Electrodeposited multilayer films with giant magnetoresistance (GMR): Progress and problems

Progress in Materials Science 55 (2010) 107–245

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Electrodeposited multilayer films with giant magnetoresistance (GMR): Progress and problems I. Bakonyi *, L. Péter Research Institute for Solid State Physics and Optics, Hungarian Academy of Sciences, P.O.Box. 49, H-1525 Budapest, Hungary

a r t i c l e

i n f o

Article history: Received 27 April 2009 Received in revised form 13 July 2009 Accepted 16 July 2009

a b s t r a c t The giant magnetoresistance (GMR) effect was discovered in 1988 in nanoscale metallic ferromagnetic/non-magnetic (FM/NM) multilayers. By now, devices based on this phenomenon have been widely commercialized which use multilayered structures manufactured via physical deposition (PD) methods, mainly sputtering. It was shown in the early 1990s that electrodeposition (ED) is also capable of producing multilayered magnetic nanostructures exhibiting a significant GMR effect. These layered structures include multilayer films similar to those prepared by PD methods on macroscopic substrates and multilayered nanowires deposited into nanosized template pores, the latter ones being unique to the ED technique. Whereas ED multilayered nanowires can exhibit a GMR effect comparable to the values obtained on PD multilayer films, the GMR values achieved on ED multilayer films still remain inferior to them and, quite often, require high magnetic fields for saturation. Therefore, in spite of the relative simplicity and cost-effectiveness of the ED method, the GMR characteristics of ED multilayer films are still not competitive with the corresponding parameters of their PD counterparts. The main purpose of the present review is to give a summary of the progress achieved over the last one and a half decades on ED multilayer films with GMR effect and to critically evaluate the GMR results reported for various element combinations accessible to the ED technique for the preparation of FM/NM multilayer films (ED multilayered nanowires will be treated very briefly only). In order to promote an understanding of the inferior behavior of ED multilayer films, a detailed discussion of the magnetoresistance effects occurring in bulk homogeneous ferromagnets as well as in magnetic nanostructures (FM/NM multilayers and granular alloys) will be provided. Particular attention will be paid to the case

* Corresponding author. Tel.: +36 1 392 2628; fax: +36 1 392 2215. E-mail address: [email protected] (I. Bakonyi). 0079-6425/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.pmatsci.2009.07.001

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of non-ideal magnetic nanostructures which contain both FM and superparamagnetic (SPM) regions. This is an essential ingredient in explaining the high saturation field of GMR commonly observed in ED multilayer films. In addition to the GMR magnitude, this is another characteristic decisively influencing the magnetic field sensitivity, a key feature concerning applications in sensor devices. The controversial results reported for the spacer layer thickness dependence of GMR in ED multilayer films will also be discussed. It is pointed out that the still inferior GMR characteristics of ED multilayer films can be to a large extent ascribed to microstructural features leading to the appearance of SPM regions, pinholes in the spacer layers and probably not sufficiently perfect interfaces between the FM and NM layers. The origin of the latter deficiency is not yet well understood although it is clearly one of the main causes of a weak interlayer coupling (if there is any coupling at all) and, thus, a small degree of antiparallel alignment leading to a reduced GMR effect. Works will also be described in which attempts were made to produce ED multilayer films with view on possible applications in GMR sensor devices. Finally, problems will be identified which should still be solved in order to make the properties of ED multilayer films attractive for GMR applications. Ó 2009 Elsevier Ltd. All rights reserved.

Contents 1.

2.

3.

4.

5.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.1. Background and scope of the review: GMR in electrodeposited multilayer films . . . . . . . . . . . 1.2. GMR in ED multilayer films and ED multilayered nanowires . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.3. Giant magnetoresistance in magnetic nanostructures and giant magnetoimpedance in bulk ferromagnets . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Magnetoresistance in metallic magnetic materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. Definition of magnetoresistance and overview of various magnetoresistances . . . . . . . . . . . . . 2.2. Measurement of magnetoresistance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. AMR in bulk homogeneous ferromagnets . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4. GMR in magnetic nanostructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4.1. GMR in FM/NM multilayers. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4.2. Exchange coupling and oscillatory GMR in multilayers. . . . . . . . . . . . . . . . . . . . . . . . . 2.4.3. GMR in granular magnetic alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4.4. GMR in multilayers containing SPM regions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary of FM/NM multilayer film preparation by electrodeposition . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Electrochemical background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Single-bath technique with pulse plating . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1. Bath composition and its relation to the deposit composition . . . . . . . . . . . . . . . . . . . 3.2.2. Deposition pulse modes, pulse combinations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.3. Relationship with other chemically modulated structures . . . . . . . . . . . . . . . . . . . . . . 3.3. Dual-bath technique . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4. Cell construction, electrodes and electrode arrangement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5. Role of additives and pH . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.6. Control of electrochemical processes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Structure of ED multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. TEM study of ED multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. XRD study of ED multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3. Influence of preparation conditions on structure in ED multilayer films . . . . . . . . . . . . . . . . . . Critical evaluation of GMR results reported on ED multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1.1. Overview of systems investigated and preparation methods applied . . . . . . . . . . . . . 5.1.2. Guidelines in evaluating the experimental results on GMR in ED multilayer films . .

109 109 111 112 113 113 113 115 116 116 119 120 125 128 128 130 130 133 135 136 137 140 143 145 146 149 152 155 155 155 155

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5.2.

6.

7. 8.

