacent grains, and this current dominates total conduction for x > 0.50. The value of 1/ρt also increases exponentially with x, indicating good coincidence with Ni-SiO2 granular films [11]. In Figure 8 the data of the parameter c for Co-SiO2 and FeCo-SiO2 are also plotted. As shown here, the data agree well with the values obtained previously for Co-SiO2,4) and Au-Al2O3 [11]. A theoretical expression for c(x) was calculated in ref.11 by assuming a model, in which the granular metals construct a simple cubic lattice of spheres of diameter d with a lattice constant, s + d (see Figure 4), and was given by

(s / d ) 2 c =η , (1/ 2) + (s / d )

(6)

where

2 χe 2 s π . = ( )1 / 3 − 1 , and η = ε d 6x

(7)

The solid line in Figure 8 is the best fit to Eq.(6) with η = 1 eV. This value of η, however, is far small compared with the theoretical value; e.g., η = 75 TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

Tunneling Giant Magnetoresistance …

43

was obtained from Eq.(7) for SiO2 matrix [11]. This discrepancy may lie in the over-simplification for the irregularly shaped and arranged metal granules [11]. Fe-SiO2

  

(a)

x = 0.16 - 1.0

0

1.0

(b)

x = 0.26

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- 1.0

0

1.0

(c)

x = 0.37 - 1.0

0

1.0

(d)

x = 0.45 - 1.0

0

1.0

H (T)

Figure 9. In-plane magnetization curves for Fe-SiO2 films of various volume fractions x measured at room temperarure. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

44

S. Honda and I. Sakamoto

 

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Fe-SiO2

Figure 10. Saturation magnetization Ms defined from the magnetization values at 1.5 T in-plane field as a function of x for Fe-SiO2 films. The open circles are the experimental data measured at room temperature, and the solid circles are the values extrapolated to H = ∞ using Eq.(8). The solid line indicates that the magnetic moment per Fe volume MFe is constant 1714 emu/cm3.

4. MAGNETIC PROPERTIES Typical magnetization curves measured at RT are shown in Figure 9 for Fe-SiO2 films. The films of x ‫ أ‬0.37 [(a) – (c)] show the nonhysteretic and unsaturating curves which are characteristics in the superparamagnetism. This is consistent with M ssbauer observation reported by Mitani et al. [14]. The superparamagnetism comes from the small granules isolated magnetically from the neighbors in these x-regions as mentioned in the last section. On the other hand, for x larger than percolation threshold 0.44, the magnetization is saturated easily as shown in Figure 9 (d), indicating the ferromagnetism. The magnetization at 1.5 T measured at RT are plotted in Figure 10. For x ‫ ؤ‬0.45, the data fit well on the linear line. This indicates that the Fe grains have the same saturation magnetization as pure -Fe bulk, 1714 emu/cm3

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Tunneling Giant Magnetoresistance …

45

independent of x; The result is consistent with M ssbauer data [14]. For x ‫أ‬ 0.43, however, the data are slightly smaller than the line. This is probably due to the unsaturation at 1.5 T because of the superparamagnetism, or the existence of the magnetically dead granules due to oxidation at the surface [15]. The magnetization curves exhibiting the superparamagnetic nature for x ‫أ‬ 0.43 can be well expressed by the summation of two Langevin functions, indicating that the grain size is distributed around two groups of the smaller and larger ones. Figure 11 shows an example analyzed for the experimental curve for x = 0.26 using two Langevin functions;

M ( H , T ) = M Ss [coth

μS H kT

−

μ H kT kT ] + M Ls [coth L − ] μS H μ L H (8) kT

= M Ss LS ( H , T ) + M Ls LL ( H , T )  

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Fe-SiO2

- 1.5

- 1.0 - 0.5

0

H (T)

0.5

1.0

1.5

Figure 11. Experimental and calculated magnetization curves for Fe-SiO2 film of x = 0.26. The circles indicate the experimental data. The solid and broken lines are the calculated values of Eq.(8) for total and S and L components, respectively.

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46

S. Honda and I. Sakamoto

where the subscripts S and L mean the smaller and larger grains, respectively, and MSs and MLs are the saturation magnetizations, μS and μL are the magnetic moments of unit magnetic cell of smaller and larger granules, respectively. T is the temperature, and k is the Boltzmann constant, In Figure 11, the open circles indicate the experimental data. The broken lines indicate each term in the right-hand side of Eq.(8), and the solid line indicates the summation of two broken lines, namely the calculation of Eq.(8). Here, the values of parameters used are shown in the figure. From the X-ray diffraction (Figure 3) and Figure 10, we can suggest that the superparamagnetic granules are bcc crystalline structure and they have the magnetic moment of 2.22 μB per Fe atom. Then, the volume of each grain, Vi (i = L or S) can be estimated by the relation, V i = μi/2.22 μB·a03, where a0 = 0.287 nm is the lattice constant of bcc Fe. Assuming a spherical grain structure, the grain radius is given by ri =[Vi/(4π/3)]1/3. From Figure 11, rS = 1.49 nm and rL = 2.51 nm can be estimated. This indicates that the size of Fe granule is distributed around 1.4 – 2.5 nm, in consisted with the TEM observation. Thus, we can estimate the values of MSs, MLs, rS, and rL for various x as shown in Figure 12. In Figure 12(a), the linear line is the plot of saturation magnetization estimated from the simple volume dilution of Fe with the magnetic moment MFe = 1714 emu/cm3, and the circles are the sums of MSs and MLs, the saturation values of the superparamagnetic moments, which are plotted also in Figure 10 by solid circles. The saturation values are slightly smaller than the line, indicating the existence of the magnetically dead layers at the granular surfaces. Furthermore, the saturation magnetization of the smaller granules MSs depends only slightly on x, while that of the larger granules MLs increases linearly with x, and dominates the x-dependence of the total magnetization. Figure 12(b) indicates that the size of smaller granule is around 1.5 nm in radius independent of x, while the size of the larger granule increases from 1.5 nm to 5.5 nm in radius with increasing x. Similar results were obtained for Co-SiO2 and FeCo-SiO2 films. The saturation magnetization are plotted in Figure 13 as a function of metal compositions. They fit approximately on the linear lines indicating the simple dilution of the magnetic metal component, even though the data show slightly smaller values in the region of 0.60 ‫ أ‬x ‫ أ‬0.80. The superparamagnetic curves for x < 0.55 are analyzed using Eq.(8) similarly to Fe-SiO2, and the saturation magnetizations of larger and smaller granules, MLs and MSs, were observed as shown in Figure 14, which indicates that the data behave parallel to Fe-SiO2;

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47

that is, the MLs increases with x, while the MSs is approximately constant independently of x.

Fe-SiO2

 

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(a)

(b)

Figure 12. Analyzed data using Eq.(8) for Fe-SiO2 films: (a) Magnetizations, MSs, MLs, and MSs + MLs, (b) radii of granules, rS and rL, for S and L granules, which are plotted by the squares and triangles, respectively. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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S. Honda and I. Sakamoto

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Figure 13. Saturation magnetization Ms defined from the magnetization values at 1.5 T in-plane field as a function of metal composition for Fe-SiO2, Co-SiO2 and FeCo-SiO2 films.

5. TUNNELING MAGNETORESISTANCE The magnetoresistance (MR) curves were measured by applying in-plane field at RT, and the typical curves for Fe-SiO2 films are shown in Figure 15, which are for the same samples as those used in Figure 9. Similarly to the magnetization curves of Figures 9(a) – (c), the MR curves also exhibit the nonhysteresis and unsaturation characteristics. When x increases, the field sensitivity of MR at low fields increases, and the MR ratio increases up to about 3.6% at x = 0.37 [Figure 15(c)]. However, the MR ratio becomes again very small for x larger than the percolation threshold x = 0.44 as shown in Figure 15(d). Thus, the large MR occurs in the tunneling transport region exhibiting the superparamagnetism. In other words, the large MR comes from tunneling process between neighboring granules with superparamagnetism isolated electrically by insulating SiO2 matrix. In this case, the MR is cited as tunneling magnetoresistance (TMR).

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49

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Figure 14. Calculated values of MSs and MLs as a function of metal composition for FeSiO2, Co-SiO2 and FeCo-SiO2 films.

The TMR in insulating granular films was examined theoretically by Inoue and Maekawa [16]. According to them, the constant σ0 in Eq.(4) depends on the external field H, and the conductance σ is rewritten by

σ ( H , T ) = σ 0 ' (1 + P 2 m 2 ) exp[ −2 c / kT ] ,

(9)

where σ0’ is the constant independent of H, and P is related to spin polarization in ferromagnetic metals, and given by

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50

S. Honda and I. Sakamoto Fe-SiO2

  

(a)

x = 0.16 - 1.0

0

1.0

(b)

x = 0.26

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- 1.0

0

1.0

(c)

x = 0.37

- 1.0

0

1.0

(d)

x = 0.45 - 1.0

0

1.0

H (T) Figure 15. MR curves at room temperature for Fe-SiO2 films of various x. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

.

Tunneling Giant Magnetoresistance …

P=

D↑ − D↓ D↑ + D↓

51

.

(10)

Here, D↑ and D↓ are the state-densities for the up and down spins at the Fermi energy. The value m is the relative magnetization of the granular system and given by

m 2 =< cos φ ij > ,

(11)

where φ ij is the relative angle between magnetic moments of paired grains i and j (see Figure 18 in Chp.1). The MR ratio is given by

Δρ ( H , T ) ρ ( H , T ) − ρ (0, T ) σ ( H , T ) −1 − σ (0, T ) −1 = = ρ (0, T ) ρ (0, T ) σ (0, T ) −1 (12) σ (0, T ) 1 =1 − = 1− . σ (H ,T ) 1 + P2m2

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R( H , T ) =

 

Fe-SiO2

- 1.5

- 1.0

- 0.5

0

H (T)

0.5

1.0

1.5

Figure 16. Experimental and calculated MR curves for Fe-SiO2 film of x = 0.26. The circles indicate the experimental data. The solid and broken lines are the calculated values of Eq.(14) for total and S and L components, respectively, using the parameters obtained from Figure 11.

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S. Honda and I. Sakamoto For the small value of P, the above relation is approximated as follows

R( H , T ) ≈ 1 − (1 − P 2 m 2 ) = P 2 m 2 .

(13)

Thus, the MR curve is proportional to the square of the magnetization curve in granular film. However, the experimental MR curve did not fit on the relation of Eq.(13), while it fitted well on the sum of the square of each Langevin function for smaller and larger granules, namely

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R( H , T ) =

Δρ ( H , T ) = − AS [ LS ( H , T )]2 − AL [ LL ( H , T )]2 , ρ (H ,T )

(14)

where LS(H,T) and LL(H,T) are the Langevin functions for smaller and larger granules given in Eq.(8). The constants AS and AL are the MR ratios for H = ∞ arising from the smaller and larger granules, respectively. Figure 16 shows a typical example for the TMR analysis using the parameters obtained in Figure 11. Here, the open circles indicate the experimental data, the broken lines show the two terms of Eq.(14) and the solid line is the sum of two terms. We can see the good agreement between the experimental data and the calculation. The relation of Eq.(14) suggests that the film might consist of the double layers having the smaller and larger granules as sketched in Figure 17. This structure might be caused by the annealing effect because of temperature rise of the substrate during sputtering.

 

substrate

Figure 17. Schematic model for the granular films prepared in this experiment. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

Tunneling Giant Magnetoresistance …

53

Fe-SiO2

  

(a)

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(b)

Figure 18. Analyzed data of MR ratio using Eq.(14) for Fe-SiO2 films as a function of x. (a) values at H = 1.5 T; the open circles are experimental data, the squares, triangles and solid circles are calculated values for S- and L-components, and sum of them, respectively. (b) values extrapolated to H = ∞.

The MR ratio measured at RT with maximum field Hm = 1.5 T is plotted by open circles in Figure 18(a) as a function of x. The MR ratio increases with x and reaches the maximum 3.6 % at x ~ 0.37 and then decreases rapidly. This decreasing phenomenon is similar to the GMR in granular metallic films described in Chp.1, in which this is considered to be caused by the phase-

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54

S. Honda and I. Sakamoto

change from superparamagnetism to ferromagnetism. For the present case, however, the decrease in TMR comes mainly from the change in the conduction-mechanism at the percolation threshold (x = 0.44). In this figure, the results analyzed with Eq.(14) are also plotted; the squares and triangles are the components from the smaller and larger granules, and the sum of two components are shown by the solid circles, which agree well with the experimental data.

Fe-SiO2

  

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(a)

(b)

Figure 19. Relation between AS (AL) and MSs (MLs) for Fe-SiO2 films.

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Figure 18(b) indicates the values of As, AL and AS + AL calculated for H = ∞. The values of AS are slightly dependent on x corresponding to those of MSS in Figure 12(a), while AL increases with x until 0.37, similarly to x-dependence of MLs in Figure 12(a). This suggests that the MR ratio is proportional to the saturation magnetization of superparamagnetic granules, namely to the density of smaller and larger granules, in consistent with the discussion in Chp.1. Therefore, we plotted in Figure 19 the MR ratio versus Ms. For the larger granules, the data fit well on the linear line of the coefficient αL = AL/MLs = 0.0058. For the smaller granules, the coefficient is about 2.4 times as large as the larger granules; αS = AS/MSs = 0.014. The values of AL and AS correspond to the square of the spin polarization, P2, in Eq.(13). However, the real value of the spin polarization depends on the granular density or the separation between surfaces of neighboring granules, as found above. For the same atomic density, the separation is narrower for the smaller granules than the larger granules. Therefore, the spin polarization between smaller granules becomes higher than for the larger granules, as shown in the obtained relation α S < αL.

Δρ (1.5)/ρ 0 (% )

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Figure 20. MR ratios at 1.5 T as a function of metal composition for Fe-SiO2, Co-SiO2 and FeCo-SiO2 films.

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S. Honda and I. Sakamoto

Figure 21. Calculations of MR ratio at H = ∞ for smaller and larger granules, AS and AL, in Fe-SiO2, Co-SiO2 and FeCo-SiO2 films.

Co-SiO2 and Fe38Co62-SiO2 granular films show similar behavior in MR ratio versus x to Fe-SiO2 films as shown in Figure 20. Because the percolation threshold is different from each other, the peak position of MR ratio is also different from each other in their films. The peak for Co-SiO2 films appears at the largest composition x = 0.45 and shows higher MR ratio of 3.8% than FeSiO2 films. The FeCo-SiO2 films achieve the highest value 4.5% at x = 0.43. The coefficients AS and AL are calculated also for Co-SiO2 and FeCo-SiO2 films, and the data are shown in Figure 21. They behave similarly to those of Fe-SiO2. Therefore, we plotted again the relation between AL and MLs in Figure 22. Good linear relation is obtained similarly to Figure 19, and the

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Tunneling Giant Magnetoresistance …

57

coefficients, αL, are given by 0.0062 and 0.0083 for Co-SiO2 and FeCo-SiO2. They are larger than that of Fe-SiO2. The average ratios of AS and MSs, αS, are also obtained to be 0.015 and 0.019 for Co-SiO2 and FeCo-SiO2. Thus, the highest values for α L and α S are obtained in the FeCo-SiO2 films. This is caused from the highest value of the spin polarization P in the FeCo alloy.

A L (%)

3

α FeCo = 0.0083 FeCo Co

2

α Co = 0.0062

1

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0 0

100

200

300

40

M Ls (emu/cm3) Figure 22. Relation between AL and MLs for Co-SiO2 and FeCo-SiO2 films.

CONCLUSION We prepared heterogeneous granular (Fe,Co)-SiO2 films on glass substrates by rf magnetron sputtering, and investigated their magnetic and transportation properties. Their properties changed drastically with Fe and Co contents, x, at the percolation points. For the content below the percolation point, the Fe, Co or FeCo grains were isolated electrically with each other and the electrical conductance changed to the tunneling process characterized by the activation energy constant c proportional to both the tunnel-barrier thickness and the charging energy of the metallic grain. The value c increased with decreasing x without regard to granular systems (Fe-SiO2, Co-SiO2, FeCo-SiO2, Au-Al2O3 and W-Al2O3).

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S. Honda and I. Sakamoto

In the tunneling region with the metal content smaller than the percolation threshold, the Fe, Co or FeCo granules showed the superparamagnetic nature. Similarly to the magnetization curve, the MR curves exhibited the unsaturating and non-hysteretic characteristics caused by the giant magnetoresistance due to the tunneling between the superparamagnetic granules. The MR ratio was proportional to the saturation magnetization of the superparamagnetic granules and showed the maximum value near the percolation threshold. The maximum MR ratio at room temperature was observed to be about 3.5 for Fe-SiO2, 3.8 for Co-SiO2 and 4.5 for Fe38Co62-SiO2.