Co–Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 5.2.1. General overview of deposition conditions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 5.2.2. Optimization of deposition parameters . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 5.2.3. GMR of ED Co–Cu/Cu multilayer films prepared in P/P or G/P mode from a single bath 159 5.2.4. GMR of ED Co–Cu/Cu multilayer films prepared in G/G mode from a single bath . . . 182 5.2.5. GMR of ED Co/Cu multilayer films prepared by dual-bath technique . . . . . . . . . . . . . 185 5.2.6. Influence of Cu content in the magnetic layer on GMR in ED Co-Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 187 5.2.7. Dependence of GMR in ED Co-Cu/Cu multilayer films on Cu layer thickness: evidence for a non-oscillatory GMR behavior and for the absence of coupling . . . . . . . . . . . . . 189 5.2.8. Temperature dependence of GMR in ED Co–Cu/Cu multilayer films . . . . . . . . . . . . . . 192 5.3. Co–Ag/Ag multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 195 5.4. Co–Au/Au multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 197 5.5. Co–Ru/Ru multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 197 5.6. Co–Zn/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 198 5.7. Ni-Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199 5.7.1. Deposition from a single sulfamate/sulfate bath in P/P mode . . . . . . . . . . . . . . . . . . . 199 5.7.2. Deposition from a single sulfate/citrate bath in G/G mode. . . . . . . . . . . . . . . . . . . . . . 203 5.7.3. Deposition by dual-bath techniques . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 206 5.8. Co–Ni–Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 207 5.8.1. Deposition from variants of sulfamate/sulfate type baths in P/P mode . . . . . . . . . . . . 207 5.8.2. Deposition from a sulfamate bath in G/0/G mode . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 216 5.8.3. Deposition from a sulfate/citrate bath in P/P mode. . . . . . . . . . . . . . . . . . . . . . . . . . . . 217 5.8.4. Deposition from a sulfate bath with additives in G/G mode. . . . . . . . . . . . . . . . . . . . . 217 5.8.5. Deposition from a sulfate bath in G/P mode . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 218 5.9. Co–Ni–Ag/Ag multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 218 5.10. Fe–Co–Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 218 5.11. Fe–Ni–Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 219 5.12. Co–Zn–Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 220 5.13. Fe–Co–Ni–Cu/Cu multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 221 GMR of ED spin-valve type structures and multilayers with view on possible applications. . . . . . . . . 222 6.1. Electrodeposition of multilayer films on semiconductor wafers . . . . . . . . . . . . . . . . . . . . . . . . . 223 6.2. Application of the GMR effect in magnetoresistive sensors. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 224 6.2.1. Field-sensitivity of GMR multilayers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 224 6.2.2. GMR spin-valve type structures and applications of GMR sensors . . . . . . . . . . . . . . . . 225 6.3. GMR results on ED pseudo spin-valve type structures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227 6.4. GMR results on application-oriented ED multilayer films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 231 Summary of progress achieved and problems ahead with GMR in ED multilayer films . . . . . . . . . . . . 233 Concluding remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 236 Acknowledgements. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 237 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 237

1. Introduction 1.1. Background and scope of the review: GMR in electrodeposited multilayer films In recent years, new devices produced by nanofabrication routes have become commercialized which operate by using the two possible spin states (spin-up and spin-down) of electrons. The appearance of such nanoscale devices has laid the foundation of spin-electronics (or spintronics) [1], an industry of broad future perspective. This progress has been made possible via the rapid development of thin film technologies for the preparation of nanoscale metallic layered structures (multilayers). In such structures, the thickness of the constituent layers can be smaller than the characteristic length scales of electron transport. In case one of the constituent layers is ferromagnetic (FM) and the magnetization orientation changes on a scale smaller than these characteristic lengths, so-called spindependent electron transport phenomena different from those known in bulk materials may arise

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due to the travel of conduction electrons between two adjacent FM layers through a non-magnetic (NM) spacer layer. An outstanding example of such phenomena is the giant magnetoresistance (GMR) effect discovered independently by Grünberg [2] and Fert [3] in metallic FM/NM multilayer films in 1988–1989, which discovery was awarded with the Nobel Prize in Physics in 2007 [4]. The magnetoresistance is the change in the electrical resistance of the material under study upon the application of an external magnetic field. The physical mechanism of the GMR effect of FM/NM multilayers differs basically from the mechanism of magnetoresistance in homogeneous bulk FM metals and alloys (the latter known as anisotropic magnetoresistance = AMR). Furthermore, the size of the GMR can be by an order of magnitude larger than the AMR; this is the origin, from one of the first reports [3], for the name ‘‘giant”. It should be noted, however, that in the current review, we will assign the name giant magnetoresistance if the underlying physical mechanism is the same as was in the original discovery of GMR (i.e., spin-dependent scattering of electrons when travelling through a nanoscale non-magnetic metallic spacer between two magnetic regions having their magnetizations not fully aligned), irrespective of the size of the actual magnetoresistance change (e.g., even if smaller than for AMR). The measurement/detection of magnetic fields is of great technological importance and a large variety of magnetic field sensors have been developed [5]. Evidently, the above described magnetoresistance effects can also be used for this purpose and magnetoresistive sensors based on the AMR effect found widespread applications for measuring or detecting a magnetic field already several decades ago. In present-day magnetic hard disk drives, the read-out heads operate almost exclusively with the GMR principle by utilizing a so-called spin-valve structure. Their introduction in 1997 made it possible to further maintain the high rate of magnetic storage density increase and GMR performance ensures that this trend can continue even in the future [6]. The significant progress achieved in the 1970–1980s in thin film technologies (evaporation, sputtering and, especially, molecular-beam epitaxy = MBE [7]) was a prerequisite that enabled the preparation of fairly defect-free very thin metallic films and FM/NM multilayers necessary for the observation of the new phenomena described above. The difficulties of the preparation of layered structures capable of operating as spintronic devices can be illustrated with the fact that the individual layer thicknesses should indeed be in the nanometer range whereby, for metals, a 1 nm thick layer comprises about five atomic layers only. This is a great challenge for materials technology. The details of the preparation methods have been properly elaborated for the above listed physical deposition (PD) techniques of which sputtering is mainly used for fabricating nanoscale spintronic devices. Electrodeposition (ED) has long been considered as a viable alternative to PD techniques to provide a simple and cost-effective technology for the preparation of high-quality films and multilayers [8]. The extended research activity on ED multilayer films with GMR behavior resulted in more than 140 papers [9–149] in the last 15 years (note: for a distinction between ED multilayer films and ED multilayered nanowires, see Section 1.2). In spite of these efforts, many aspects of the GMR characteristics of ED multilayer films, especially the application-relevant field sensitivity, have still remained inferior with respect to the corresponding parameters of PD multilayer films. This fact has permanently provided motivation for the continuing research efforts in this field. The differences in properties include the smaller magnitude of the GMR and the usually larger saturation field of the magnetoresistance of ED multilayer films, both factors contributing to a reduced field-sensitivity. In addition, whereas clear oscillations in the GMR magnitude have been demonstrated for many (but by far not all) PD multilayer films, the presence or absence of such oscillations has remained a controversial issue in ED multilayer films till now. It must also be admitted that in the case of the ED technique, reproducibility is still a problem: whereas the basic GMR features appear to be roughly the same when comparing the results of various authors for a given system, some ‘‘outstanding” results reported could be definitely not reproduced by other researchers. It has also been the experience that due to some still unknown details the GMR magnitude cannot always be reproduced under nominally identical preparation conditions. The first report demonstrating a GMR effect in ED multilayer films was published in 1993 [9]. Early works on ED multilayer films with GMR were summarized by Schwarzacher and Lashmore in 1996 [26]. In a recent review [137], some aspects of the latest progress made have been summarized, by