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REFERENCES [1] M.N.Baibich, J.M.Broto, A.Fert, F.Nguyen Van Dau, F.Petroff, P.Etienne, G.Creuzet, A.Friedrich and J.Chazelas, Phys. Rev. Lett. 61 (1988) 2472. [2] J.I.Gittleman, Y.Goldstein and S.Bozowski, Phys. Rev. B5 (1972) 3609. [3] J.S.Helman and B.Abeles, Phys. Rev. Lett. 37 (1976) 1429. [4] S.Barzilai, Y.Goldstein, I.Balberg and J.S.Helmann, Phys. Rev. B23 (1981) 1809. [5] S.Maekawa and U.Gafvert, IEEE Trans. Magn. MAG-18 (1982) 707. [6] T.Miyazaki and N.Tezuka, J. Magn. Magn. Mater. 151 (1995) 403. [7] J.S.Moodera, L.R.Kinder, T.M.Wong and R.Meservey, Phys. Rev. Lett. 74 (1995) 3273. [8] H.Fujimori, S.Mitani and S.Ohnuma, Mater. Sci. Eng. B31 (1995) 219. [9] S.Honda, T.Okada. M.Nawate and M.Tokumoto, Phys. Rev. 56 (1997) 14566. [10] P.Sheng, B.Abeles and Y.Arie, Phys. Rev. Lett. 31 (1973) 44. [11] B.Abeles, P.Sheng, M.D.Coutts and Y.Arie, Adv. Phys. 24 (1975) 407. [12] S.Honda, M.Ohkoshi and T.Kusuda, J. Magn. Magn. Mater. 35 (1983) 238. [13] L.I.Schiff, “Quantum Mechanics” Mcgraw-Hill Book Comp. New York (1955). [14] S.Mitani, H.Fujimori, S.Furukawa and S.Ohnuma, J. Magn. Magn. Mater. 140-144 (1995) 429. [15] S.Honda, T.Shimizu, Y.Une, M.Sakamoto, K.Kawabata and T.Tanaka, J. Appl. Phys. 94 (2003) 4279. [16] J.Inoue and S.Maekawa, Phys. Rev. B 53 (1996) R11927.

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Chapter 3

FORMATION OF GRANULAR LAYERS BY ION IMPLANTATION N. Hayashi1, I. Sakamoto2 and T. Toriyama3 1

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Kurume institue of Technology, Fukuoka, Japan 2 Hosei University, Tokyo, Japan 3 Musashi Institute of Technology, Tokyo, Japan

1. INTRODUCTION Material engineering at nanometer scale could provide smaller technological devices than those currently available. During the last decade the composite systems of metal nanoclusters embedded in an insulating or dielectric matrix have been demonstrated to be very fascinating because of their unusual optical, electrical, and magnetic properties, and their wide potential applications; for example, they could be used for fabricating all optical functional devices [1]. The possibility to process light signals without converting them to electronic form should, in principle, allow all-optical devices to operate in a time range faster than the electronic ones. The potential optical applications of composite layers on the basis of dielectrics with metal nanoparticles include high-speed nonlinear optical switches, directional couplers, waveguides, and so on. Furthermore, in addition to the optical applications the nanocomposites synthesized with transition elements are important for their magnetic properties [2, 3].

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Thermodynamically nanocrystalline metal particles are metastable and chemical active due to the large surface/volume ratio, and need a protective coating to resist oxidation when exposed to ambient atmosphere. The use of an insulating matrix as a host for nanocrystalline particles has been confirmed to be an effective means for stabilizing the metal nanoclusters against oxidation and offering the possibility of tailoring particles size, size distribution and homogeneity of the dispersion of the ultrafine clusters. Surface layers and films including metal-ceramic nanocomposite can be prepared in various ways. One of the most promising fabrication methods is the implantation of metal ions, because it allows us to gain a high metal filling factor in a solid matrix beyond the equilibrium limit of solubility. By choosing beam energy and ion fluence it is possible to form magnetic layers with controlling the depth, thickness and configuration. In addition, ion implantation is completely compatible with the current matured Si semiconductor technology and offers possible future applications with a combination of memories and devices. In this chapter we make a description of some case studies on metal composite-ceramics obtained by either ion implantation or ion irradiation of multilayer films, pointing out the flexibility and capability of these techniques from a point of view of tailoring granular structure. The structure, size and size distribution of granules depend on the ion species and on their reactivity with the substrate, and, for a given substrate, on the various preparation parameters. Furthermore, it will be presented that the size control is essential in improving the physical properties such as giant magnetoresistance in magnetic naocomposite materials. Therefore, we can say from a viewpoint of application purposes that the materials designer has various opportunities for controlling the composite properties. Some schematic diagrams on a variety of ion-beam based methods are presented in Figure 1 referring to the creation of nanocluster composites used in this article. Figures 1(A) and 1(B) sketch the process of direct implantation with one or two sorts of metallic ions into insulating oxide matrixes, respectively. It should be noted in Figure 1(B) that alloying of two ion species could be achieved in ion implantation even if they are immiscible in thermal equilibrium. In Figure 1 (C) the irradiation with energetic ions in MeV range are used to tailor the structure of the granules which were prepared by the preceding implantation. Many experimental studies on the irradiation effects have been reported to modify the size, density and radial distributions of the nanoparticles [4]. Figure 1(D) shows a scheme of ion irradiation to metal-doped matrixes or metal/oxide multilayered substrates for the controlled formation of metal nanoparticles. In the last two schemes the irradiation leads to the precipitation of metal species through the

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process of mixing, ballistic displacement, migration, and clustering. Ion implantation is inherently a non-equilibrium process, and the response of the target materials and the behavior of implanted ions in the solids are influenced by the materials’ chemical and structural nature. The interaction of an energetic ion beam with an insulator differs significantly from that with a metal in some respects. Whereas the lattice in metals with free electrons is flexible to some extent and resistant to radiation damages, it is likely that the insulators whose bonding between lattice atoms is rigid are more easily damaged after the atoms are exposed to displacement due to the elastic collisions with the incident beam. Therefore, it is important to obtain valuable data on the properties and amount of defects to understand structural changes in the implanted materials, which leads to the better understanding of the behavior of implanted impurities in the matrices. On the other hand, the implanted impurities, if they are probe of any analyzing tool, can also give significant data on the defects remained in materials through the association of defect-impurity complexes. Although various methods including TEM (Transmission electron microscopy), XPS (X-ray Photoelectron Spectroscopy) and EXAFS (Extended X-ray Absorption Fine Structure) have been used to study the valence states and the aggregation characteristics of the implanted ions in insulating solids, in the case of iron implantation Mössbauer spectroscopy can afford the most definitive information [5]. Through the measurements of hyperfine parameters and the spectral line profiles one can study the crystalline configurations surrounding the sites of Mössbauer probe 57Fe in the implanted insulators. We have prepared Fe and Fe-based alloy granules in various oxides matrixes [6]. The production of magnetic nanocomposites layers was achieved through the implantation with high fluences of Fe and/or Co ions. Figure 2 demonstrates a typical example of conversion electron Mössbauer (CEMS) spectrum taken from the 57Fe implanted SiO2 whose implantation condition is given in Table 2 (sample A) [7]. In the bottom of the figure the decomposed lines representing various iron sites are denoted by bar diagram. The granule formation with the particles’ size of about 6 nm in diameter (estimated by glancing XRD patterns) is depicted by a central peak at near zero velocity (Fe0, superparamagnetic, SPM peak) originating from α phase iron. This presents that magnetically split lines in metallic α-iron (ferromagnetic, FM peaks) were collapsed to one peak due to superparamgnetic relaxation. In the bar diagram three doublets coming from iron oxides are also presented. The details are discussed in section 3.1.

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(A) Single ion implantation

(B) Sequential implantation  by double ions beam 

(D) Ion beam mixing

(C) Irradiation with high  energy ions beam 

Figure 1. Ion beam based methods for forming metal and alloy clusters formation. The process (B) is for alloy formation. The process (C) shows post- implantation irradiation without mass injection of second ions, and in the process (D) multi-layered films are irradiated and mixed by the ion beam.

Table 1. Calculation of the projected ranges for some typical combination of implanting ions and matrixes Projectiles

Energy (keV)

Matrix

Fe+ Co+ Fe+ Fe+

100 100 100 74

Al2O3 Al2O3 MgO SiO2

Projected Range (nm) 49 51 60 54

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SPM

SPM

Fe2＋

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FM 

Velocity Figure 2. (a) CEMS spectra from Fe implanted crystalline SiO2 composites with fluences of 1.0 × 1017 ions/cm2 at 100 keV and additional 0.4 × 1017 ions/cm2 at 140 keV (correspond to the sample A in Table 2). (b) Bar diagrams of the spectrum corresponding to (a) spectrum.

2. EXPERIMENTAL DETAILS FOR CHARACTERIZING TMR GRANULES Most granular solids are composed of common metals such as Fe, Au, Co, Cu, etc., and insulators such as Al2O3, SiO2, MgO, and so on. The granular solids can be synthesized by a variety of deposition techniques; above all sputtering deposition is one of the most versatile methods as described in

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Chapters 1 and 2. Although it is not so common as sputtering deposition, ion implantation has many unique merits that the technique can be realized with almost any metal species, and it can be performed to allow the direct designing of the structural geometry. The production of buried magnetic granular layers is achieved through the implantation of an amount of magnetic metal species [8]. To produce magnetic nanocomposites crystalline or amorphous oxide substrates were implanted with a mass analyzed beam to high fluences of (1~3.5) × 1017 ions/cm2, using a 400 keV ion implantor (Nissin Ion Equipment Co.,LTD, NHV79). With such high fluence, for example, the implantation of 2 × 1017 ions/cm2 gives rise to average atomic concentrations of about 30 at. % over the straggling range (in percentage of Fe per cations). The crystalline target samples were tilted 5 o ~ 7o with respect to the ion beam to avoid channeling implantation. The implantation was mostly carried out at room temperature under vacuum condition of 10-4 Pa or less. For the most research works the ion current density is kept at a level of 1~ 5 μA/cm2 in order to suppress temperature rise of the targets during implantation. An ion source of hollow cathode type is adopted for 57Fe implantation to prevent wasting source materials because the iron isotope is expensive. The ion beam is scanned over the sample surface in order to assure a homogeneous implantation. Since granular solids of nano-alloy-particles are expected to display a variety of interesting physical properties [2], sequential implantation of two species of metal ions was also performed to synthesize binary alloy granules. The alloy composition is adjusted by setting the amount of injected ions to a desired fluence. Moreover, ion beam technology has been used to modify the surface properties. Besides the direct injection of metal species, ion beams can be used for promoting cluster aggregation in the surface layers of oxides previously doped with other metal or multilayered films. High energy ion-beam mixing and ion-beam assisted deposition have also been successfully exploited for the preparation of cluster-doped oxides [9] and the tailoring of the granular structures. Some of these schemes are shown in Figure 1. In most implantation works the beam energies were used in the ranges of 105 eV to incorporate the ions in the near surface regions of the targets. The estimation of the ion penetration depth was done by TRIM (Transport of Ions in Matter) code [10], and typical examples of the implantation data concerned with the present review are summarized in Table 1. For the samples prepared with moderate dose rate the mean depth values obtained experimentally from RBS (Rutherford Backscattering Spectroscopy) measurements were found to agree roughly with the projected ranges R p calculation [11, 12]. However, it is

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known that the implantation condition affects the penetration depth and/or range straggling in the case when the extreme experimental parameters such as highly intensive flux are employed.

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Figure 3. Depth profiles of the Cu atomic density in the amorphous SiO2 implanted with 60 keV Cu-, at four dose rates of 1, 3, 10, and 45 μA/cm2 (Kishimoto, et al. [13]). Solid lines are calculation and histogram show experimental distribution.

Kishimote et al. [13] reported that the depth profiles in the amorphous SiO2 implanted with 60 keV Cu- ions deviates considerably from the calculation by TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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TRIM code for the higher flux beams more than a few 10 μA/cm2. The changes in the target morphology against dose rates are shown in Figure 3 [13]. While the Cu depth-profiles agree with the TRIM code under moderate dose rates, the profiles become shallower near the surface and narrower with increasing dose rate. For example, sphere nanoparticles of 10 ~ 15 nm in diameter was observed in amorphous SiO2 matrix implanted by Cu –- ions at 45 μA/cm2 and to a fluence of 3 × 1016 ions/cm2, evolving a specific twodimensional distribution at a specific depth less than the projected range. The dose-rate-initiated process and temperature rise during the implantation are supposed to induce radiation-enhanced diffusion of Cu solutes and promote the Cu precipitation towards the surface. The narrowing of the depth distribution seems to be useful in considering technological application of the nanocomposite. Another result is that the size obtained at the low fluence of 3 × 1016 ions/cm2 is eminently larger than the other cases observed for Fe implantation in Al2O3 matrix at higher fluences (see Table 2). Thus, it is suggested in SiO2 matrix that the implanted layers exposed to subsequent projectiles are inevitability subjected to the competition between in-beam growth and decomposition of the precipitates. On the other hand, the crystallographic structure of matrices does not seem to affect the coarsening or precipitation of implanted irons.

 

Figure 4. An arrangement of CEMS spectrometer, designed to measure spectra under external magnetic field of 0.4 T at maximum.

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(quartz) [7]. It should be rather mentioned that the chemical reactivity between projectiles and matrices plays a role in the precipitation process, as discussed later. 8.E+05

1.E+06

 

1.E+06

(A ) 1.5×1017　Fe/ｃｍ2　in A l2O 3 R oom Tem perature

8.E+05

C ounts

C o u n ts

M easured at 4.2 K

7.E+05

1.E+06 1.E+06 9.E+05

7.E+05 7.E+05 7.E+05

8.E+05 7.E+05 -10

(B ) 1.5×1017　Fe/A l2O 3

7.E+05

-8

-6

-4

-2

0

2

V elocity (m m /s)

4

6

7.E+05 -10 8 10

-8

-6

-4

-2

0

2

4

6

8

10

V elocity

Figure 5. CEMS spectra of Fe-Al2O3 granules prepared by ion implantation with a fluence of 1.5 ×1017 ions/cm2. Spectrum (A) was taken at room temperature and (B) at 4.2K.

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Figure 6. TEM micrographs of Fe-Al2O3 granules implanted to (a) 0.8 × 1017 ionsFe/cm2 and (b) 1.5 × 1017 ions/cm2.

Different measuring techniques have been employed for determining the physical and chemical properties of implanted atoms and the aggregation states of the implanted impurities in refractory metal oxides and ionic materials. A combination of various experimental techniques including TEM, XPS, EXAFS, GXRD (Glancing X-ray diffraction) and RBS (Rutherford backscattering and channeling spectroscopy) allows us to investigate phenomena in the implanted zone in a complementary way. When iron-oxide granules are the object of research and when mass-analyzed 57Fe beam is used

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as implanting ions, conversion electron Mössbauer spectroscopy (CEMS) offers a highly useful method [8].

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Table 2. Conditions of Fe implantation into SiO2 and Al2O3 matrixes and size data of granules obtained from GXRD measurement Sample symbol

Matrix

Fluence at 100 keV (ions/cm2)

Fluence at 74 keV

A B C

SiO2 SiO2 Al2O3

1.0 × 1017

0.4 × 1017 1.5 × 1017

1.5 × 1017

Granules lattice parameter (nm) 0.286 0.284 0.286

Estimated Diameter (nm) 6 15 7

There are two essential features of 57Fe CEMS application important for the present studies. One of them is that CEMS probes only the surface layer of samples, at the depth of about 200 nm, which means a thickness comparable with the range of implanted ions in the most cases. This arises from the fact that the mean range of 7.3 keV conversion electrons from 57Fe nuclei is about 200nm in refractory-metal oxides. Another benefit is that the probing 57Fe atoms themselves compose the precipitations or clusters which we want to characterize. This means that CEMS measurements make it possible to investigate sensitively and effectively the physical and chemical properties of implanted zone. For example, CEMS measurement is so sensitive to investigate samples implanted with fluences as low as 1 × 1015 Fe/cm2, which is difficult to be measured by other conventional techniques. Through the measurements of hyperfine parameters and the spectral line profiles one can illuminate the crystalline configurations of Mössbauer probe 57Fe sites in the implanted insulators. Whereas the valence states and condensed phase of the iron with particle sizes limit can be identified through the CEMS measurement, TEM observations can determine directly precise particle sizes, morphologies, and crystal structures. This suggests a combination of these two techniques to be extremely useful for characterizing implanted layers as the complementary means [14]. In the most studies CEMS spectra are measured in the backscattering geometry, using a 57Co source in a Rh matrix with the activity of 0.3 ~ 3.7 GBq. In the geometry the incident angle of γ-rays is 90o to the sample. Conversion electrons are detected with a He-CH4 gas-flow proportional counter for room temperature measurements. Generally, the spectra were taken

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with a constant acceleration triangular motion. Although CEMS characterizes sensitively and nondestructively the magnetic granules, the available information is still restricted for extremely small particles owing to their magnetic relaxations. One of the effective countermeasures against this problem is to conduct measurements at low temperatures and another is to measure CEMS under applied magnetic fields. With our other CEMS apparatus we can do measurements under applied magnetic field to 1.3 T. Figure 4 shows an original layout of CEMS equipment which was developed using a permanent magnet to study aggregation states of iron atoms in an α– Al2O3 single crystal granular films [15].