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focusing on some general features of the electrodeposition process and the GMR of ED multilayer films. The current review aims at giving a detailed critical evaluation of all observations reported on the GMR effect in ED multilayer films [9–149]. This review is organized as follows. Section 2 provides a general discussion of the magnetoresistance and a comparison of the magnetoresistance effects in homogeneous ferromagnets and in various magnetic nanostructures. A summary of multilayer film preparation by electrodeposition and the control of electrochemical processes will follow (Section 3). Section 4 will be devoted to the methods of studying multilayer structure and to the general microstructural features revealed for ED multilayer films. The GMR results reported on ED multilayer films in the literature will be critically evaluated in Section 5 for the electrochemically accessible element combinations of the FM/NM multilayers. Section 6 will be devoted to introducing some aspects which may be important for the application of GMR multilayer films and spin-valve type film sandwich structures. Section 7 will briefly summarize (i) the progress achieved in the understanding of how microstructure formation in these systems is governed by the underlying electrochemical processes and (ii) some issues which should be seriously dealt with for an improvement of the GMR characteristics of ED multilayer films in order to make them competitive for magnetic-field sensor applications. Some concluding remarks follow then in Section 8. However, before proceeding, two issues will be first discussed in order to put the GMR in ED multilayer films into proper perspective in comparison with ED multilayered nanowires and to describe briefly the difference between giant magnetoresistance in magnetic nanostructures and giant magnetoimpedance in bulk ferromagnets. 1.2. GMR in ED multilayer films and ED multilayered nanowires It was mentioned above that, similarly to the PD method, multilayer films exhibiting a significant GMR effect can be produced also by ED techniques. Such PD or ED multilayer films are obtained on the surface of fairly smooth flat substrates with typical lateral dimensions of at least several millimeters and with the total deposit thickness usually not exceeding 1 lm. It is a unique feature of the ED method that, by using nanosized template materials [150], it is also capable of preparing high aspect-ratio nanowires (typically 50–100 nm in diameter and several micrometers in length) with a nanoscale multilayered structure. The first GMR studies of such ED multilayered nanowires were reported some 15 years ago [151]. It has been shown [152,153] since then that the GMR in these ED multilayered nanowires can be much higher than in ED multilayer films and can be as high as the values reported for corresponding PD multilayer films. In spite of the importance of multilayered nanowires and their unique properties very attractive for specific applications, in the current review, discussion will be intentionally restricted to the case of ED multilayer films only for several reasons. Due to the inherent differences in the electrochemistry for the two geometries as well as in the typical individual layer thicknesses, quite different approaches are required for the two cases, both for the preparation and for the relevant physical phenomena involved. First, there are significant differences in the electrochemical conditions and problems when comparing the preparation of multilayer films and multilayered nanowires. Whereas a multilayer film is deposited on a macroscopic-sized flat surface, multilayered nanowires are prepared in a rather restricted geometry since deposition is performed in nanoscale pores previously formed with a high areal density (ca. 1011 nanowire/cm2) in track-etched polycarbonate membranes or in anodized aluminum oxide templates. A major issue is that the transfer of electrochemical conditions between the two kinds of multilayers is not straightforward; a few remarks on this will be made in Section 3.2.3. Second, the physical parameters being decisive for the magnetotransport properties are also partly different for the two kinds of magnetic nanostructures [154–156]. This becomes clear if we take into consideration that for multilayer films the magnetoresistance is measured in the so-called current-in-plane (CIP) mode whereas for multilayered nanowires in the current-perpendicular-to-plane (CPP) mode as will be discussed in Section 2.2. As a consequence, there are differences in the GMR magnitude in favor of the multilayered nanowires. As it was pointed out by Valet

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and Fert [154], the characteristic length-scales governing the size of the GMR in FM/NM layered structures are not the same for the two measuring geometries. For multilayer films measured in the CIP geometry, the electrons average the properties of the multilayer in the perpendicular direction on the length scale of the electron mean-free path (kmfp), the average distance over which an electron can travel without undergoing a (spin-conserving) bulk-type scattering in the layer material. This implies that the GMR vanishes when the multilayer period (bilayer length) becomes larger than kmfp. On the other hand, for multilayered nanowires measured in the CPP geometry, the relevant characteristic length-scale is the spin-diffusion length (ksdl) being the average distance an electron can travel without undergoing a spin-flip process. Since usually the spin-diffusion length is much larger than the electron mean-free path (ksdl  kmfp), as a consequence of the differences in the magnetotransport-measurement geometry, the typical layer thicknesses for obtaining significant GMR values are much larger for multilayered nanowires (CPP mode) than for multilayer films (CIP mode). On the other hand, this layer thickness difference influences the preparation in the sense that maintaining a thickness uniformity and reproducibility is easier at larger layer thicknesses. In view of all the above listed differences between the two kinds of ED magnetic nanostructure, the present authors had in mind a discussion of GMR in ED multilayer films only since it was considered that covering this topic alone needs an extensive review in itself. The structure of the review was formed in a way to give a proper account of electrochemical and GMR issues specifically pertinent to the case of ED multilayer films. By taking into account the wealth of literature on multilayered nanowires, their description would require a separate discussion what is beyond the scope of the present review. The interested readers should consult reviews on electrodeposited multilayered nanowire arrays [26,133,152,157] or specific research papers with attempts to investigate the magnetotransport properties of a few, or even a single multilayered nanowire [158]. It follows from the above described topical restriction of the current review that all GMR results on ED multilayer films to be discussed later refer to measurements in the CIP geometry only. 1.3. Giant magnetoresistance in magnetic nanostructures and giant magnetoimpedance in bulk ferromagnets In view of the fact that quite recently a review has been published in this journal on the so-called giant magnetoimpedance (GMI) effect [5], it might be appropriate here to point out the difference between GMI and GMR, both including the term ‘‘giant”. The GMI is a classical effect which can be induced by subjecting first a FM conductor to a small alternating current (ac) which then creates an alternating magnetic field within the conductor, causing periodic magnetization reversals. If now this ac current-carrying FM conductor is placed into a static external magnetic field, its complex impedance changes. This is due to the random domain structure of a bulk FM material in zero magnetic field which is then saturated (magnetically homogenized) with an externally applied magnetic field, as a consequence of which the magnetization reversal mechanism can be quite different. The GMI effect is the difference between ac complex impedances in the demagnetized and saturated state. The actual dominating physical mechanism of GMI changes with the frequency of the ac current due to changes in magnetization reversal processes with frequency (measuring frequencies range up to the GHz region). The saturation fields are kept low (typically below 100 Oe) by using a soft ferromagnetic material in a proper geometry (e.g., ribbon or wire). In principle, the GMI effect can occur in any bulk ferromagnetic material (either homogeneous or inhomogeneous). By contrast, the GMR effect can be observed in appropriate metallic magnetic nanostructures only for which the magnetization direction in zero magnetic field changes on a scale smaller than the spin-diffusion length of conduction electrons. This situation leads to a high-resistance state. When in an applied magnetic field the magnetization of each individual magnetic region is aligned parallel, this results in a low-resistance state and the difference between the resistances of the two states just provides the GMR effect. The observation of a GMR effect can be achieved even with a dc measuring current since it relies on the spin-dependent scattering processes of conduction electrons only and is not related to the actual remagnetization mechanism.

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2. Magnetoresistance in metallic magnetic materials 2.1. Definition of magnetoresistance and overview of various magnetoresistances As already noted in the Introduction, the phenomenon of magnetoresistance (MR) is the change in the electrical resistance of the material under study upon the application of an external magnetic field H: DR = RH  Ro where RH and Ro are the resistances measured in the presence and absence of a magnetic field, respectively. It is customary to define a magnetoresistance ratio in the following form:

MRðHÞ ¼ DR=Ro ¼ ðRH  Ro Þ=Ro :

ð1aÞ

For convenience, the resistance R is used in this definition but it is evident that by using the resistivity q, the MR ratio will be the same. Although the quantity MR(H) is defined as a ratio and is usually expressed in percentages, it is very common to apply the term ‘‘magnetoresistance” only for this quantity. Expression (1a) represents a conservative definition for the magnetoresistance since, by taking into account that usually RH < Ro, the absolute value of the quantity MR(H) can only take values between 0% and 100%. It is also common to introduce an inflationary definition

MRinfl ðHÞ ¼ DR=RH ¼ ðRH  Ro Þ=RH ;

ð1bÞ

according to which, evidently, we can have also MR values above 100%. For not too high GMR effects as is often the case for ED multilayer films, the two definitions lead to very similar MR values (a GMR magnitude of 10% by the conservative definition converts to about 11% by the inflationary definition). Therefore, we shall not make a strict distinction between the two definitions when discussing the reported GMR values for ED multilayer films although we would suggest the use of the conservative definition, at least from an experimental point of view. For completeness, it should be noted, however, that the inflationary definition is especially preferred for treating the magnetoresistance of some magnetic perovskite compounds which arises as a consequence of a phase transformation under the application of an external magnetic field [159]. Due to the strong decrease of the resistivity in magnetic field with respect to the zero-field resistivity in these perovskites, the inflationary definition (1b) yields MR values much larger than the GMR values of metallic magnetic nanostructures and, therefore, the magnetoresistance observed in these perovskites was termed as colossal magnetoresistance (CMR). The physical mechanism of CMR is, however, quite different from that of the GMR and the CMR phenomenon will not be discussed further here. In the present paper, we shall use the conservative definition for the magnetoresistance. In an external magnetic field, the electrical resistivity changes to some extent for any metallic material. In NM metals, however, the change is very small (even in ordinary laboratory magnetic fields as about 10 kOe, their MR is the fraction of a percent only) and this effect is called as ordinary magnetoresistance (OMR) [160]. In the present review, we shall be concerned only with the magnetoresistance of homogeneous FM metals and their alloys (AMR) as well as of magnetic nanostructures composed of FM and NM metals (GMR). Such magnetic nanostructures include nanoscale multilayers consisting of alternating thin layers of FM and NM metals (e.g., Fe/Cr or Co/Cu) and granular magnetic alloys which consist of a matrix of a NM metal (e.g., Ag or Cu) in which nanoscale clusters of a FM metal (e.g., Co or Fe) are embedded randomly with their separation distance being also in the nanometer range. It is noted that when nanoscale magnetic layers are separated by typically 1 nm thick insulators such as, e.g., Al2O3 or MgO, spin-dependent tunneling of electrons through the thin insulating spacer can give rise to the tunneling magnetoresistance effect [161]. The same effect occurs if nanoscale magnetic particles are embedded in an insulating matrix. This topic is, however, beyond the scope of the present review since such nanostructures cannot yet be prepared by the electrodeposition technique and we shall restrict the discussion to fully metallic systems only. 2.2. Measurement of magnetoresistance The measurement of the electrical resistivity is usually carried out by using a four-point probe. In early studies of GMR in ED multilayers [26], the so-called van der Pauw geometry [162] was used. In