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Depth (nm)

Figure 7. Element depth distribution of Fe, Al, O and C in Fe-Al2O3 granules implanted to 1.5 × 1017 Fe/cm2 analyzed by EDX.

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It is demonstrated in section 4 how the CEMS with applied fields is useful to explore magnetic granular systems. CEMS 13measurements below room temperature were performed using an equipment with a channeltron detector which was placed in the inner evacuated tube of a liquid helium cryostat, developed at Duisburg University in Germany [16]. Such an explicit example of CEMS spectra is shown in Figure 5 for Fe nanocomposites in Al2O3 matrix measured at room temperature and also at 4.2 K [17]. The nanocomposite granules were prepared by ion implantation of 57Fe to a fluence of 1.5 × 10 17 ions/cm2 where the peak Fe concentration was calculated to be 3.8 × 1022 atoms/cm3 from TRIM code. The nanoparticle size was estimated to be about 3 ~ 6 nm in diameter from GXRD pattern. Whereas the CEMS in Fig 5 (A) is in contrast with that shown in Figure 11 taken from Fe-SiO2 granules with the almost same Fe fluence, it presents features similar to the CEMS of Fe-SiO2 with lower peak concentration shown in Figure 2; the characteristics of Fe granules in SiO2 are discussed in the following section 31). The weak magnetic splitting lines in Figure 5 (A)manifest that a lot of granules in the Al2O3 matrix are in superparamagnetism. As the granules are apparently smaller in size than those in SiO2 matrix (see Figures 2 and 3), it is manifested that the size of implanted granules depend on the matrixes. On the other hand, Figure 5(B), where the measuring temperature was lowered, clearly shows that superparamagnetic relaxation is now blocked in the most granules at 4.2 K. The details will be discussed in Chapter 4. The measured spectra are analyzed with the computer least-squares fitting with the assumption of Lorentzian shapes to obtain hyperfine parameters. One of the parameters is isomer shift which originates from the s-electron density at the nucleus; e.g., in the case of 57Fe oxidation the change in chemical states leads to the increase in s-electron density. Comparing the values of isomer shift with standard compounds, we can obtain unequivocal information on the 57 Fe valence states. Electric quadrupole and magnetic dipole hyperfine interactions with the nucleus cause splitting of the line energies in Mössbauer spectra. By the quadrupole interaction, in the case of 57Fe, the nuclear excited state is split to produce two lines, i.e., a doublet. The magnetic hyperfine field induces the Zeeman splitting of the nuclear energy levels and produces a sextet pattern (six lines). The relative intensity in a spectrum gives data on the direction between the hyperfine field axes and the direction of γ–ray beam. In observed spectra the line intensity depends on the 57Fe concentration and the recoilless fraction which in turn reflects the binding of the Mössbauer-active atoms with the environments in the compound or lattice.

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TEM is an effective technique to analyze the structure of granular solids since the granules are distributed on nanometer scale. Our TEM observations were done with a Hitachi HF-2000 field emission electron microscope operating at 200 kV, equipped with an EDAX (Energy-dispersive spectrometer) at Kochi University of Technology. TEM micrographs for the Al2O3 samples implanted to 0.8 × 1017 and 1.5 × 1017 ions/cm2 are displayed in Figures 6 (a) and (b) [14]. These TEM samples were prepared by the same conditions as those used in CEMS measurements. Therefore, the physical properties characterized with CEMS and GXRD measurements can be referred to the observed morphology; especially, the evaluated sizes have been proved to be comparable with the size distribution depicted in Figure 15 [23]. The comparison of the images in Figures 6 (a) and (b) reveals that Fe injection leads to not only the concentration increase but the coalescence of the nanoparticles. The growth of the particles larger than 3 nm is evident in the implantation to 1.5 × 1017 ions/cm2. The microstructures are characterized by the four regions extending from the surface to a depth of 150 nm or so. Clustering of the implanted Fe ions is clearly observed in the zone B, while the zone A and C are the layers of amorphous Al2O3 containing Fe ions; the underlying zone D is considered to be radiation-induced polycrystalline Al2O3 matrix. Figure 7 shows the depth distribution of Fe, Al, O and C through the zones by EDAX analyses. The Fe curves represent the implants distribution, and the difference between the two fluences reflects sputtering effect, i.e., retardation of the surface for the higher fluence. Furthermore, it is suggested from Figure 6 (a) that the clustering is induced at the Fe concentration attained with rather lower fluence, i.e., the Fe contents of a few at. %. This observation is not in contradiction with the report by McHargue et al. describing a starting fluence for Fe clustering in Al2O3 matrix to be (2 ~ 3) × 10 16 ions/cm2 [18]. X-ray diffraction is a complimentary technique to clear up the crystalline structure of metal granules. GXRD of implanted composite surface layers was recorded to determine the crystal structure and lattice parameter of the granules, in spite of their small sizes. The GXRD pattern was dated using CuKα radiation at incident angle of 0.5o – 2.0o to specimen surface. The identification of crystalline phases was done using JCPDS database card (Joint Committee on Powder Diffraction Standard, Powder Diffraction File, ASTM, Philadelphia, PA, 1992). Furthermore, the nanoparticles size (structural coherent length) was estimated from the linewidth of X-ray diffraction peak using Scheller’s formula. In GXRD analysis the diffraction peak was fitted assuming Gaussian profile.

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3. ION IMPLANTATION 3.1. Single Ion Implantation The technique of ion implantation is extensively used for microelectronic applications in semiconductors and also for modification of material surface so as to improve hardness, corrosion resistance, etc. in metals. Also it holds applications to develop some interesting prospects for insulators. However, the phenomena accompanying the implantation, especially in insulating solids, are complex because in the solid phase the implanted impurities can be incorporated in various charge states that can associate in turn with various lattice defects, thus creating a variety of metastable impurity-defects complexes.

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Figure 8. XPS depth profile in the 5 × 1016 ions/cm2 implanted silica evidencing the different chemical bonds of dopant atoms (a), and O/Si ratio on the same sample.

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Therefore, it is noted that in spite of the great possibilities offered by metal ion implantation in glassy and ceramic composite materials the physical and chemical mechanisms governing the clustering process are still under debate. The interaction of an energetic ion beam with an insulator differs from that with metals at several points. Insulator compounds generally composed of two or more elements have at least two sublattices, usually in an ordered structure. The most eminent effect induced in substrates is the displacement of the constituting lighter element by the incident beam towards deeper inside. Furthermore, it is unlikely that constituent atoms or ions of one type reside at the lattice sites of another species after the displacement due to elastic collisions, especially in ionic substrates. Oxides such as MgO and Al2O3 are characterized by significant amounts of ionic bonding, with ionic character to lesser extent for Al2O3, and are not allowed to deviate from stoichiometry to considerable extent, because local electrical neutrality must be maintained even in the damaged microstructure. Thus, F-centers are formed by anion vacancies with trapped electrons and V-centers by cation vacancies with holes. Otherwise, stoichiometry may be maintained by the defect complexes such as Schottky pairs, Frenkel pairs, and/or dislocation tangles. The implanted impurity ions are supposed to exist as impurity-point defect complexes, or precipitate as metallic particles/clusters in many cases [18]. In addition to the mass difference of the constituents the production of defects is influenced by the type of chemical bonding during the cascade cool-down period. A number of works including theoretical ones has been done to study a modification of the stoichiometry and of the chemical state in SiO2 or silica substrates induced by ion irradiation; the early implantation study commenced in 1973, intending optical applications [19]. Later developments have expanded from one kind of metal implant to the use of many ion species and to the active formation of compounds, including the inclusions of metal alloys and different composition precipitation. As an example the change in SiO2 composition and the silicide formation in the Mn ion implanted layer are displayed in Figure 8 [19]. The XPS measurements exhibit that oxygen atoms are displaced towards the inner side of the substrate, and that the deviation from stoichiometry in the implanted zone leads to the formation of Mn-O bond and Mn3Si silicides depending on the depth from the surface. The depth profiles of implanted metal ions shown in Figures 7 and 8 give us significant information on the precipitation process of the metals. In Co implantation Co oxides and Co silicides including metallic particles expected to be formed, judging from the values of Gibbs energy for chemical are reactions between Co and silica substrate [20]. However, at the fluence of

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4 × 1016 Co/cm2 XPS measurement detected metallic Co, but Co oxides did not amount to the detection limit nor exhibited the formation of Si-Co chemical bonds. In the experiments the formation of Co nano-clusters with 2.8 nm in diameter was confirmed by TEM and GXRD observations [20]. On the other hand, we found different results for Fe ion implantation into amorphous and crystalline SiO2. Fe ion implantation into amorphous SiO2 matrixes was investigated with changing implanting conditions, which are summarized in Table 2, together with the size of the nanoparticles estimated from GXRD measurement [7]. Figure 9 shows simulation curve of the Fe distribution calculated from the TRIM code. It is noted that the implantation with dual energies of 100 keV and 74 keV offers two third peak concentration of the 74 keV implantation. The curve for sample C presented for comparison was calculated for 100 keV Fe implantation into Al2O3. To obtain the structural information of the implanted layers GXRD patterns were measured as shown in Figure 10. We can see a rather clear peak around 2θ ≃ 44.5o corresponding to the diffraction from bcc α-Fe (110) plane, which indicates the formation of Fe nanoclusters. The GXRD peaks are characterized by broad linewidth typical of the specimens with small grain size. The peaks were analyzed by least-squares fitting assuming Gaussian curves, and the obtained parameters are listed in Table 2. The lattice parameters are in good agreement with the value of bulk α–iron, 0.287 nm. It is obvious from the size estimations that the process of the clustering or precipitation depends on the implanted matrices, which is also suggested by CEMS measurements. Thus, it is demonstrated that GXRD measurement is effective means to study the crystalline structures of implanted granules. The different granule sizes are clearly reflected in CEMS spectra through the effect of superparamagnetic relaxation, as shown in Figures 2 and Figure 11. The two spectra were taken from Fe granules implanted in SiO2 with the almost same total fluence but different peak intensities. As for granule size, the increment in diameter (15 nm) of the sample B in Table 2 whose CEMS is shown in (Figure 11), compared to the 6 nm in diameter; CEMS as shown in sample A (Figure 2), is too much to explain by the increase in Fe concentration depicted in Fig 9. Therefore, the size increase is considered to be brought by the coalesced nanoparticles in SiO2 matrix. Thus, CEMS spectrum of Figure 11 exhibits ferromagnetic (FM) lines arisen from the multi-domain lusters, which grew at the expense of superparamagnetic Fe0 (SPM). It is noted that the appearance of FM lines is related to the longer relaxation time τ due to the increase of granule volume V or the rise of blocking temperature TBi (see the equations (1) and (2) in Chapter 4). On the other hand, CEMS of the

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sample C shown in Figure 5 reveals that the granule size is much smaller in Al2O3 than SiO2 matrix and that TBi for a number of the granules in Al2O3 matrix is lower than liquid He temperature. In the Fe-SiO2 granules implanted to high fluence the FM lines in the Figure 11 spectrum are decomposed into three sets of magnetically split six lines, whose spectrum is very similar to that observed for bulk disordered FeSi alloys [21]. The result suggests that Si atoms are dissolved from the SiO2 matrix into Fe clusters. Referred to the hyperfine parameters depicted in Table 1 of the reference 23, the values of three resolvable hyperfine fields and isomer-shift indicate that there exist three kinds of iron sites with different environment i.e., the nearest-neighbors’ coordination of 8Fe, 7Fe-1Si and 6Fe2Si. Figure 12 shows the variation of hyperfine field Bhf at an iron nucleus with the number of nearest-neighbor Fe atoms. Besides the data from Fe-SiO2 granules, two sets of data from disordered Fe-Al alloys with 10.6 and 15.2 at. % Al are included in the figure [21]. The linear variation shows that the change ΔBhf in Bhf relative to pure α–Fe for any site can be expressed as ΔBhf 6

= ∑ m n Δ B n , where mn is the number of nth nearest-neighbor sites occupied by

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n =1

solute atoms and ΔBn is the contribution to ΔBhf by the single solute atom at the site [22]. This relationship and small increase in isomer shift observed at the same time are caused by a transfer of electrons to the iron 3d-band from the solute atoms, Si. It should be noted that the relationship is useful to make an evidence of alloying itself and characterize the alloy nanoparticles by CEMS. In the case of Fe implantation into SiO2 matrix the granules are presented to be not composed of pure Fe nanoparticles but of Fe-Si intermetallic alloy particles. Table 3. GXRD data and lattice parameters of the Fe-Co/Al2O3 granules prepared by Fe and Co co-implantation to a total fluence of 2 × 1017 ions/cm2 Granule

2θ (degree)

Fe (100%) Fe-Co(25%) Fe-Co(60%)

44.68 44.79 45.16

Lattice parameter (nm) 0.287 0.286 0.284

Particles’ diameter (nm) 7.5 6.2 5.9

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Figure 9. Fe ion distribution calculated from TRIM code for the implantation conditions given in Table 2. While Fe ions were implanted with 150 keV for sample B and C, the implantation for sample A was performed with multi-energy of 100 plus 74 keV.

 

Figure 10. GXRD patterns from implanted layers of the sample A implanted with (100 + 74) keV Fe into SiO2, sample B with 100 keV Fe into SiO2, and sample C with 100 keV Fe into Al2O3.

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77

Relative Intensity

1.48 1.38

1.5×1017 Fe/cm2

C

1.28 1.18 1.08 0.98 -10

-8

-6

-4

-2

0

2

4

6

8

10

Velocity [mm/s]

SPM

 

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Intensity

Fe2＋

FM

Velocity Figure 11. CEMS spectra from Fe implanted SiO2 to a fluence of 1.5 × 1017 ions/cm2 at 150 keV, that corresponds to the sample B in Table 2 and has a higher peak Fe concentration than that in Figure 2 (also, see Figure 7).

The Si concentration in the nanoclusters is supposed to be in the range of 5~10 at. % from the intensity ratio of the three Fe sites in CEMS spectra and from Figure 12. Although the Si incorporation into Fe precipitation was predicted from thermochemical consideration [20], we have succeeded in presenting it experimentally by the analysis of hyperfine parameters, for the first time. The result indicates that there is a case where the implanted granules are “contaminated” by the incorporation of matrix constituents, and such possibility could be guided by the article of Cattaruzza [20]. It should be noted that the alloying or the contamination causes the unexpected or unconscious change in physical properties of the granules and may result in incorrect application of optical and magnetic nanocomposites.

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H yperfine field (T )

 

34 32 30 28 26 24 9

8

7

6

5

4

N earst neighbor Fe atom s

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Figure 12. Variation in internal magnetic field at an iron site in Fe-Si nanoparticles (open circle) with the number of nearest-neighbour Fe atoms, compared with those in Fe-Al disordered alloys (closed squares for 10.6 at. % Al and closed triangles for 15.2 at.% Al).

The CEMS spectra of the granules synthesized by implantation are composed of two or three quadrupole doublets in addition to the Fe0 peak, whose origin can be assigned to ferric and ferrous ions. It is known that the ferrous and/or ferric irons as charge states are observed in most refractory metal oxides at the beginnings of implantation in the range less than 5 × 1016 ions/cm2, and Fe0, precipitation of metallic α-iron, thereafter grows progressively with increasing fluence. Thus, a characterization of the nonequilibrium systems was a subject of considerable interest to understand the evolution process of the nanoparticles. Perez et al. calculated the probability of finding different structural configurations of implanted atoms as a function of the atomic fraction by a statistical model, using a binomial distribution [12]. They reported that in the lower fluence range the calculations is consistent with CEMS observation for Fe ions implanted in MgO substrates, and showed that the implanted Fe ions reside in various charged states, which are closely correlated with local configurations of iron impurities and various lattice defects, leading to impurity-defect complexes. Figure 13 displays an example of CEMS, where 70 keV Fe ions is implanted into MgO at a fluence of 8 × 1016 ions/cm2; the three sets of bar diagram in the figure denote three quadrupole doublets, one from Fe3+ and two from Fe2+ states. Considering the

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neutrality in the implanted oxides, it is quite natural that Fe3+ states exist in the Fe implanted MgO as an isolated impurity at the relatively lower fluences and form complexes with cationic vacancies.

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Figure 13. CEMS spectra for MgO crystals implanted with a fluence of 8 × 1016 ions/cm2 of 100 keV Fe ions, measured at 300 K and 77 K. The two doublets with larger IS are for Fe2+, a doublet with IS of ~0.3 mm/s is for Fe3+, and a single line for metallic Fe0 [12].