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this method, a fixed set of four point contacts arranged in a square form is placed on the sample (Fig. 1a). This method was originally elaborated for measuring the electrical resistivity of flat conductors (typically thin films) with laterally homogeneous transport properties. However, ED films/layers often exhibit lateral variation of thickness, microstructure and/or chemical composition, all these parameters having an influence on the electrical transport properties. Indeed, resistivity measurements by using narrow strips of the deposit and by arranging the four point contact on a line along the strip have revealed [82] a significant lateral inhomogeneity of the GMR in ED multilayer films under certain preparation conditions as to be discussed in Section 3.4. A schematic view of the resistivity measurement by the four-point method in the latter strip geometry is displayed in Fig. 1b. When measuring the magnetoresistance, the magnetic field can be parallel (longitudinal MR = LMR) or perpendicular (transverse MR = TMR) to the flow direction of the measuring current (the abbreviation TMR is not to be confused with tunneling magnetoresistance mentioned in Section 2.1, for which the same notation is commonly used). The significance of distinguishing between the LMR and TMR components of the magnetoresistance will become clear later when comparing the AMR and GMR effects. Since in multilayer films of interest the magnetization lies in the film plane, the external magnetic field is also oriented in the multilayer plane (field-in-plane or FIP configuration). As depicted in Fig. 2, for a multilayer the current can flow parallel (current-in-plane, CIP, configuration) or perpendicular (current-perpendicular-to-plane or CPP configuration) to the layer planes. In the study of multilayer films, the magnetoresistance can be conveniently measured in the CIP configuration only; due to the small total multilayer thickness (below about 1 lm), in the CPP configuration the resistivity to be measured is so small that its variation upon the application of an external magnetic field cannot be reliably established (at least at room temperature). Although multilayered nanowires are not discussed in the present review, it is noted at this point that for these latter nanostructures the measurement configuration always corresponds to the CPP case since the measuring current flows along the nanowire and, thus, it is perpendicular to the multilayer planes. For the multilayered nanowires, H can be applied either parallel or perpendicular to the wire, these two cases corresponding to the field-perpendicular-to-plane (FPP) and the FIP configuration, respectively. It should be pointed out that, when using the van der Pauw geometry for measuring the magnetoresistance, there is always some degree of admixture of the LMR and TMR components [26] because the local current flow direction within the sample inclines to the external magnetic field at various angles between 0° and 90°. By contrast, in the strip geometry it can be ensured that the current lines are strictly parallel or perpendicular to the external magnetic field and, therefore, no admixture of the LMR and TMR components can occur.

I

I

(a)

H transverse

I

I

I

I U

(b)

U

H longitudinal Fig. 1. Measurement of magnetoresistance (MR) for a multilayer film by the four-point-contact method in the van der Pauw geometry (a) and by the ‘‘four-point-in-line” method in the strip geometry (b), in both cases in the FIP/CIP configuration. The four contact points are indicated by the grey dots and stripes, respectively. In case the magnetic field H is oriented parallel to the measuring current I, the longitudinal magnetoresistance (LMR) is obtained whereas for a magnetic field H perpendicular to the measuring current, the transverse magnetoresistance (TMR).

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CIP configuration

CPP configuration

nanowire axis

film plane

multilayer film

multilayered nanowire

current flow Fig. 2. Configurations for measuring the magnetoresistance of multilayers (the grey and white stripes indicate the FM and NM layers, respectively). The measuring current can flow parallel (CIP configuration) or perpendicular (CPP configuration) to the layer planes. For multilayer films, the CIP geometry is used with the magnetic field H applied in the layer planes either parallel (LMR) or perpendicular (TMR) to the measuring current. For multilayered nanowires not discussed here, the magnetoresistance is measured in the CPP geometry whereas the magnetic field can be oriented along the wires axis or perpendicular to it.

2.3. AMR in bulk homogeneous ferromagnets It was observed already 150 years ago by Thomson (known later as Lord Kelvin) [163] that upon the application of a magnetic field the electrical resistivity of the FM metals Ni and Fe changes by a few tenths of a percent. He also established that if the external magnetic field is oriented parallel to the flow direction of the measuring current (I), i.e., H || I, then the resistivity increases (LMR > 0) and in the perpendicular orientation of current and field (H\I), the resistivity decreases (TMR < 0). It turned out from subsequent studies [160,164] that the situation is the same also for Co and for most FM alloys. The magnetic field dependence of the MR components are shown in Fig. 3 schematically for homogeneous FM metals by the (continuous and dashed) thick lines. The rapid resistivity variation at small fields and the opposite sign of the two components was explained by Bozorth [165,166] on the basis of the domain theory as follows. In the absence of an external magnetic field, a piece of a FM metal is always demagnetized to some extent (the magnetization of the individual magnetic domains is more

2 LMR

MR, ΔR/Ro (%)

0

AMR

-2

homogeneous bulk FM metal

TMR

-4 GMR -6 -8

LMR AMR

-10

TMR

FN/NM multilayer

-12

-8

-6

-4

-2 0 2 Magnetic field, H (kOe)

4

6

8

Fig. 3. Typical MR(H) curves shown schematically for a homogeneous FM metal (thick solid and dashed lines) and for a FM/NM multilayer (thin solid and dashed lines). The vertical lines with arrows at both ends denote the magnitudes of the quantities AMR (=LMR  TMR) and GMR.