The existence of the complexes is considered to be associated with a large displacement and polarization of the oxygen ion situated between the trivalent ion and the vacancy. One of the ferrous irons is supposed for the replaced Fe TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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atoms which reside in cationic sites because a good solubility of Fe2+ in MgO lattice is realized by a small mismatch in size between the Fe2+ and Mg2+ ions, i.e., 0.76 and 0.65 Å in radius, respectively. As a matter of fact, the magnesiowüstite solid solution is achieved in ion implanted MgO, where Fe2+ ions are randomly distributed in the cationic sublattice. The observed magnitude of IS (isomer shift) = 1.0~1.1 mm/s, for one of the Fe2+ doublet is in good agreement with that in magnesio-wüstite [12]. CEMS reveals the existence of another Fe2+ component with different hyperfine parameters. Judging from the large value of QS (quadrupole splitting) ≈ 2 mm/s, another Fe2+ seems to be specific for implanted systems and to be associated with lattice defects. Multiple charge states of implanted iron have been reported also in Al2O3 matrix [18]. CEMS obtained from implanted zone are complex and the spectra have been analyzed by assuming various types of quadrupole doublets including Fe2+, and Fe3+ or high spin states, Fe4+, although the evolution of these ionic states is often dependent on the implantation conditions. Taking account of the hyperfine parameters of the observed CEMS, it is reasonable to consider that the spectra at room temperature can be analyzed with the components of two types of ferrous iron Fe2+I, II metallic α–iron Fe0, and γ–like Fe in Al2O3 matrix with the fluence range higher than 1 × 1017 Fe/cm2 [23]. When we refer to the results from Fe implantation in MgO matrix and also take into account the insolubility of Fe in Al2O3 matrix, we can infer that one of the Fe2+ doublets corresponds to the irons with Fe-O bond in the oxides, Fe1-xO, which may be formed in the inner oxygen rich layers as shown in Figure 5. Judging from the isomer shift (1.1 ~1.5 mm/s), another ferrous doublet could be assigned as Fe2+ state arising from irons in the octahedral sites combined with some defects. With increasing fluences the excess cations become significant, accompanied with the excess oxygen vacancy and lattice disorder, and these promote the aggregation of Fe0. The Fe0 component is supposed to appear with irons aggregating three or more atoms in the nearestneighbor cationic sites, working as nucleation centers for further cluster growth. Under the matrix implanted to high fluences it is reasonable that the nanocompsite system is stabilized by aggregating the excess metal impurities and with increasing fluence the clusters grow via a mechanism such as Ostwald ripening. Above percolation limit, i.e., fluence more than 1.5 × 1017 Fe/cm2, coalesced clusters are supposed to be formed in the granules.

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The study on Fe implantation into Al2O3 to higher fluences has been a subject of interest to synthesize magnetic granules [6, 8, 23]. After the first observation of tunneling magnetoresistance (TMR) effect in Fe-implanted Al2O3 granules was reported by Sakamoto et al. [6], Wakabayashi et al. studied the Fe-Al2O3 granules by measuring CEMS under applied magnetic field [23], intending to obtain detailed information on the aggregation states of implanted Fe ions and optimize the granular characteristics to achieve higher GMR effect. Figure 14 shows a series of CEMS from the samples prepared by the implantation of 100 keV 57Fe+ ions at a fluence of 1.5 × 1016 ions/cm2 into the R-plane of a single crystalline α–Al2O3 substrate and measured at room temperature with applied fields of Ba = 0, 0.1, 0.2, 0.3, and 0.4 T. With increasing applied field, the broad magnetic split lines around the tail of the peak at zero-velocity evolve into a well-resolved sextet with I.S.≈ 0 mm/s, while the central peak diminishes relatively. The spectra analysis was done by assuming five components, where superparamagnetic Fe0 line was distinguished from that of ferromagnetic Fe0 precipitation so as to realize clearly the effect of applied fields. The hyperfine field Bhf for the superparamagnetic Fe0 component observed under applied field Ba at a temperature T is given by Mørup et al. [24], B hf = B0 L( μBa/kT) – Ba

(1)

where L(μBa/kT) is Langevin function, B0 is the saturation hyperfine field, μ is a magnetic moment of the particle and k is the Boltzman constant. After deriving the distribution of hyperfine field by the analysis of the spectra, we have succeeded in obtaining the particle size distribution through equation (1). The result is shown in Figure 15. The estimated sizes are in good agreement with the TEM observations shown in Figure 6, and the bimodal distribution is in consistence with other works on implantation in oxide matrixes [12]. From the measured distribution the average diameter was obtained as about 3.5 nm for the smaller superparamagnetic particles and the larger particles was about 6.2 nm in diameter. The latter value is close to the estimation by GXRD.

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Figure 14. CEMS spectra (open circle) of Fe into Al2O3 granules implanted with a fluence of 1.5 × 1017 ions/cm2 observed under applied fields of 0, 0.1, 0.2, 0.3 and 0.4 T Fits to the spectra (thin lines) are also shown with the contributions from ; two kinds of -Fe particles (screened), Fe like clusters (hatched), and two kinds of ferrous oxides (thick lines) [23]. .

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Figure 15. Particle-size distribution for metallic Fe0 where α–FeI corresponds to superparamagnetic components (Fe0) and α–FeII to ferromagnetic Fe0, respectively, obtained from the analysis of CEMS in Figure 10.

 

Figure 16. Changes of Bhf for ferromagnetic Fe0, Fe0 (f), and relative intensity for the main components of Fe0(f), superparamagnetic Fe0(s), and gamma like Fe0, Fe(γ), plotted against Fe ion fluence. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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CEMS measurements make it possible to estimate not only the size but the shape of nanoparticles. The intensity ratio of the magnetically split lines in Figure 14 reflects the direction of the particles’ magnetization relative to the incident γ–ray direction. The direction is determined by the condition of minimizing free energy, F. It is given by the equation of F = KV sin2 (π /2 – θ) –Ba MsV cos θ, where Ms is the saturation magnetization of α–Fe, V is the volume of spherical nanoparticles, and the angle θ is measured from γ ray direction. The external field Ba is applied perpendicular to the sample surface, i.e., perpendicular to the easy axis direction and parallel to γ ray direction. The preferred direction θ0 corresponding to ∂F/∂θ = 0 is given by, θ0 = π − sin−1 ⎛⎜ BaMs ⎞⎟ . 2 ⎝ 2K ⎠

(2)

Therefore, when the line intensity ratio for six lines is assumed as 3 : X : 1 : 1: X : 3, the preferred direction of magnetization is determined by,

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θ0 = sin−1

2X 4+ X

(3)

Using equations (2) and (3), the average anisotropy constant has been estimated to be 2.3 × 105 J/m3 from the observed dependence of the parameter X on applied fields. The value is coincident with the anisotropy constant of the Fe-SiO2 granules estimated previously. The difference from the bulk anisotropy energy constant could be attributed to the shape anisotropy. When the larger nanoparticles are assumed to be oblate ellipsoids, the ratio of semimajor axis to semi-minor axis was calculated as 1.35 [23]. The estimated size and particles’ shape are in agreement with the TEM observations [18, 24]. Thus, it is demonstrated that CEMS, as a nondestructive means, is highly useful in estimating the size distribution of the nanoparticles embedded in oxide matrixes. In the later section we will discuss again how useful guide such analysis offers in improving TMR characteristics. We have studied the magnetic properties of the Fe-Al2O3 granules implanted up to such high fluences as 3.5 × 1017 ions/cm2 [8]. An example of the CEMS spectra at a fluence of 1.5 × 1017 ions/cm2 is given in Figure 5(A). The split lines exhibit broad peak width and the apparent average hyperfine field is smaller than the bulk value of 33 T. This suggests that the magnetization vectors in the ferromagnetic particles are fluctuating in

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directions close to the easy axis of magnetization, i.e., collective magnetic excitations [25]. With increasing the fluence the ferromagnetic (FM) component of α–Fe grows eminently at the expense of the superparamagnetic (SPM) α– Fe. The relative intensity of the SPM peak diminishes to a few percent of the spectral components and most Fe nanoparticles are in multidomain clusters by the implantation to 3.5 × 1017 ions/cm2. Figure 16 shows the variation of hyperfine field and the relative intensity of the main spectral components plotted against Fe fluence; the cited components are denoted as Fe0(s) and Fe0(f) for SPM and FM lines, respectively. It is noteworthy that the Fe0(f) component and its hyperfine field intensity in the narrow fluence range of 1.5 ~ 2.0 × 1017 ions/cm2, and the range is considered to be a percolation limit of α -iron nanoparticles in Al2O3 matrix. Percolation operates more effectively when particle densities are higher. The nanoparticles size formed by implantation at 2.0 × 1017 ions/cm2 under the conditions similar to us is observed to have the distribution with diameters was of 5 - 20 nm by TEM [24]. Thus, the dependence of the physical properties on the fluence will give us useful guide to discuss TMR effects in nanocomposites formed by implantation.

3.2. Multiple Ion Implantation

Intensive effort devoted to the research of synthesis and properties of nanocomposites is due to the fact that nanocomposites exhibit novel and promising properties for a variety of application as mentioned before. The properties not only depend on the crystalline structure, but can also be tuned by their composition. Metallic alloy nanoparticles, composed of two or three species of metal, are more complicated and have more extensive applications in the optical and magnetic research field compared with single-component metal. Ion implantation technique allows an easy control of the nanostructure in both size and dispersion by properly tuning the implantation conditions. Moreover, the composition of binary nanoparticles can be varied by sequentially implanting two species. Since the important physical properties of nanocomposites can be optimized for a particular application by controlling the concentration and size, implantation techniques can be used to create novel materials by combining with many other experimental techniques including post-implantation treatments [26]. Such typical implanting procedures have been developed to modify surface plasmon resonance absorption and enhance the nonlinear optical properties. Various types of binary alloy nanocomposites

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have been reported by implantation with two ion species, including the combination of Ag-Sb, In-Ag, In-Cu, Ag-Cd, Cu-Ni and Au-Cu [27, 28, 29]. It is reported that binary nanoparticles can be present as alloys but also as coreshell structure. The latter structure may import new properties arising from the interaction or combination of the properties of the two pieces, offering another possibility to produce novel metallic precipitates [30]. Au and Cu have complete miscibility at any concentration in the bulk phase. While single Au or Cu implantation in silica leads to the nanoparticle formation of the respective metals, the formation of metallic nanoparticles with an average diameter of 3.8 nm at the total fluence of 6 × 1016 ions/cm2 was also found in the sequential implantation of Au and Cu by TEM observation, suggesting the formation of AuxCu1-x intermetallic alloy [29]. Because x ~ 0.68 was obtained from Vegard’s law using the observed lattice constant of the nanoparticles, the Au/Cu atomic ratio was supposed to be about 2.1 or alloy of composition around Au70Cu30. However, it was found that subsequent thermal treatments modify the aggregation characteristics such as the alloy composition, particle size, Au precipitation, and so on [31]. Cu and Ni also have complete miscibility at any concentration in the bulk phases. TEM micrographs on the Cu and Ni co-implanted layer with equal fluence of 6 × 1016 ions/cm2 into soda glass and silica proved the existence of nanoparticles of fcc phase with a lattice parameter value slightly smaller than that of bulk Cu [20]. The result suggested that Cu-Ni alloy nanocluster of ~6 nm in diameter was formed with equal composition of about 50:50 as expected from Vegard’s law, but the proof of the formation of alloy nanoparticles seems not to be enough. Above all, what is important is how one can obtain evidences for alloying in the nanoparticles. We have shown the alloying process can be observed by combined means of GXRD and CEMS [32, 36]. To synthesize magnetic alloy granules the co-implantation of Fe and Co ions was done into Al2O3 matrix and the clustering process was investigated mainly by Mössbauer spectroscopy [32]. Cobalt is known to form a continuous range of solid solutions with iron, whose structures at room temperature are bcc for 0 - 73 at. % Co and fcc for 73 - 92 % Co in bulk alloys [33]. In the experiments sequential implantation of 57Fe ions followed by 59Co ions were performed at room temperature to two total fluences of 1.5 and 2.0 × 1017 ions/cm2 with projectiles’ energy of 100 keV. The proportion of the fluence of two ions was changed to obtain alloy nanoparticles with various Fe-Co compositions. Firstly the formation of alloy composites was confirmed by GXRD measurements. The X-ray diffraction data from α– Fe (110) plane are summarized in Table 3, where granules diameter estimated from Scheller’s

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formula is also listed. While the lattice parameter of the granules for single Fe implantation (Fe 100 %) agrees with bulk α– iron, the decrease in lattice parameters with increasing Co concentration is consistent with the corresponding change in bulk Fe-Co alloys, when referred to 1977 JCPDSInternational Center for Diffraction Data. The results indicate that Fe-based alloy nanoparticles are precipitated in the implanted layers. Further evidence for the alloying has been obtained from CEMS measurements. It was observed that CEMS spectra taken from the samples co-implanted with Fe and Co ions exhibit patterns similar to the spectra of single Fe-implanted samples and also the similar change in the spectra when the total fluence increased from 1.5 × 1017 to 2.0 × 1017 ions/cm2. However, the values of hyperfine magnetic field Bhf and isomer shift of both FM and SPM lines are in good agreement with bulk Fe-Co alloys [32]. The magnetic FM lines grow in exchange for the SPM singlet line with increasing fluence, because superparamagnetic relaxation in the “alloy” granules is blocked due to the volume increment. Another eminent event is the increase of hyperfine magnetic field Bhf observed in CEMS of the Fe-Co nanocomposites. Figure 17 shows the comparison of Bhf distribution curves obtained by the analysis of CEMS spectra from the nanoparticles with three Co concentration of 0, 25 and 60 at. %. We do not notice any essential differences in the shapes of these distribution curves, i.e., they have similar FWHM. Although the Bhf distribution exhibits some broadening, the average value is well defined for any Co content. It is remarkable that the Bhf in Fe-25%Co denotes the distribution with the highest hyperfine field. Figure 18 shows the variation of the averaged Bhf value as a function of Co concentration, together with the relative intensity ratio of FM to SPM line. The maximum value of Bhf is achieved at Co concentrations near 25 at. %. Many studies have been published on the concentration dependence of the average hyperfine field at iron sites in bulk Fe-Co alloys [34], and the maximum Bhf value has been reported to be 36.5 T at the Co concentration around 30 at. %, whereas the average Bhf in the nanoclusters with 25 at. % Co is observed to be 34.2 T in Figure 17. The reduction of the average value in the Fe-Co granules can be explained by the influence of collective magnetic excitations [25], whose effect comes out in the wide Bhf distribution. The difference of the concentration dependence between the granules and bulk alloys is likely to be partly caused by the size effect that is depicted by the plots of the relative intensity of superparamagnetic and ferromagnetic peaks and indicates in the figure that the granules become smaller over the 25 at.% Co. Nevertheless, it is significant that the variation of Bhf with Co concentration in the nanoclusters

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exhibits characteristics similar to the well known Slater-Pauling curve, that is, the variation of saturation magnetic moment for the bulk alloys of the first transition group metals, plotted as a function of electron concentration (for example, see Figure 11.7 of Ref. 22).

 

0.14 0.12

25% C o at.% 60% C o at.%

P (B hf)

0.1

Fe

0.08 0.06 0.04 0.02 0 15

20

25

30 hf

35

40

(T )

Figure 17. Distribution of Bhf in Fe-Co nanoparticles with Co concentration of 0, 25, and 60 at.% at the total fluences of 1.5 × 1017 ions/cm2, obtained from analysis of CEMS spectra. B hf D IS T R IB U T IO N

36

100

B hf Ferrom agnetic lines S uperparam agnetic line

34

80

60 32 40 30

Relative Intensity (%)

 

InternalField (T)

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B

20

0

28 0

10

20

30 40 C o C oncentration

50

60

70

Figure 18. Concentration dependence of the averaged hyperfine field Bhf at iron sites (solid curve), and relative intensity of ferromagnetic and superparamagnetic Fe0 line, plotted against Co concentration. The total fluence of Fe and Co implantation was 2.0 × 1017 ions/cm2.

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Intensity (a.u.)

A

B

1.0×1017 Fe/A l2O 3

(1.0+1.0)×1017 Fe+C u/A l2O 3

C D

40

42

44

2θ

1.2×1017 Fe/A l2O 3

(1.2+0.8)×1017 Fe+C u/A l2O 3

46

48

50

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Figure 19. GXRD pattern of Fe-Al2O3 implanted to 1.0 × 1017 ions/cm2 (curve A) and its change after additional Cu implantation to 1.0 × 1017 ions/cm2 (curve B). Curve C and D are GXRD patterns of Fe-Al2O3 implanted to 1.2 × 1017 ions/cm2 and after Cu implantation to 0.8 × 1017 ions/cm2, respectively.

The empirical relation between Bhf and atomic moment is given by the expression [34], Bhf = a · μ(Fe) + b · μ’

 

Figure 20. Plots of magnetic ordering temperatures Tc against Fe composition x in sputtering films of FexCu1-x alloys [38]. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

(6)

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where μ(Fe) is the magnetic moment of Fe atom, μ’ is the average magnetic moment of the alloy, and a and b are proportionality constant. Using the constants empirically determined by Vincze et al. [34], i.e., a = 7.0 T/μB and b = (8.0 – 3.3 cCo) T/μB, we obtain μ(Fe) = 2.4 μB for the Bhf value of 34.2 T observed at 25 at. % Co. It should be noted that the equation (6) gives μ(Fe) = 2.6 μB for bulk Fe-Co alloys at 30 at.% Co when we use the Bhf value of 36.5 T, and thus the μ(Fe) value obtained for the nanoparticles apparently is underestimated because of collective magnetic excitation. Hamdeh et al. [35] evaluated the magnetic moment at Fe atoms by assuming Co magnetic moment to be constant value of 1.85 μB in Fe-Co alloys and counting changes in the number of 3d electrons in Fe atoms alloyed with Co. The Fe magnetic moment was found to increase linearly from 2.22 μB for pure Fe to 2.75 μB at 30 at. % Co, and finally 3.0 μB at 70 at. % Co. We notice that, whereas the maximum mean magnetic moment as well as maximum Bhf is observed at 30 at. % Co, a transition from weak to strong ferromagnetism is supposed at the Co concentration [35]. The above results on the Bhf distribution and its variation with Co concentration in the alloy granules indicate that the sequentially implanted Co ions are uniformly distributed within the previously precipitated Fe clusters; the intimate interactions between the Fe and Co atoms are comparable to that in bulk Fe-Co alloys. It is added that the peak value in Bhf distribution is 35.5 T, which is close to the maximum Bhf of 36.5 T observed in bulk Fe-Co alloys, and that the Bhf distribution in the granules extends to over 37.5 T. It is worth noting that the increment of atomic magnetic moment is expected to enhance TMR effect in the Fe-Co granules as discussed later. Above all, these studies by means of CEMS have verified that the sequential implantation of two element ions is effective to synthesize homogeneous alloy nanocomposites and that almost all of the implanted Fe and Co ions are incorporated in the alloying particles as appeared from the dependence of Bhf on the relative Co/Fe fluence in Figure 18. Since the injection of energetic ions induce the non-equilibrium process in the matrix, it is expected that metastable alloy nanoparticles are synthesized by co-implantation with immiscible elements. Fe-Cu solid solution has been prepared by means of non-equilibrium production techniques, such as mechanical alloying, sputtering co-deposition and melt-spinning rapid quenching. We studied the formation of Fe-Cu alloy nanoparticles by sequential implantation of two immiscible Fe and Cu ions into Al2O3 matrix [36]. The implantation was performed at room temperature to a total fluence of 2.0 × 1017 (Fe+Cu) /cm2. The change in crystalline structure due to the implantation of Cu following Fe ions was monitored by GXRD measurement

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and shown in Figure 19. The diffraction peaks are clearly enhanced to a considerable extent after implantation of the subsequent Cu ions. The positions of peaked diffraction corresponding to the bcc (110) and fcc (111) planes were observed at 2θ = 44.5° and 43.1°, and the lattice parameters of the both peaks were obtained as 0.288 nm and 0.363 nm, respectively, which are close to the values of 0.287 nm in bulk α-iron and 0.365 nm in bulk copper. The slight expansion in the bcc lattice with Cu alloying is reasonable when the larger radius of Cu atom is taken into account. It is remarkable in the figure that both fcc and bcc peaks coexist in the granular Fe60Cu40 sample (curve D) while only fcc peaks are observed in the Fe50Cu50 sample (curve B). The result is in consistent with that reported by Chien et al. [37]. They investigated magnetic properties of FexCu100-x solid solution, using the alloy films synthesized by dc sputtering in magnetic field. Figure 20 presents the plot of magnetic ordering temperature Tc versus Fe at. %, x [37]. In the Fe70Cu30 sample the Tc data clearly display two distinct values, i.e., the composition 60 ≤ x ≤ 75 is in the range of structural transformation from bcc to fcc structures. The behavior of Tc and the internal field in fcc phase is thought to exhibit characteristics of simple magnetic dilution. The sizes of Fe-Cu alloy nano-particles in the implanted granules were estimated to be about 5 nm in diameter from the width of GXRD peaks. Then, CEMS spectroscopy has been used to obtain quantitative information on the local magnetic properties around Fe atoms in metastable alloy-particles, that is, on both the existing Fe phase distribution and the composition and size of the magnetic particles. Figure 21 (a) shows CEMS spectrum from the Fe implanted Al2O3 samples at a fluence of 1.6 × 1017 ions/cm2 and Figure 21(b) is the CEMS after subsequent implantation of Cu ions to 0.5 × 1017 ions/cm2 intending to form an alloy with Cu concentrations of 24 at.%. To be remarkable, the sequential Cu implantation resulted in the increase of ferromagnetic FM lines relative to superparamagnetic Fe0 line, while the internal magnetic fields Bhf of FM component decrease after Cu implantation. Table 4 summarizes some hyperfine parameters of the spectra in Figure 21. The intensity increase in the FM lines could be simply due to the blocking of superparamagnetic relaxation caused by the increase in KV term in equation (1) or (2) , but the increase by only ~10 % in relative FM amount (40 → 50 %) with increasing the total fluence is rather gentle because the alloying is accompanied with magnetic dilution、in contrast with pure Fe implantation shown in Figure 14. Moreover, the decrease in Bhf values is caused by the dissolution of implanted Cu atoms in the iron clusters, which is consistent with the observation by Chien et al. [37].

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1.4E+05 17  Fe/cm22   C ) Fe1.5E17/cm .((A) 1.6 ×10

C ounts

1.3E+05

1.2E+05

1.1E+05

1.0E+05 -10 -8

-6

-4

-2

0

2

4

6

8

10

6

8

10

V elocity (m m /s) 1.6E+05

(B)  (D ) Fe-24%C u

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1.4E+05

1.3E+05

1.1E+05

-10 -8 -6 -4 -2

0

2

4

V elocity (m m /s) Figure 21. CEMS of the granular sample implanted with Fe to a fluence of 1.6×1017 ions/cm2 (a), and subsequent Cu implantation to 0.5×1017 ions/cm2 after the Fe implantation (b).

The CEMS spectra of Fe-Cu alloys are also characterized by broader linewidth, contrasted with Fe-Co and Fe-Si alloy granules. Although the broadening or wider Bhf distribution is commonly thought to come from multiple inequivalent Fe sites, it is also likely that the size of the formed alloy clusters widely distribute or the mixing of Cu atoms into Fe clusters may not be homogeneous in the disordered Fe-Cu alloys, as suggested by the coexistence of bcc and fcc structures; above all, by collective magnetic

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excitation the dispersion of particle volume leads to a Bhf distribution through the equation, Bhf = B0(1 – kT /2K’eff V), where B0 is the saturation of hyperfine field. From the measurements of GXRD and CEMS it is quite safe to conclude that metastable Fe-Cu alloy nanoparticles are produced by the sequential implantation of Fe and Cu ions. The precipitation of metastable alloy nanoparticles suggests that the clustering of implanted metal atoms occurs via the process of collision cascades and cooling down of thermal spike [22]. However, we can not guarantee in the case that all the implanted metal species are involved in Fe and Cu alloying. And also it is suggested that the clusters with core-shell structure were not formed in the Fe-Cu co-implanted layer, because the observed values of 57Fe hyperfine parameters can not be attributed to a interaction with core-shell boundary [30] but to homogeneous solid solution.

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3.3. Modification of Implanted Granules by High Energy Ion Irradiation Ｗhen one of the injecting ions has much higher energy than other ion, a quite different aspect is evolved in the implanted layers [4], and it offers a new technique for tailoring of the implanted granular layers. The control of the size itself and its distribution is important to realize high performance in nanocomposite devices, since the broad spatial and size distribution of nanoparticles are inevitable in ion-beam material synthesis. For example, the eminent size effect can be seen in size dependence on specific magnetic properties or plasmon frequency for non-linear optical application. The control can be done by varying the implantation parameters such as ion energy, flux and sample temperature. Recently, the modification or tailoring of the size distribution was found to be achieved by post-implantation irradiation, where high-energy ions were irradiated through a layer of the granular films [38, 39]. Rizza et al., studied irradiation effect of ~ 4 MeV Au ions to Au nanoparticles embedded in a dielectric matrix [40] and a Au layer sandwiched between two silica films, i.e., SiOx/Au(7 nm)/SiOx [41]. They showed the possibility of obtaining nearly monodispersed distributions by the irradiation of high energy ions. Using Au nanoparticles chemically-synthesized in amorphous SiO2, they observed a refinement in nanocomposite morphology from the original size of 15.9 nm (centroid of the distribution) to a mean size of 2.0 nm after irradiation up to 8 × 1016 ions/cm2. In Figure 22 the change in TEM micrographs and corresponding size distributions with increasing 4 MeV Au fluences are shown

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[39]. The evolution of the smaller nanoclusters as function of the irradiation fluence is presented in Figures 22 (b) and (e). The size distribution of (f) - (j) histograms clearly manifests that with increasing the fluence the continuous dissolution of the original cluster finally results in the very narrow distribution of the small nanoparticles. Figures 23 show another TEM observation, where (a) and (b) exhibit TEM micrographs of as-prepared sandwiched films and the films after 4.5 MeV Au irradiation to a fluence of 1.5 × 1016 ions/cm2, respectively [40]. The island structure in a percolative Au film was modified to the precipitated states of spherical clusters, i.e., the formation of a halo of 5 nm-size clusters surrounding larger inclusions. From these observations they suggest that the radiation induced inverse Ostwald ripening is working through the thin Au layers in SiO2, where the ion beam induces two competing processes of (1) dissolving the nanoclusters by collisional mixing of Au into SiO2 matrix and (2) precipitating as a number of smaller nanoparticles without coalescing to larger inclusions. Although the ripening depends on the irradiation conditions (projectiles’ energy, flux, etc.), the precipitates redistribution by specified condition is expected to lead to a unimodal distribution of the particles size, important for nanocomposite applications [40]. D`Orleans et al. also used high energy ions irradiation to modify the morphology in nanocomposite systems where spherical nanoparrticles of 10 nm in mean diameter in SiO2 matrix were formed in advance by 160 keV implantation of Co ions at a fluence of 1017 ions/cm2 [39]. TEM observations after 200 MeV 127 ions irradiation at fluence range of 1011 – 1014 ions/cm2 indicated the spherical growth in size below a fluence of 1012 ions/cm2, and for higher fluences a growth of elongated particles with the dimension of 9 × 35 nm. Magnetization measurements supported the observed modification of the size and shape, manifesting the magnetic anisotropy that the easy axis is normal to the sample surface or parallel to the irradiation beam direction. The results are in contrast with those of Rizza et al. [40]. The ion flux in this experiment (109 ions/cm2/s) was lower by three orders than that by Rizza et al. and also the fluence range to realize ion-beam modification is apparently different. It is likely that radiation enhanced diffusion may influence the recoalescing of the mixed out solute atoms in the irradiated zone. Although we still need further irradiation experiments and data accumulation to obtain useful information and guides for nanocomposites tailoring, the technique is expected to make possible a precise control of the modification, superior to thermal treatments. The advance in TMR characteristics made by such postimplantation treatments will be described in the next chapter.

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Figure 22. (a)~(d), a sires of bright field TEM micrographs in the course of 4 MeV Au irradiation of the fluence from zero (as-prepared) to 8 ×1016 ions/cm2 at 300 K. The size distributions of the nanoclusters and satellites corresponding to (a) ~ (d) are displayed in (f) ~ (i) [41].

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Table 4. Relative intensity and internal magnetic fields of α –iron component in CEMS spectra from the samples as Fe implanted to a fluence of 1.6 × 1017 ions/cm2 and subsequent Cu implantation to a total fluence of 2.1 × 1017 ions/cm2 As Fe implanted Relative amount of superparamagnetic Fe0 Relative amount of ferromagnetic Fe0 Hyperfine field

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22 %

Fe and Cu coimplantation 15 %

40 %

50 %

30.8 T

24.4 T

(a)

(b) 

Figure 23. TEM micrographs of as-prepared SiOx/Au/SiO2 sandwich film showing percolative structure (a), and of the film after 4.5 MeV Au irradiation to a fluence of 1.5 × 1016 ions/cm2, showing two different types of precipitate [40].

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CONCLUSION

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It has been presented that ion implantation is an excellent technique to fabricate thin granular layers because the technique can be practiced with almost any metal species and performed to allow the direct designing of the structural geometry, which could be advantages in device applications. 1) Magnetic granular layers were formed in the near surface regions of insulating oxides such as MgO, SiO2 and Al2O3 by Fe ion implantation with energy of 100 keV. The projected range was calculated as 50-60 nm. Measurements by means of GXRD, TEM and CEMS were found to be useful in identifying the crystalline and magnetic properties of the granules, giving the size estimation in a non-destructive way . 2) Alloy nanoparticles were successfully synthesized in Al2O3 matrix at as-implanted condition by sequential implantation of Fe and then Co or Cu ions. The magnetic characteristics of the co-implanted Fe-Co granules are very similar to the corresponding bulk alloys, as exhibited by the dependence of internal field on constituent concentration showing likeness to Slater-Pauling magnetization curve. Moreover, metastable FeCu nanoparticles were also synthesized with co-implantation of the two immiscible metal species. 3) In most case the granule sizes are inevitably distributed over a wide range of 1 ~ 10 nm in diameter, depending on implantation fluence; the peak diameter was 2-3 nm for a Fe fluence of 1.1 × 1017 ions/cm2 in Al2O3. Thus, it is important to modify granular structures in order to reconstruct the size distribution and achieve high performance by realizing unique and desired size. While thermal annealing leads to precipitates coalescence and the formation of larger granules, irradiation of high energy heavy ions shows a possibility of tailoring granular layers to the narrower distribution with smaller size than original ones. 4) The measurement of conversion electron Mössbauer spectroscopy has been proved to be useful and excellent means to characterize the physical properties of Fe and Fe-based alloy nanoparticles. The spectra offer effective information on magnetic ordering or superparamagnetism in granular systems. Above all, it should be marked that we can derive the size distribution of the particles by the measuments under external magnetic field or at low temperatures. The

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N. Hayashi, I. Sakamoto, and T. Toriyama estimated sizes are in good agreement with other TEM and GXRD observations.

REFERENCES [1] [2] [3] [4] [5] [6]

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[7] [8] [9] [10] [11] [12] [13] [14] [15]

[16] [17]

P. Mazzoldi, G.W. Arnold, G. Battaglin, F. Gonella, R.F. Haglund Jr., and J. Nolin. Opt. Phys. Mater. 5 (1996) 285. C.L. Chien, J. Appl. Phys. 69 (1991) 5267. D.N. Lambeth, E.M.T. Velu, G.H.Bekkesis, L.L. Lee, and D.E> Laughlin, J. Appl. Phys. 79 (1996) 4496. K.H. Heinig, T.Muller, B. Schmidt, M. Strobel, and W. Möller, Appl. Phys. A77 (2003) 17. B.D. Sawicka and J.A. Sawicki, Topics in Current Physics, ed. By U. Gonser (Springer, Berlin, 1981), vol. 25, pp.139. I. Sakamoto, S. Honda, H.Tanoue, N. Hayashi, and Y. Yamane, Nucl. Instrum. Methods B, 148 (1999) 1039. T. Moriwaki, N. Hayashi, I. Sakamoto, H. Tanoue, T. Toriyama, and H. Wakabayashi, Trans. Mater. Research Soc. Japan, 29 (2004) 607. N. Hayashi, I. Sakamoto,H. Tanoue, H. Wakabayashi, and T. Toriyama, Hyperfine Interact., 141/142 (2002) 163. F. Gonella, Nucl. Instrum. Methods B, 166/167 (2000) 831. J.F. Ziegler, Handbook of Ion Implantation Technology (North-Holland, Amsterdam) 1992. N. Hayashi, T. Toriyama, H. Wakabayashi,,I. Sakamoto, T. Okada, and K. Kuriyama, Surf. Coat. Technol., 158/159 (2002) 193. A. Perez, C. Marest, B.D. Sawicka, J.A. Sawicki, and T. Tyliszczak, Phys. Rev. B 28 (1983) 1227. N. Kishimoto, N. Umeda, Y. Takeda, V.T. Critsyna, T.J. Renk, and M.O. Thompson, Vacuum, 58 (2000) 60. T. Moriwaki, Phd Thesis, Kochi University of Technology, 2006. T. Hirai, H. Wakabayashi, T. Toriyama, I. Sakamoto, and N. Hayashi, Proc. 18th Symp. Mater. Sci. Eng. Res. Cent. Ion Beam Tech., HoseiUniv., 1999, p.159. Y.C. Wang, C.P.Luo, W. Kleemann, B. Scholz, R.A. Brand, and W. Keune, J. Appl. Phys., 73 (1993) 6907. N. Hayashi, F. Stromberg, W. Keune, T. Toriyama, M. Yamashiro, and I. Sakamoto, J. Appl. Phys. 101 (2007) 104304.

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Formation of Granular Layers by Ion Implantation [18] [19] [20] [21] [22] [23] [24] [25]

[26]

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[27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

99

C.J. McHargue, P.S. Sklad, C.W. White, G.C. Farlow, A. Perez, and G.Marest; J.Mater. Res., 6 (1991) 2146. J. Davenas, A. Perez, P. Thevenard and C.H.S. Dupuy, Phys. Stat. Sol. A19 (1973) 679. E. Cattaruzza, Nucl. Instrum. Methods B, 169 (2000) 141. M.B. Stearns; Phys. Rev., 129 (1963) 1136. N.N. Greenwood and T.C. Gibb, Mössbauer Spectroscopy, Chapman and Hall, London, 1971, p.304. H. Wakabayashi, T.Hirai, T. Toriyama, N. Hayashi, and I. Sakamoto, phys. stat. sol. (a) 189 (2002) 515. E. Alves, C. McHargue, R.C. Silvia, C. Jesus, O. Conde, M.F. da Silvia, and J.C. Soares, Surf. Coat. Technol., 128/129 (2000) 434. S. Mørup, J.A. Dumesic, and H. Topsøe, Application of Mössbauer Spectroscopy, vol.11, ed. R.L. Cohen, Academic Press, New York 1980, p. 1. A. Meldrum, L.A. Boatner, and C.W. White, Nucl. Instrum. Methods B, 178 (2001) 7. T.S. Anderson, R.H. Magruder III, R.A. Zuhr, and J.E. Wittig, J. Electron. Mater. 25 (1996) 27. M. Falconieri, Appl. Phys. Lett., 73 (1998) 288. F.Gonella, G. Mattei, P. Mazzoldi, C. Sada, G. Battaglin, and E. Cattaruzza, Appl. Phys. Lett. 75 (199) 55. E. Cattaruzza, Nucl. Instrum. Methods B, 148 (1999) 10079 R. Bertoncello, F. Trivillin, E. Cattaruzza, P. Mazzoldi, G.W. Arnold, G. Battaglin, and M. Catalano, J. Appl. Phys. 77 (1955) 1294. N. Hayashi, I. Sakamoto, H. Wakabayashi, T. Toriyama, and S. Honda, J. Appl. Phys. 94 (2003) 2597. C. E. Johnson, M. S. Ridout, and T. E. Granshaw, Proc. Phys.Soc. 81 (1963) 1079. I. Vincze, I.A. Campbell, and A.J. Meyer, Solid Sate Commun. 15 (1974) 1495. H.H. Hamdeh, B. Fultz, and D.H. Pearson, Phys. Rev., B 39 (1989) 11233. N. Hayashi, T. Moriwaki, M. Taniwaki, I. Sakamoto, H. Tanoue, T. Toriyama, and H. Wakabayashi, Thin Solid Films, 505 (2006) 152. C.L. Chien, S.H. Lior, D. Kofalt, Wu. YU, T. Egami, and T.R. McGuire, Phys. Rev., 33 (1986) 3247. M. Strobel, K.H. Heinig, W. Moeller, A. Meldrum, D.S. Zhou, and C.W. White, Nucl. Instrum. Methods B, 147 (1999) 343.

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[40]

C.D’Orleans, J.P. Stoquert, C. Estourmes, C. Cerruti, J.J. Grob, J.L. Guille, F. Haas, D. Muller, and M. Richard-Ploulet, Phys. Rev. B 67 (2003) 220101. G.C. Rizza, H. Cheverry, T. Gacoin, A.Lamassen, and S. Henry, J. Appl. Phys., 101 (2007) 014321. G.C. Rizza, M. Strobel, K.H. Heinig, and H. Bernas, Nucl. Instrum. Methods B, 178 (2001) 78.

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[41]

N. Hayashi, I. Sakamoto, and T. Toriyama

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Chapter 4

TMR EFFECT IN GRANULAR LAYERS BY ION IMPLANTATION N. Hayashi1, T. Toriyama2 and I. Sakamoto3 1

Kurume institue of Technology, Fukuoka, Japan Musashi Institute of Technology, Tokyo, Japan 3 Hosei University, Tokyo, Japan

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2

1. INTRODUCTION The nanocomposites synthesized with transition elements are important for their magnetic properties [1, 2], as well as the optical applications of composite layers on the basis of dielectrics with metal nanoparticles. As a matter of fact, when materials possessing long-range magnetic order are reduced in size, the magnetic order can be replaced by some other magnetic states such as superparamagnetism [3]. Especially, nanocomposite of magnetic granules dispersed in insulating matrixes have been very fascinating because of their unique properties associated with quantum-size effects and the possibility of applications as enhanced magnetic refrigerants and high density magnetic recording media [4]. One of the aims in this work is to review briefly the studies done to investigate the characteristics of the magnetic nanoparticles synthesized by ion implantation into insulating or dielectric materials. General characteristics of the implanted granules have been discussed in the preceding chapter 3. We are now concerned with TMR effects in Fe and Fe-alloy granules in Al2O3, MgO

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and SiO2 matrixes, and the features peculiar to the nanocomposites of implanted magnetic layers are described in this chapter. It is noted that the TMR effects in the granules prepared by implantation differs from those in the granular films prepared by sputtering technique and, above all, that the Fe-Co alloy granules implanted in Al2O3 matrix show a highest value of MR ratio of about 13 % under applied magnetic field of 1.1 T at room temperature [5, 6]. Conversion electron Mössbauer spectroscopy (CEMS) was used to investigate the physical properties of Fe and Fe-alloy granules including their size estimation. As the magnetic properties of these nanocomposites depend on the cluster structure, composition and size, it is important to prepare procedures for modifying and controlling these granular characteristics. Post implantation treatments are shown to be useful in modifying the size distribution of the implanted granules. We have done high energy Au ion irradiation through the implanted layer which resulted in the increase of TMR ratio. Finally, it is demonstrated that a combination of implantation and low temperature annealing are effective to improve TMR effects in Fe-Al2O3 granules.

2. EXPERIMENTAL TECHNIQUES FOR CHARACTERIZING TMR GRANULES Magnetic nanocomposites were synthesized by 57Fe ion implantation to high fluences of (0.7~3.5) × 1017 ions/cm2. Synthesis of binary nano-alloy granules was performed by implanting sequentially two species of metal ions. The alloy composition was determined with the relative fluence of two injected ions. Moreover, such ion beam technology has been used to modify the surface properties. Ion irradiation with high energy enough to penetrate through the pre-injected layers was done for tailoring granular structures. In the experiments high energy irradiation was performed by a 1.7 MeV Tandem accelerator (HVEE Corp., HV-4117HC) at Tokyo Metropolitan Industrial Technology Research Institute. CEMS is an effective means to characterize the implanted granules, providing quantitative information on the local magnetic properties around the Fe atoms that includes both the existing Fe phase distribution and the composition and size of the magnetic particles. 7.3 keV conversion electrons were detected by a He-CH4 gas flow proportional counter for CEMS at room temperature and by a channeltron detector for low temperature CEMS measurement, using a ~3 GBq 57Co source in Rh matrix. It is noted that the

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size distribution of the granules can be estimated from CEMS measurements with changing temperatures and/or applying external fields. This is due to the fact that the granules size is sensitively reflected on CEMS pattern through superparamagnetic relaxation. The relaxation effect is explained by the balance of the magnetic anisotropy energy barrier in single-domain particles against thermal energy, and in zero applied field it can be described by Arrhenius law, i.e., τ = τ0 exp( KV / kT),

(1)

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where τ is the relaxation time, k is Boltzmann’s constant, K is an anisotropy energy constant, T is the temperature, and V is the volume. The preexponential factor τ0 is of the order of 10-10 ~ 10-13 s. Thus, the increase in τ brought by T lowering or V increase causes the changes in magnetic order, i.e., superparamagnetic to ferromagnetic transition. Observation of the superparamgnetic behavior depends on the used instruments or their characteristic time constant τi [1], since the blocking temperature TBi measured by an instrument is defined by the equation, T Bi =

KV . k [ln( τ i/ τ 0 )]

(2)

In 57Fe Mössbauer spectroscopy τ i is 10-8 s (lifetime of 14.4 keV excited state). This means that the blocking temperature observed by CEMS with rapid time constant is higher than that by static measurements such as VSM or SQUID with τ i of ~10 s [1]. We can say that CEMS can more sensitively detect the transition in a granular solid. While six lines spectra appear below TBi, the CEMS pattern above TBi is similar to those of paramagnetic systems. Thus, the Mössbauer spectroscopy can be a useful means to analyze the magnetic properties of the individual species within nanocomposite layers and then the size distribution of the granules. The typical examples of the changes in CEMS spectra have been presented in Figures 5, 2, and 11 in Chapter 3. These demonstrated observations give us useful information on the granules characteristics, e.g., faster growth in Fe-SiO2 granules than Fe-Al2O3. The size estimation and crystalline structure of the granules were also monitored by glancing angle X-ray diffraction (GXRD). To check the structure of granular layers TEM observation was done using a Hitachi HF2000 field emission electron microscope operating at 200 kV. It was shown that TEM measurements combined with GXRD and CEMS are effective in

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characterizing structural and magnetic properties of the implanted granular layers [7]. The magnetoresistance of granular layers was measured at room temperature by a dc method employing two-probe contacts with silver paste and pressured contacts. The magnetization was measured with a vibrating sample magnetometer (VSM), applying perpendicular and in-plane external fields up to 1.1 T. All the measurements were performed at room temperature.

3．MAGNETIC PROPERTIES AND TMR EFFECTS IN IMPLANTED GRANULES

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3.1. Direct Implantation of Magnetic Elements When magnetic particles in composite systems are in the size of nm range, the particles have properties that are significantly different from those of the corresponding bulk materials. Magnetic nanocomposites may exhibit enhanced coercivity, thermally activated demagnetization process, changes in the hysteresis loops and magneto-transport properties [8]. A suitable way to reduce the dimension of ordered magnetic regions is to prepare the magnetic particles embedded in non-magnetic targets and isolate them. Consequently the granules become single-domain particles and the magnetization direction fluctuates spontaneously (superparamagnetic relaxation) with a characteristic time given by the equation (1). In the preceding chapter it was demonstrated that ion implantation is one of the most versatile methods to synthesize this kind of materials with the capability of introducing virtually any magnetic element in any solid substrate and overcoming solubility limits. Thus, the composite material is expected to gain new technological possibilities, e.g., in the field of magnetic recording substrates for high-density information storage. In order to elucidate details of the physical properties of nanocomposite systems it is desirable to perform measurements at both low and high temperatures and/or under large applied magnetic fields, so as to bring to light on phenomena smeared by the relaxation effect. A simple means to determine particle size is to measure the magnetic field dependence of Mössbauer spectra below the superparamagnetic blocking temperature TB. While above TB the magnetic hyperfine splitting in zero field spectra is collapsed due to the fast superparamagnetic relaxation, below TB the observed hyperfine field, Bobs, at the nucleus, when a magnetic field Ba is applied, is given as the approximated form of the equation (1) in Chapter 3 for

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large μBa /kT, Bobs = B0 (1 – kT / μBa) – Ba. A plot of (Bobs + Ba) as a function Ba-1 thus offers a rather simple method to determine the particles’ volume, because the plot gives a straight line whose slope and intercept give the values of μ and B0. Even for relatively large values of the magnetic anisotropy this procedure can be used to obtain reliable values of μ. If the saturation magnetization Ms of the particles is known, the particle volume, V = μ/Ms, can be calculated from Mössbauer spectra [9]. On the other hand, if one have any information of particle size, anisotropy energy constant K is estimated from the temperature dependency of Mössbauer spectra without applied field, using the relationship, Bobs = B0 (1 – kT /2KV). Bødker measured magnetic anisotropy energy constant for α -Fe particles as a function of the reciprocal particle diameter [10]. We simply estimated the anisotropy constant K by observing the change in CEMS spectra. Assuming that most granules in Figure 2 of Chapter 3 are distributed around 6 nm and in the critical size for blocking of superparamagnetic relaxation at 300 K, we obtained the K value as 2.0 × 105 J/m3, where 10-8 s and 10-10 s were assigned for the time constant τ i and τ 0 in equation (1), respectively. Although the estimation is rather rough, it is reasonably in agreement with the value (~1.2 × 105 J/m3) reported by Bødker et al.[10]. But, as it is larger than the anisotropy energy constant for bulk α–Fe (~0.5 × 105 J/m3), the additional contribution of the shape or surface anisotropy should be taken into account. The surface anisotropy constant for tinier particles would be smaller than that of thin iron film because of the small deviation from spherical symmetry [10]. Pereira de Azevedo et al. [11] applied the technique of ion implantation for preparing metal-granular magnetic films to investigate magnetization and magnetoresistance, where 57Fe ions were implanted in Cu and Ag thin films, immiscible with iron. They reported that the implanted Fe ions precipitate as either very small clusters or large iron α–phase particles; From the analysis of magnetization curve the former precipitations were suggested to be consisted of 5 atoms at most while the latter has more than 6000 atoms. These values were estimated from the value of magnetization and CEMS measurements. Whereas the MR ratio, Δρ/ρ, at low external fields H exhibited H 2 dependency in the similar way as usual GMR effect due to spin-dependent scattering of electrons passing between the larger Fe clusters, the Δρ/ρ measured at 4.5 K for higher field more than 1 T was observed to be dominated by a linear term in H. The linear field-dependency can be attributed to the electron scattering between the larger size clusters with blocked magnetization and the small superparamagnetic ones because in the latter granules the thermal average of magnetic moment is linear in H. However, the

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MR ratio itself is less than 0.3 % even at fields up to 13 T, due to the poor texture with the implanted granules occupying only a part of the sample. We notice that in section 3 of Chapter 1B more definitive analysis to distill the superparamagnetic contribution to GMR effect is presented for the sputtered Co-Ag films. A number of studies have been done on the tunneling type giant magnetoresistance in metal-oxides granular films prepared by rf sputtering. In chapter 2 theoretical and experimental works on the TMR were discussed, placing an emphasis on the Fe and Co granules in SiO2 matrixes. It was reported in the sputtered films that the highest MR ratios are 3.6 %, 3.8 % and 4.5 % at applied field of 1.5 T for Fe-, Co- and FeCo-SiO2 granules, respectively [12]. As for metal-alumina composites Huang et al. [13] have investigated the physical properties of Fe0.5 (Al2O3)0.5 granular films and observed a maximum MR ratio of 5 %. This MR ratio seems to be the highest of the sputtered Fe-Al2O3 films. In order to characterize the magnetic layers synthesized by implantation and search for advanced TMR effect we have performed systematic CEMS measurements, focusing on two subjects of 1) the fluence dependence of both electric resistance and MR ratio in the range of (1 ~ 3.5) × 1017 ions/cm2 and 2) the correlation of CEMS spectra with magnetic properties in granular systems. Magnetic and TMR characteristics were studied in Fe implanted Al2O3 granules, for the first time to our knowledge, by Sakamoto, et.al. [14]. Figure 1 (a) shows MR curves of Fe-Al2O3 granules plotted against external fields. In the granules implanted with Fe at a fluence of 1 × 1017 ions/cm2 the MR ratio, defined by Δρ/ρ = [ρ(0) – ρ(15)]/ρ(15), is observed to reach up to 7.5 ~ 8.5 % at room temperature, where ρ(0) and ρ(15) is the resistivity at external fields of 0 and 1.5 T, respectively. The MR ratio increased from 7.5 % to 8~9 % by annealing at 300oC. Thus, it is noteworthy that the obtained MR ratios are nearly twice as large as those of sputtered granular films [13, 15] and that the presented results were reproduced by Wakabayashi et al. [16], indicating the reproducibility and reliability of implantation technology. Figure 24(b) shows the in-plane magnetization curve measured with the same sample as Figure 24(a). The measured magnetization curve (solid line) was analyzed using Langevin function L(μH/kT) whose relationship to the magnetization M of superparamagnetic granules is given by, M = M s · L(μH/kT),

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(3)

TMR Effect in Granular Layers by Ion Implantation

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where M s is the saturation magnetization. The best fit to the equation was obtained by analyzing the experimental curve with two Langevin functions that correspond to two size distributions of the Fe-Al2O3 granules. Two broken lines in the figure represent the two fitted curves, and the dotted curve overlapping to the measured one is the sum of the two curves.

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  （ｂ ）

Figure 1. (a) MR curves for the Fe-Al2O3 granules implanted with a fluence of 1.0 × 1017 ions/cm2. The solid and broken lines are the experimental and calculated curves, respectively. (b) Magnetization curves for the granules. The solid and broken lines are the experimental and calculated curves, respectively; chain lines I and II indicate the superparamagnetic components from the small and large nanoparticles.

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10 M R ratio (%)

 

(A) 

8 6 4 2 0 0

1 2 3 17 Fluence (10 Fe/cm 2 )

4

10000

R 0 　（ M Ω）

(B)  1000

100

1

1

2

3

F luence (10

17

4 2

F e/cm )

1400

16 Intensity (counts)

  D iam eter

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10

12

2.5×1017 Fe/cm 2 1000

600

8 200

38

40

4

42

44 46 2θ (degree)

48

50

52

0

1

2

3

4

F luence (10 1 7 ions/cm 2 )

Figure 2. (a) Dependence of MR ratio in Fe-Al2O3 granules on Fe fluence. (b) Dependence of zero field resistance R0 on Fe fluence. (c) Change of nanoparticles size against fluence, obtained from the analysis of GXRD patterns. TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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From the analysis we obtained the magnetic moment of 3183 μB for the smaller and 7353 μB for the larger granules, respectively, and from the two values the granules radii were estimated as 1.2 nm and 1.6 nm for the small and large nanoparticles by assuming spherical shape. It is noted that these radii are in good agreement with the peak values in the histogram of Figure 15 in Chapter 3 which was obtained by the analysis of CEMS under applied magnetic field [16]; the samples used in these two measurements were prepared in separated experiments. The result is in contrast with the granule size of 3 to 10 nm in diameter reported for sputtered Fe-SiO2 granular films and Fe-implanted SiO2 granules, that were discussed in Chapters 2 and 3, respectively. Then it is possible to express MR curves by applying the field dependence of the magnetization, i.e., using the above Langevin functions. According to the empirical relation given by Honda et al. [17] and described in Chapter 2, MR ratio is expressed by the equation,

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Δρ /ρ = – α [MI(H)/M s I]2 – β [MII(H)/M s II]2

(4)

In the equation suffix I and II are used for M and M s of the small and large particles in the granules, respectively. The parameters α and β are the weight of contribution from the small and large particles to MR ratios, respectively. The calculated MR curve is shown by the broken line in Figure 1 (a). The best fit was obtained by taking α = 10.08 % and β = 0 %. This means that the larger particles seem not to contribute to magnetoresistance, even if they exhibit superparamagnetic behavior, and on the whole the result agrees with those found in sputtered granular films with low metal content, as shown in section 4.2 of Chapter 2. These results indicate that there should be an optimal fluence or size distribution of metal nanoparticles to achieve larger TMR effects, and we have done investigations searching for the optimal conditions; it is described in section 4-2. For this purpose it is necessary to take into account the experimental result that a percolation limit of the FeAl2O3 granules is in the fluence range of (1.5 ~ 2.0) × 1017 ions/cm2 where the ferromagnetic Fe 0 components become predominant. So it is useful to investigate the change in the magneto-transport and correlation of TMR effect with the magnetic characteristics of the granules by monitoring of the size and its distribution through the fluence up to 3.5 × 1017 ions/cm2. Figures 2 (A), (B) and (C) show the dependence of MR ratio, zerofield resistance, and the size of Fe0 nanoparticles on Fe fluence, respectively [18]; a spherical shape was assumed for the grain size estimation. It is noted in Figure

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N. Hayashi, T. Toriyama and I. Sakamoto

25 (A) that the MR ratio still gains a rather high value of 2.2 % even at such a high fluence as 3.5 × 1017 ions/cm2, while the maximum magnetoresistance was observed at around 1.2 × 1017 ions/cm2. On the other hand, in Figure 25 (B) the zerofield resistance R0 shows drastic but continuous drop over 4 orders for the studied fluence range. In Figure 25 (C) the diameters of the nanoparticles with two sizes are plotted against the fluence. The inset in Figure 25 (C) shows a typical GXRD pattern from bcc (110) diffraction of Feimplanted Al2O3 granules at 2.0 × 1017 ions/cm2, together with two fitted Gaussians. The measured GXRD curves apparently have much broader periphery than the expected diffraction, suggesting a superposition of two diffraction peaks. For example, the width of one tenth intensity in the diffraction from the implanted layers is wider by about 50 % than that from bulk polycrystalline grains. Therefore, they are supposed to represent a distinct size distribution of the large and small particles. This is the reason why we simply analyzed it with fitting of two Gaussians. It is difficult to do quantitative estimation of their intensities because of the simplified analysis, although the area of two fitted curves looks comparable. Nevertheless, it is plausible that an amount of the smaller particles exist even in the granules implanted at the high fluence. The size of the smaller particles itself is kept at around 4 ~ 5 nm, exhibiting only slight increase with fluence. It should be noted that the plots in Figure 2 (C) display behavior quite similar to that in Figure 12 (b) of Chapter 2 where the size changes in large and small clusters are plotted against the volume fraction of metal composition x in the sputtered Fe-SiO2 films. We can suggest that the granules in insulating oxide matrixes commonly have a bimodal distribution with two distinct size peaks. In CEMS measurement the blocking temperatures TBi for granules of 6 nm and 5 nm in diameters are estimated to be 300 K and 174 K, respectively, from equation (2) with the assumption of K = 2 ×105 J/m3. Thus, CEMS measurements at room temperature emphasize ferromagnetic characteristic of the granules, as shown in Figure 16 of Chapter 3 that claimed for the most clusters to be ferromagnetic in fluence ranges more than 2 ×1017 ions/cm2. When we take into account of magnetic percolation or magnetic correlation with nearby ferromagnetic clusters, it is likely that even the smaller granules with size of 3 – 5 nm, depicted in Figure 2 (c), behave ferromagnetically in the concerned fluence range. Franco-Puntes [19] reported that by the observation of magnetic force microscopy sputtered CoFe-Au granular films displayed domain structures with long-range magnetic ordering or magnetic percolation which were evolved at ferromagnetic grains contents well below volume percolation limits. On the other hand, since from the equation (2) TBi of the

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granules with 6 nm is estimated as 100 K for static VSM or MR measurements with longer time constant, the estimation leads to the possibility of observing still TMR effect even at the high fluence range, as shown in Figure 2 (a), i.e., the dependence of MR ratio on fluence is caused by the existence of the smaller granules presented in Figure 2 (c).

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Figure 3. Hysteresis loops of the Co4Ni1, Co1Ni1 and Co1Fe1 formed in silica matrix by sequential implantation. Inset depicts the details at low field [51].

The larger granules cause the resistance reduction by metallic conduction. It has been accepted that electron transport can be described with the parallel conduction of two current paths in the granular films characterized by a bimodal distribution of nanoparticles. One of them is tunneling conduction with exponential temperature dependence (see equation (4) in chapter 2), and the other is conduction flowing through connected nanoclusters. It has been experimentally confirmed that the latter becomes predominant at the high concentration of metal species over percolation limit [12]. The equation (5) in chapter 2,

1

ρ (0 )

=

1

ρt

⎧ exp ⎨ − 2 ⎩

c ⎫ 1 , manifests that the resistivity ρ(0) ⎬+ kT ⎭ ρ c

at zero field consist of both the tunneling, ρt, and temperature independent TMR Research in Insulating Granular Magnetic Materials, Nova Science Publishers, Incorporated, 2011. ProQuest Ebook

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N. Hayashi, T. Toriyama and I. Sakamoto

conduction, ρc, in the granular films. Thus, when ρc < ρt, the conductivity is mainly determined by ρc term, the resistivity of metal-like path. Figure 2 (b) indicates that the metallic conduction increases exponentially with ion fluence. Then, the variation in MR ratio with fluence in Figure 2 (a) could be explained by taking account of the tunneling and metallic conduction in parallel links in the granules. The shunting effects due to the metallic part of coalesced Fe0 particles cause the reduction in MR ratio. It is possible to estimate the shunting coefficient define by k = (Δρ/ρ)int/ (Δρ/ρ)meas [20], if necessary, when we take (Δρ/ρ)int as the maximum value in Figure 25 (a). (Δρ/ρ)int and (Δρ/ρ)meas denote intrinsic and measured MR ratios, respectively. 1.2E17C o/A l2O 3

2

M R R atio (% )

0

M R R atio (%)

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0

-3

(c) pure Fe

-2 -4 -6 -8

-6

-6

-4

-2

0

2

4

6

E xte rn al F ie ld (kO e )

-9

(b) FeC o(21%)

-12 (a) FeC o(25% -15 -12

-9

-6

-3

0

3

6

9

12

E xternal Field (kO e) Figure 4. Field dependence of MR ratio in FeCo-Al2O3 granules with Co concentration of 0 % (triangular), 21 % (dot) and 25 % (open circle).

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M R ratio

 

113

15 12 9 6 3 0

0

20

40

60

80

100

C o concentration (%)

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Figure 5. Concentration dependence of MR ratio in FeCo-Al2O3 granules.

Magnetic properties of alloy-based granules including Cu-Ni, Co-Ni and Co-Fe, that were prepared by co-implantation in silica, were investigated by Fernández et al. [21]. The alloy granules are materials of interest because the characteristics of bimetallic nanocomposites are influenced by both alloy composition and size effects. But, first of all, the alloying itself in the nanoparticles should be assured after sequential implantation of two metal species. It could be done by comparing the structural and magnetic properties with those of corresponding bulk alloys. The granules implanted with Co and Fe ions were observed to be soft magnet and to have bcc structure by TEM observation as expected in bulk FeCo alloys. Their magnetic properties were studied by measuring the hysteresis loop at 3 K [20]. Figure 3 shows magnetization curves for the granular samples of 1Co-1Ni, 4Co-Ni and 1Co1Fe, prepared by 180 keV ion implantation to a total fluence of 3 ×1017 ions/cm2. The figure indicates that the co-implantation of Co and Fe ions leads upto an increase of the magnetization to 2.2 μB from the magnetic moment of 1.6 μB in pure Co granules, and that implanting Ni in Co clusters decreases the magnetization. Although the alloy formation was inferred from these measurements, especially in the case of sequential implantation it is still necessary to investigate the microscopic physical properties, its compositional dependence, solid solution of two implanted species, etc., and also to confirm the alloying process.

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FeCo alloy granules are fascinating systems to investigate magnetotransport phenomena from the viewpoint of basic research and technical application. Many studies have been published on the TMR effect in Fe-Co nanocomposites in thin films prepared by rf sputtering or ion plating. The MR ratios in Fe-Co sputtered films have been reported to be 6.9 % at most by Wang et al. [22]; they cited it as the highest value in FeCo/Al2O3 sputtered films. Such slight increase of TMR was also found in FeCo-SiO2 sputtered film with the highest MR ratio of 4.5 %, compared to 3.6 % in Fe-SiO2 film (see Chapter 2). Since we can expect the enhancement due to the observed increase in magnetic moment of ion implanted Fe-Co nanocomposites, systematic studies on the TMR effect have been done with changing Co concentration in the alloy granules. Table 1 summarizes typical MR data and zero field resistance R0 for the Fe-Co granules with Co concentration of 0 100 % at the fluences of (1.2 – 1.5) × 1017 ions/cm2. Some typical MR curves of the granules are shown in Figure 4 [8, 9]. The inset in the figure presents a MR curve of Co/Al2O3 granules by implantation. The remarkably enhanced TMR effects are demonstrated in the FeCo alloy-granules, compared with Fe/Al2O3 and Co/Al2O3 granules, and the TMR reaches a value of more than 13 % at an applied field of ~1.1 T, that is so far the highest among the MR ratios reported for Fe/ and Co/Al2O3 granules at room temperature. The notice is that in the table and figure the ratios are not necessary the optimal on, despite the fluences were set with intention of obtaining higher MR ratio. For example, the lower MR ratio in FeCo40% than in FeCo43% granules is correlated with lower R0 and then the coalescence of FeCo nanoparticles in the former granules; here the notation of FeCo40% and FeCo43% stands for the FeCo-Al2O3 granular samples with 43 at. % Co and 45 at. % Co, respectively. For the implanted Co-Al2O3 granules a hysteresis was observed in the MR curve as well as in the magnetization measured by VSM. These features in appearance resemble the observation in Co-Al2O3 granules in rf sputtered film [23]. It was confirmed that the hysteresis of MR curve follows exactly the behavior of the magnetization in the granules and the field dependence is explained by the square of the normalized magnetization (M2) [23]. But the TMR value obtained in ion implanted granules is 6.3 % at H =1.1 T, much lower than ~10 % reported in Co-Al2O3 sputtered film. Although the MR magnitude of 10 % is the predicted value from the spin polarization P of 0.34, i.e., from the relation of P2 /(P2 + 1) [3], it is still necessary to examine the magneto-transport process in granular films including Co-Al2O3, as discussed below.

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1

Δ R /R 0 (%)

-2 -5 2

M fitting

M

1.6

fitting

-8 -11 -14 -1.4

-1

-0.6 -0.2 0.2 0.6 　ExternalField （T)

1

1.4

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Figure 6. Variation of MR ratio in FeCo-Al2O3 granules with 43 at. % Co against magnetic field, together with the dependence of magnetization M. Open circle is experimental data, solid curve is calculated one fitted by M1.6, and broken curve by M2.

Figure 5 shows plots of MR ratios against Co concentration in FeCoAl2O3 granules implanted to a total fluence range of (1.0 ~ 1.2) × 1021 ions/m2 [24]. The behavior of TMR change with Co concentration is very similar to that of hyperfine field B hf change shown in Figure 18 shown in Chapter 3, where the maximum of both MR ratio and B hf curves occurs around Co concentration of 25 at.%. The result comes up to our expectations, and this means that we should consider TMR effect to be influenced not by Fe magnetic moment μ(Fe’) but by the mean moment μ’ in Fe-Co granules, especially in high Co concentration range over 25 at. %. It is reasonable that the MR ratio take the values of 13.2 % and 6.3 % at Co concentrations 25 and 100 at. %, respectively, when we take account of the relation, {μ(Fe’)/ μ(Co)}2 ≈ 2.0, assuming that the MR ratio is proportional to square of magnetic moment [13]. Here μ(Fe’) = μ’ was assumed in the FeCo granules at Co composition of 25 at.% [25]. However, the difference in the MR ratio between Fe-Al2O3 and FeCo-Al2O3 granules can not be explained by the change in Fe magnetic moment. We could expect MR ratio of about 11 % in Fe-Al2O3 granules, if a relation of {μ’/μ(Fe)}2 = (2.4μB/2.22μB)2 is used in the estimation. But the highest value is 8.5 % at most even after post implantation annealing. When we think about TMR of 10 % reported in sputtered Co film,

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it is likely that the structural factor such as particle distribution will contribute to magnetoresistance in order to explain the additional increase of TMR effect in Fe-Al2O3. As for this respect, it is noted that Tsymbal and Balashchenko [26] have explained the enhanced TMR effect in sputtered Co-Al2O3 films by a polarization effect for tunneling electrons, where the polarization is supposed to be primarily determined by electronic and structural properties of the Co/insulator interfaces rather than by bulk properties in Co/Al2O3/Co tunnel junctions. Different technique for granules preparation may induce different type of interface Co-O bonding and result in different evolution of spin dependent conduction. We may expect further advance in TMR of FeCoAl2O3 granules by optimizing the morphological, magnetic and chemical properties. In these estimations it was implicitly assumed that MR ratio is proportional to the squares of the magnetization M in granular films. However, we have found that the magnetoresistance in the Fe-Co alloy granules exhibits a deviation from M2 dependency. Figure 6 shows a typical example of experimental and calculated MR curves plotted against external magnetic field for the FeCo43% samples implanted at a total fluence of 1.4 × 1017 ions/cm2. Apparently the curves are characterized by the high sensitivity to applied field, [27]. The S value is higher in whose coefficient is defined as S = d ( MR ) | dH

H →0

the FeCo than Fe-Al2O3 granules and S = 0.013 %/Oe was obtained for the FeCo43%. The sharp change near zero field indicates that it is necessary to compare the behavior of MR curve with the magnetization data that was obtained by VSM as a function of external field. Figure 6 reveals that the best fitting is obtained by taking M1.6 dependency instead of M2.0. Therefore, it is likely that electron tunneling occurs not only between superparamagnetic particles but the particles with a sort of magnetic ordering participate in the TMR effect, and that a part of nanoparticles in the FeCo granules are under magnetic correlations between neighboring clusters or magnetic percolation. Such percolation with respect to magnetic order was suggested to explain characteristic magnetoresistance in Co80Fe20/Al2O3 discontinuous multilayers with granular structures [28]. Santos et al. reported that in the nanocomposites the highest S value of ~ 0.025 %/Oe was obtained at the thickness of layered Co80Fe20 where the remanence and coercivity first appeared. A FeCo40% sample was prepared with Fe and Co ion fluences of 0.9 × and 0.6 × 1017 ions/cm2, i.e., with slightly higher Fe fluence than FeCo43%. In the FeCo40% granules MR ratio drops to 10.3 % at H = 1.1 T in spite that the composition is near to FeCo43% granules with MR ratio of 12.6 %. CEMS

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spectra from both samples were measured at liquid Helium temperature to reveal the origin of the different TMR characteristics, and are shown in Figure 7 [5]. The low MR ratio in FeCo40% seems to arise from the facts that FM sextet peaks appear in the CEMS at room temperature while they are hardly observed in FeCo43% CEMS. At the low temperature spectra the difference of the granules’ size in the two samples appears more clearly in the features that FM component in FeCo40% CEMS (A) is more intensive than in FeCo43% CEMS (B). The CEMS was analyzed under assumption that the spectra consist of two magnetic sextets, two single lines, and two quadrupole doublets from ferrous Fe2+ irons in the same way as Fe-Al2O3 granules. The single line components located near zero velocity are assigned to superparamagnetic state (SPM) of FeCo clusters (α-FeCo) and to γ-like FeCo state. Thus, the variation of CEMS patterns normally accompanies the change in magnetoresistance, as discussed in the followings. Some hyperfine parameters for the FM sextets in FeCo granules are summarized in Table 2 together with the parameters for Fe/Al2O3 granules implanted at a fluence of 1.5 × 1017 ions/cm2, for comparison. In the table the intensity ratio (spectral area ratio) of ferromagnetic to superparamagnetic peak (FM/SPM ratio) is given. The increase of Bhf values with Co concentration in the table gives us an evidence for Fe-Co alloying by sequential implantation, as mentioned before. The change in isomer shift (IS) offers another support for the alloying, because the IS values of Fe atoms increase with Co doping due to the increase in the number of Fe d-electrons. The IS increase at low temperature is due to the second-order Doppler shift. And among the hyperfine parameters the intensity of superparamgnetic phase or intensity ratio of FM/SPM is a key factor to considered TMR effect in FeCo alloy granules. Table 1. Zero field resistance R0 (H = 0) and MR ratio of Fe-, FeCo-, and Co-Al2O3 nano-composites Granules Fe Fe-Co Fe-Co Fe-Co Co

Co contents (at.%) 0 25 43 40 100

R0 (GΩ) 11.9 9.5 3.9 0.27 0.96

MR ratio (%) 7.5 13.2 12.6 10.3 6.3

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7320000

(A ) F eC o43% -4.2K

C ounts

7280000

7240000

7200000

-10

-8

-6

-4

-2

0

2

4

6

8

10

V elocity (m m /s)

C ounts

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3470000

(B ) F eC o40% -4.2K

3440000

3410000

3380000

-10

-8

-6

-4

-2

0

2

4

6

8

10

V elocity (m m /s) Figure 7. CEMS spectra taken from the samples FeCo43% (a) and FeCo40% (b), measured at 4.2K.

From the CEMS at 4.2K the FM/SPM ratio in the FeCo40% and FeCo43% granules indicates that there are a number of tiny nanoparticles whose superparamagnetic relaxation is not blocked even at the low temperature. When we take TBi = 4K as the blocking temperature and K = 2 × 105 J/m3 for anisotropy energy constant, the corresponding size is estimated as about 1.4 nm in diameter from equation (2). Thus, it is estimated from the values of FM intensity and FM/SPM ratio in Table 6 that the proportion of

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TMR Effect in Granular Layers by Ion Implantation

119

particles smaller than 1.4 nm in bcc phase amounts to 17 % of for FeCo40% and 26 % for FeCo43% granular sample. The larger proportion of the tiny granules in the FeCo43% samples is directly related with its higher TMR effect. Furthermore, it is estimated from the FM/SPM ratios in Table 6 that the granules of 5~6 nm sizes corresponding to TBi = 300 K amounts to 44 % of αFe precipitates in FeCo40% sample, leading to size shift to the larger one in the granular layer. We should remind that Mössbauer spectroscopy is highly sensitive to the magnetic ordering or blocking of superparamagnetic relaxation because of its rapid time constant. This means that some FM granules in CEMS spectra could possibly contribute to TMR effect by functioning as superparamagnetic particles in the measurement of magnetoresitance at room temperature. Table 2. Hyperfine parameters of the ferromagnetic sextets in Fe- and Fe-Co/Al2O3 granules

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Parameters

Fe/Al2O3 RT 4.2K

Fe-Co40%/Al2O3 RT 4.2K

Fe-Co43%/Al2O3 RT 4.2K

I.S. (mm/s) Area Intensity (%) Bhf (T)

0.01 0.113 32 49

0.062 0.185 27 53

(0.045) 0.177

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[42] H. Wakabayashi, S. Mochizuki, T. Toriyama, I. Sakamoto, and N. Hayashi. Proc. of the Specialist Research Meeting on New Developments in Solid State Physics with Probes of Radiations and Nuclei (KURRI-KR-106), Research Reactor Institute, Kyoto University, 2003, pp.183-188( in Japanese).

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INDEX

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A accelerator, 102 activation energy, 41, 57 aggregation, 61, 64, 67, 69, 80, 81, 86 anisotropy, 21, 84, 94, 103, 105, 118, 124 annealing, 7, 14, 15, 16, 17, 20, 25, 26, 27, 28, 52, 97, 106, 115, 126, 127, 128, 129, 130, 131 antiferromagnet, 3 Arrhenius law, 103 asymmetry, 20 atmosphere, 12, 60 atoms, 8, 9, 12, 16, 20, 24, 25, 61, 67, 69, 70, 72, 73, 75, 78, 80, 90, 91, 92, 94, 102, 105, 117, 124 Au nanoparticles, 93

B backscattering, 20, 34, 67, 68 basic research, 114 beams, 64, 66 behaviors, 122 bias, 1, 7, 9, 12, 20, 25 Boltzmann constant, 46, 81 bonding, 61, 73, 116 bulk materials, 104

C case studies, 60 cation, 73 ceramic, 60, 73 chemical, 60, 61, 67, 70, 72, 73, 116 clustering, 61, 66, 71, 73, 74, 86, 93 clusters, 60, 62, 68, 73, 74, 75, 80, 85, 90, 91, 92, 94, 105, 110, 113, 116, 117, 124, 127 collaboration, 131 collisions, 61, 73, 122 competition, 66 composites, vii, 60, 63, 86, 106, 117 composition, 1, 20, 22, 23, 27, 34, 35, 36, 48, 49, 55, 56, 64, 73, 85, 86, 89, 91, 102, 110, 113, 115, 116, 123 compounds, 70, 73 computer, 70 conductance, viii, 34, 35, 42, 49, 57 conduction, 26, 35, 42, 54, 111, 116, 122, 128 conductivity, 35, 36, 39, 112 configuration, 7, 60 constituents, 73, 77 contamination, 77 contradiction, 71 cooling, 93, 123 coordination, 75

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136

Index

copper, 91 correlation, 29, 106, 109, 110, 116 corrosion, 72 crystal structure, 68, 71 crystalline, 6, 7, 34, 35, 46, 61, 63, 64, 66, 68, 71, 74, 81, 85, 90, 97, 103 crystals, 79

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D database, 71 decomposition, 66 decoupling, 9 defects, 61, 72, 73, 78, 80 deposition, 7, 20, 63, 64, 90 depth, 60, 64, 65, 68, 69, 71, 72, 73, 129 detection, 74 deviation, 73, 105, 116 dielectrics, viii, 36, 59, 101 diffraction, 15, 71, 74, 91, 110 dislocation, 73 disorder, 80 dispersion, 60, 85, 93 displacement, 61, 73, 79, 123 distribution, 60, 65, 66, 69, 71, 74, 76, 78, 81, 83, 84, 85, 87, 90, 91, 92, 93, 94, 97, 102, 103, 109, 110, 111, 116, 119, 124, 125, 127, 130, 131 domain structure, 110 doping, 117

E electric current, 35 electron, 26, 36, 61, 68, 70, 71, 88, 97, 102, 103, 105, 111, 116, 130 emission, 71, 103 energy, 8, 12, 36, 51, 57, 60, 64, 70, 76, 84, 86, 93, 94, 97, 102, 103, 105, 118, 122, 127, 129, 130 engineering, 59 environment, 75 equilibrium, 60, 61, 90 equipment, 69

evaporation, 1, 19 evidence, 75, 87, 117 evolution, 78, 80, 94, 116, 124 EXAFS, 61, 67 excitation, 90, 93, 122 external magnetic fields, 125

F fabrication, 60 ferromagnetism, 20, 27, 44, 54, 90 ferrous ion, 78 film thickness, 20 flexibility, 60 flight, 12 force, 110 formation, 12, 60, 61, 62, 73, 74, 86, 90, 94, 97, 113, 123 formula, 22, 35, 71, 87 free energy, 84

G geometry, 64, 68, 97, 130 Germany, 70 Gibbs energy, 73 grain size, 20, 26, 35, 45, 74, 109 growth, 27, 66, 71, 80, 94, 103 guidance, 122, 131

H hardness, 72 height, 36 helium, 70 histogram, 65, 94, 109 homogeneity, 60 host, 60 hyperfine interaction, 70 hysteresis, 2, 8, 9, 48, 104, 113, 114 hysteresis loop, 2, 104, 113

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Index

I identification, 71 image, 2, 3, 9, 12, 35, 39, 71 impurities, 61, 67, 72, 78, 80 inevitability, 66 insulators, vii, viii, 33, 61, 63, 68, 72 interface, 4, 6, 16, 19, 116, 123 internal field, 91, 97 iodine, 124 ion bombardment, 13 iron, 61, 64, 67, 69, 74, 75, 78, 80, 85, 86, 87, 88, 91, 96, 105, 123, 127 irradiation, 60, 62, 73, 93, 94, 95, 96, 97, 102, 122, 123, 124, 125, 128, 129, 130 isotope, 64

J

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Japan, 1, 33, 59, 98, 101, 132, 133

L lattice parameters, 74, 75, 87, 91 lead, 16, 94, 123 lifetime, 103 light, 3, 59, 104 local configuration, 78 low temperatures, 69, 97 LTD, 64 Luo, 98

M magnet, 69, 113 magnetic characteristics, 97, 109, 127 magnetic field, 9, 34, 66, 69, 78, 81, 87, 91, 96, 97, 102, 104, 109, 115, 116, 119, 124 magnetic moment, 3, 4, 18, 19, 22, 24, 25, 29, 44, 46, 51, 81, 88, 90, 105, 109, 113, 114, 115, 130

137 magnetic particles, 91, 102, 104 magnetic properties, vii, viii, 7, 20, 34, 59, 84, 91, 93, 97, 101, 102, 103, 104, 106, 113, 123, 130 magnetic relaxation, 69 magnetism, vii, 1, 27 magnetizations, 46 magnetoresistance, vii, viii, 3, 4, 6, 7, 8, 18, 33, 34, 48, 58, 60, 81, 104, 105, 106, 109, 110, 116, 117, 120, 122, 123, 128, 130, 131 magnitude, 80, 114 masking, 16 mass, 36, 62, 64, 67, 73 material surface, 72 materials, vii, viii, 1, 2, 21, 60, 64, 67, 73, 101, 104, 113 matrix, 19, 43, 48, 59, 60, 66, 68, 70, 71, 74, 75, 80, 85, 86, 90, 93, 94, 97, 102, 111, 123, 124, 127 matrixes, viii, 60, 61, 68, 70, 74, 81, 84, 101, 102, 106, 110 matter, viii, 80, 101, 129 measurement, 8, 68, 74, 90, 97, 102, 110, 119 measurements, viii, 34, 61, 64, 68, 71, 73, 74, 84, 86, 93, 94, 103, 104, 105, 106, 109, 110, 113, 119, 122, 123, 124, 128, 130 media, viii, 101 melt, 90 melting, 123 metal ion, 60, 64, 73, 102 metal nanoparticles, viii, 59, 60, 101, 109 metals, vii, viii, 19, 33, 35, 42, 49, 61, 63, 72, 73, 86, 88 microscope, 71, 103 microscopy, 110 microstructure, 2, 73 microstructures, 71 migration, 61 mixing, 61, 64, 92, 94 morphology, 7, 9, 18, 66, 71, 93, 94 multilayer films, vii, 60

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138

Index

N nanocomposites, viii, 59, 61, 64, 70, 77, 85, 87, 90, 94, 101, 102, 104, 113, 114, 116, 127 nanometer, vii, 59, 71 nanometer scale, 59, 71 nanoparticles, 60, 66, 71, 74, 75, 78, 84, 85, 86, 87, 88, 90, 93, 94, 97, 101, 107, 108, 109, 111, 113, 114, 116, 118, 119, 124, 125, 129, 130 neutral, 36 NMR, 16 nonequilibrium, 78 nonequilibrium systems, 78 novel materials, 85 nuclear magnetic resonance, 16 nucleation, 80 nuclei, 68 nucleus, 70, 75, 104

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O opportunities, 60 optical properties, 85 oscillation, 8, 9, 14 Ostwald ripening, 80, 94, 123, 124, 127 oxidation, 45, 60, 70 oxygen, 73, 79, 80

P parallel, 42, 46, 84, 94, 111, 123 percolation, 35, 39, 44, 48, 54, 56, 57, 58, 80, 85, 109, 110, 111, 116, 122, 127 periodicity, 15 perpendicular magnetic anisotropy, 2, 21 phase diagram, 19 Philadelphia, 71 physical properties, viii, 7, 34, 60, 64, 71, 77, 85, 97, 102, 104, 106, 113, 128 polarization, 49, 55, 57, 79, 114, 116

polyimide, 4 precipitation, 60, 66, 73, 74, 77, 78, 81, 86, 93, 124 preparation, viii, 20, 26, 34, 60, 64, 116, 129 probability, 26, 36, 78 probe, 61, 68, 104 proportionality, 90 protective coating, 60

Q quantitative estimation, 110, 123 quartz, 67

R radial distribution, 60 radiation, 7, 34, 38, 61, 66, 71, 94, 123 radius, 24, 25, 27, 28, 46, 80, 91 reactivity, 60 recommendations, iv redistribution, 94, 123, 129 regeneration, 127 relaxation, 61, 70, 74, 87, 91, 103, 104, 118, 124 reliability, 106 researchers, 4 resistance, 4, 7, 19, 72, 106, 108, 109, 111, 114, 117, 120, 122, 126, 127, 130, 131 response, 61 retardation, 71 room temperature, 5, 7, 20, 21, 34, 36, 44, 50, 58, 64, 67, 68, 70, 80, 81, 86, 90, 102, 104, 106, 110, 114, 117, 119

S saturation, 3, 8, 9, 14, 18, 19, 22, 23, 24, 25, 26, 27, 30, 44, 46, 55, 58, 81, 84, 88, 93, 105, 107 scattering, 4, 6, 19, 20, 26, 29, 105

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Index semiconductors, 35, 36, 60, 72 sensitivity, 48, 116, 123 shape, 14, 24, 84, 94, 105, 109 showing, 96, 97 silica, 66, 72, 73, 86, 93, 111, 113 simulation, 74 SiO2 films, viii, 34, 35, 36, 37, 38, 40, 41, 43, 44, 46, 47, 48, 49, 50, 53, 54, 55, 56, 57, 110, 122 solid matrix, 60 solid phase, 72 solid solutions, 86 solubility, 60, 80, 104 solution, 80, 90, 93, 113 species, viii, 33, 60, 64, 73, 85, 93, 97, 102, 103, 111, 113, 122, 124, 127, 129 spectra analysis, 81 spectral component, 85 spectroscopy, 34, 61, 67, 86, 91, 97, 102, 103, 119, 130 spin, vii, viii, 1, 3, 4, 6, 9, 19, 20, 26, 29, 33, 49, 55, 57, 80, 105, 114, 116 state, 3, 20, 51, 70, 73, 80, 103, 117, 124 states, viii, 14, 61, 67, 69, 70, 72, 78, 80, 81, 94, 101, 119, 124 stoichiometry, 73 storage, 104 structural changes, 61, 123 structural modifications, 128 structure, 2, 6, 7, 9, 12, 15, 16, 18, 20, 24, 26, 34, 35, 46, 52, 60, 66, 71, 73, 85, 90, 93, 94, 96, 102, 103, 113, 123 substrate, 1, 2, 4, 7, 12, 52, 60, 73, 81, 104 surface, 64, 68, 102, 97

T tangles, 73 target, 7, 20, 61, 64, 66 techniques, 60, 63, 67, 85, 90, 123 technology, 60, 64, 102, 106

139 temperature, viii, 7, 21, 34, 35, 39, 42, 46, 52, 64, 66, 70, 74, 81, 86, 91, 93, 102, 103, 104, 111, 117, 118, 127, 130 temperature annealing, 102 temperature dependence, 7, 35, 39, 111 texture, 106 thermal energy, 103 thermal treatment, 86, 94 thin films, 105, 114 transformation, 91 transition elements, viii, 59, 101 transition metal, 1 translation, 36 transmission, 2, 34 transmission electron microscopy, 2, 34 transport, 48, 104, 109, 111, 114 transportation, viii, 34, 35, 57 treatment, 126, 127, 128, 129, 130 tunneling, vii, viii, 33, 34, 35, 42, 48, 57, 58, 81, 106, 111, 116, 122, 127, 128, 130, 131

V vacancies, 73, 79 vacuum, 7, 20, 64 valence, 61, 68, 70 variations, 7, 12, 27, 35 velocity, 61, 81, 117

W workers, viii, 34, 35

X XPS, 61, 67, 72, 73, 74 X-ray diffraction, 7, 15, 16, 34, 38, 46, 67, 71, 86, 103, 130

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