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or less randomly oriented). On the other hand, beyond technical saturation (H > Hs), a single-domain state is reached (the magnetization in the sample is uniform and everywhere oriented along the magnetic field, in the longitudinal case M || I and in the transverse case M\I). With the latter notations, our intention was to emphasize that basically the relative orientation of M and I is only important and the role of the magnetic field is merely to ensure a homogeneous magnetic state and help in establishing a definite relative orientation of M and I. On the basis of the above discussion, we can now formulate the MR behavior of FM metals also by saying that the resistivity q is larger for the longitudinal configuration than for the transverse case: qL > qT. This relation finds its explanation in the spin-orbit coupling [160,164]. Namely, due to this interaction the charge distribution of d-electrons responsible for the observed atomic magnetic moment exhibits a non-spherical shape (this lowered symmetry manifest itself also in the magnetocrystalline anisotropy [167]). Since the major contribution to the electrical resistivity in transition metals comes from the Mott-type s–d scattering mechanism [168] to be discussed later, as a consequence of this non-spherical charge distribution of d-electrons, the predominantly s-type conduction electrons will experience a different scattering cross-section for the M || I and M\I cases and this leads to a difference in the resistivity of the L and T configurations (see Fig. 3). The difference between the LMR and TMR components in field range beyond technical saturation (H > Hs) is defined as the anisotropic magnetoresistance AMR  LMR  TMR. The magnitude of the AMR is typically a few percent only and it is positive for most FM metals and alloys [160,164] except for a few specific alloys for which its sign can be negative at certain concentrations (e.g., in the Ni–Cr system [164]). The steep resistivity change at small magnetic fields is due to the approach of the originally randomly oriented domain magnetizations towards the field direction as H ? Hs. The small, linear decrease of the resistivity in the saturation region is due to the reduction of the magnetic scattering as a consequence of the increasing magnetic ordering upon increasing magnetic field (this is the so-called paraprocess) [160,164]. To get the isotropic MR from the measured LMR and TMR data, we follow the usual procedure [160,164] as described below. For a truly randomly demagnetized specimen, we have for the longitudinal (L) and transverse (T) component of the resistivity that

qo  qav ffi ð1=3ÞqL þ ð2=3ÞqT :

ð2aÞ

Here, qo is the resistivity in zero external magnetic field and qav is the averaged (isotropic) resistivity component. Therefore, we can get the field dependence of the isotropic MR contribution by the formula

MRðHÞ ¼ ð1=3ÞLMRðHÞ þ ð2=3ÞTMRðHÞ;

ð2bÞ

where LMR(H) and TMR(H) are the corresponding experimental data. 2.4. GMR in magnetic nanostructures 2.4.1. GMR in FM/NM multilayers The thin solid and dashed lines in Fig. 3 demonstrate schematically the GMR phenomenon of metallic FM/NM multilayers by displaying the field dependence of the magnetoresistance in the CIP geometry for both the LMR and TMR components. Such multilayer films can also be visualized in a manner that thin NM metallic layers (e.g., Cr or Cu) are uniformly inserted into a thicker homogeneous metallic FM layer (e.g., Fe or Co) as a consequence of which there will be a drastic change of the magnetoresistance. Firstly, the saturation values of both the LMR and TMR components of the multilayer magnetoresistance will be usually much larger in comparison with the corresponding bulk values of the original FM metal. Secondly, the LMR and TMR components have identical sign in the multilayer (in the common case shown in Fig. 3, both are negative) in contrast to the typical case of LMR > 0 and TMR < 0 for most homogeneous FM metals and alloys. Similarly to a homogeneous bulk FM metal, the difference of the LMR and TMR components defines the AMR magnitude also for the multilayers. The markedly different MR(H) behavior of a homogeneous ferromagnet and a FM/NM multilayer can be explained with the help of Fig. 4. In case one can somehow achieve (possible ways for this will be discussed later) that for H = 0 the relative orientation of the magnetization of the adjacent magnetic layers is antiparallel (AP) or, at least, the adjacent layer magnetizations have non-negligible

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117

R RAP

RP H

0

M

HS

H

Fig. 4. Top panel: variation of the electrical resistance R with external magnetic field H for a FM/NM multilayer exhibiting GMR behavior. Medium panel: schematic indication of the magnetization orientations of two adjacent magnetic layers in the absence of a magnetic field (in the middle) and for magnetic fields beyond saturation. In the case of a parallel alignment (P) achieved in high magnetic fields, the resistance (RP) of the layered structure is smaller than the resistance RAP of the antiparallel alignment (AP). Bottom panel: the variation of the total magnetization of the layered structure with external magnetic field (Hs: saturation field).

antiparallel components, then the resistance RAP of this AP state will be larger than the resistance RP of the parallel (P) state. The latter state is realized by orienting all the layer magnetizations with a sufficiently strong external magnetic field along the same direction (see the bottom part of Fig. 4). The GMR effect is the difference between the resistances of the two magnetic states. For understanding the difference in the resistances of the P and AP states of FM/NM multilayers, we need to invoke models elaborated for explaining the electrical transport properties of FM metals. In transition metals, the carriers of electrical conduction are mainly the highly mobile, delocalized s-type valence electrons (for this behavior, they are also termed as conduction electrons) whereas d-electrons do not effectively contribute to the conduction (therefore, the latter electrons can be considered as strongly localized). Mott suggested already as early as in 1936 [168] that the scattering probability of conduction electrons depends not only on the scattering potential but also on the number of final states at the Fermi level into which the electrons can be scattered during the scattering process. On the other hand, the number of these states is determined by the electronic density of states at the Femi level, i.e., N(EF). Since the electrical resistivity is actually proportional to the total scattering probability, according to Mott’s suggestion the resistivity qs of the current carried by the s-type conduction electrons will be proportional to N(EF), i.e., qs / N(EF) = Ns(EF) + Nd(EF) where the subscripts s and d identify the s- and d-contributions to the density of states at the Fermi level. In transition metals the relation Nd(EF)  Ns(EF) is usually fulfilled, thus we finally obtain that

qs / Nd ðEF Þ:

ð3Þ

This is the Mott s-d scattering model for transition metals. Evidently, for metals with completely filled d-bands (e.g., Cu), we get qs / Ns(EF). With this model, Mott [168] could successfully explain why transition metals having unfilled d-bands are poorer conductors (exhibit higher resistance) than other metals having no d-electrons at the Fermi level.

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N( N(E )

EF

d

s

s

E

d

Fig. 5. Schematic electronic density of states N(E) for FM transition metals like Co and Ni with ‘‘strong itinerant ferromagnet” behavior [160] in the Stoner model of itinerant ferromagnetism. The electronic subbands corresponding to the two possible spin states (" and ;) are separately indicated for both the s- and d-electrons. The vertical line denotes the position of the Fermi level EF.

In the case of FM transition metals, the above picture needs further refinement. According to the Stoner model of itinerant ferromagnetism [160], the d" and d; electronic bands are shifted in energy with respect to each other (the d-band becomes split) as a consequence of the exchange interaction which itself is responsible for the onset of the FM state. This is shown schematically in Fig. 5 for the case of strong itinerant ferromagnetism (SIF) of the Stoner model (e.g., for Ni and Co metals) when the majority spin (d") subband is completely filled and, at the Fermi level, there are d; states only, i.e., we have Nd"(EF) = 0 and Nd;(EF) > 0. In such a case, from the viewpoint of the electrical transport properties, it is convenient to separate also the otherwise non-split s-band into two (identical) subbands (s" and s;). This enables one to write

qs" / ½Ns" ðEF Þ þ Nd" ðEF Þ ¼ Ns" ðEF Þ;

ð4aÞ

qs# / ½Ns# ðEF Þ þ Nd# ðEF Þ  Nd# ðEF Þ:

ð4bÞ

and

In deriving these expressions, use has already been made of the relation Nd;(EF)]  Ns;(EF). Since, furthermore, also the relation Nd;(EF)  Ns"(EF) holds, we get finally for the resistivity of transition metals and alloys the following expression:

qs" qs# :

ð5Þ

It was assumed here that the role of scattering processes leading to spin reversal is not significant, i.e., the so-called spin-mixing between the two conduction channels can be neglected which is often fulfilled indeed (e.g., at low temperatures where the phonon and magnon scattering can be neglected). It should also be noted that for not FM metals and alloys for which the d-band is not split, of course, the relation Nd"(EF) = Nd;(EF) holds and, thus, we get qs" = qs;. According to the above depicted picture, we can now imagine the electrical transport in FM metals and alloys in a manner that the current flows in two parallel channels (s" and s; channels) which are usually characterized with the very different resistivities q" and q;. Due to the latter feature of the " and ; channels, it is customary to speak about spin-dependent electron scattering processes in such cases. This is the so-called two-current model the elaboration of which was carried out mainly by Fert and Campbell [169] and this picture enabled the explanation of the GMR effect fairly soon after its discovery. On the basis of the Mott model and the two-current model, a simple resistor picture can provide us an instructive physical description of the GMR effect occurring in FM/NM multilayers [170]. The resistor model is visualized with the help of the trilayer structure depicted in Fig. 6 in which the two outer FM layers are separated by a NM spacer layer. In terms of the two-current model, the notations q" and

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FM NM FM

119

FM NM FM

ρ

ρ

ρ

ρ

ρ

ρ

ρ

ρ

Fig. 6. Resistivity contributions of the FM/NM/FM trilayer structure (at the top). In the case of the P magnetization configuration (""), the conduction electrons s" aligned parallel to the majority spin direction (by definition: d") can pass with low resistivity through both FM layers whereas the conduction electrons s; aligned along the minority spin direction (d;) experience a large resistivity in both FM layers (left upper panel). For the AP alignment (";), conduction electrons in both majority and minority spin channels experience the same resistivity when passing through the trilayer structure (right upper panel). At the bottom, the corresponding resistor network schemes are indicated for the two different magnetization alignments (P and AP configuration). In these network schemes, the smaller resistors represent the total background resistivity contribution of non-magnetic origin (i.e., not spin-dependent contributions) due to, e.g., lattice defects, impurities and phonons. Sketched after Ref. [170] with permission of Physics World and IOP Publishing.

q; are used for the resistivity of the majority and minority spin channels of the FM layers, respectively. These latter quantities will be defined with respect to an average resistivity q according to the expressions q" = q (1  b) and q; = q (1 + b) where b is an arbitrary number with the constraint 0 < |b| < 1 and the relation q = (q" + q;)/2 also holds. For the sake of simplicity, the resistivity contribution due to any irregularities of the FM/NM interfaces will be neglected. The whole trilayer structure is represented by four connected resistors as indicated at the bottom of Fig. 6. In the P configuration when the magnetizations of the two FM layers are aligned parallel, it can be figured out from the resistor network model that the resistivity qP will be given by the expression

qP ¼ 2q" q# =ðq" þ q# Þ ¼ qð1  b2 Þ;

ð6aÞ

whereas the resistivity qAP corresponding to the antiparallel configuration will be given by

qAP ¼ ðq" þ q# Þ=2 ¼ q:

ð6bÞ

It is easy to see that the relation qP < qAP always holds whatever was the choice of q and b and, thus, the resistivity variation with magnetic field as depicted in the top panel of Fig. 4 is indeed obtained. Several theoretical models of the GMR phenomenon have been elaborated for both the CIP and CPP geometries. The treatment of spin-polarized transport in metallic FM/NM multilayers is a very complicated task. The elaborated theoretical approaches cover a wide range of models: they span from the simplest semiclassical theories based on the Boltzmann equation through treatments involving quantummechanical phenomena to models taking into account the electronic band structure with the use of the density functional theory. It is not the purpose of this review to discuss these theoretical approaches, for more details the interested reader should consult, e.g., the book of O’Handley [160]. 2.4.2. Exchange coupling and oscillatory GMR in multilayers It was already noticed in the Introduction that the recent decades have witnessed a remarkable progress in thin film technologies. The development was especially pronounced in the field of epitaxial

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layer growth which enabled the preparation of materials at the precision of a single atomic layer. The MBE technology was originally developed for the semiconductor industry and from the late 1970s it could already be used to grow also nanoscale metallic thin films with a fairly defect-free structure on appropriate single-crystal substrates. This progress has paved the road for the subsequent preparation of nanoscale metallic multilayers consisting of alternating FM and NM layers, each having a thickness of a few monolayers only [7,171]. In the case of a proper epitaxial growth and a uniform layer thickness, a high coherence of the subsequent atomic layer planes can be realized along the thickness; therefore, high quality nanoscale multilayers are often termed also as superlattices. In 1986, surprising results were obtained from the studies of the magnetic properties of nanoscale layered structures. It turned out from the work of Majkrzak et al. [172] on Gd/Y multilayers and of Grünberg et al. [173] on Fe/Cr/Fe trilayers that at certain thicknesses of the Y layer and the Cr layer, respectively, an antiferromagnetic (AF) coupling develops between the adjacent magnetic layers. Such a coupling created by the indirect exchange interaction of conduction electrons via their substantial spin-polarisability has long been known to act between localized moments of magnetic impurities situated in a non-magnetic host matrix. The surprising fact was that a similar coupling can be mediated between layer magnetizations by the conduction electrons of a NM spacer layer [161]. Upon studying the magnetoresistance of Fe/Cr/Fe sandwiches with 12 nm thick Fe layers and 1 nm Cr layer, Grünberg and coworkers observed in 1988 [2] that at room temperature both the LMR and TMR components decreased by about 1.5% in an external magnetic field and this magnetoresistance was by an order of magnitude larger than the value measured for a 25 nm thick single Fe layer. This large MR could be explained as arising from the large resistivity in the AP aligned state of the two Fe layers at H = 0 due to the AF coupling of their magnetizations which coupling was already known from an earlier work of Grünberg’s group [173]. In some Fe/Cr multilayers, this group observed a magnetoresistance of 10% at T = 5 K [2]. By realizing the potential for sensor applications, Grünberg immediately patented the newly discovered phenomenon for magnetic field sensing devices [174]. In the same year, Fert and coworkers reported [3] on the observation of nearly 50% magnetoresistance in some MBE-grown Fe/Cr multilayers at 4.2 K with saturating magnetic fields of 20 kOe. This unusually large change of resistivity was termed by them as giant MR. It was an important further step when in 1990 Parkin and coworkers [175] reported that the GMR phenomenon can be observed also in Fe/Cr, Co/Cr and Co/Ru multilayers prepared by sputtering, a method much simpler and less expensive than MBE. In addition, they have also observed a regular oscillation of both the GMR magnitude and the saturation field as a function of the thickness of the non-magnetic spacer layer (Cr and Ru). More specifically, at the spacer thickness where the GMR magnitude had a maximum, the saturation field also exhibited a maximum. This indicated that at these specific spacer thicknesses a strong AF coupling is present whereas between the AF maxima, the coupling was ferromagnetic with low saturation field and due to the FM coupling, the GMR disappeared at these spacer layer thicknesses (if a measurable magnetoresistance was observed at all, it was due to the bulk AMR of the magnetic layer material). Therefore, the oscillating GMR is the consequence of an oscillating exchange coupling alternating between AF and FM character and it reflects the spacer layer thickness dependence of the interlayer exchange coupling [161] as we can see it schematically in Fig. 7 where also the corresponding layer magnetization orientations are indicated. As already noticed before, this interlayer coupling is of similar origin as the interaction of localized magnetic moments in a non-magnetic matrix although the asymptotic form of the dependence of the coupling strength on spacer layer thickness differs from the dependence of coupling on the distance between localized moments in a non-magnetic matrix [176]. The real breakthrough towards applications of the GMR effect occurred in 1991 when both Fert and coworkers [177] and Parkin and coworkers [178] reported the observation of nearly 50% GMR even at room temperature in sputtered Co/Cu multilayers. The results of Fert and coworkers [177] are reproduced in Fig. 8 where we can also observe the oscillating GMR behavior. 2.4.3. GMR in granular magnetic alloys In the early stages of GMR research on FM/NM multilayers, a large GMR was usually observed only in the presence of a strong AF coupling between the adjacent layer magnetizations which was mediated by the conduction electrons of the NM spacer metal. Therefore, it was not really clear at the

I. Bakonyi, L. Péter / Progress in Materials Science 55 (2010) 107–245

121

P

J(D)

D

P D

D AP D

Fig. 7. Oscillation of interlayer exchange coupling strength J(D) for a pair of FM layers separated by a NM metallic spacer layer of thickness D. In the absence of an external magnetic field, for J(D) > 0 (FM coupling) the layer magnetizations are parallel (P) aligned, for J(D) < 0 (AF coupling) the alignment is antiparallel (AP).

Fig. 8. Oscillation of the GMR saturation values (GMRs) with the thickness of the NM Cu spacer layer in sputtered Co/Cu multilayers at T = 4.2 K and 300 K. Reprinted from Ref. [177] with permission of Elsevier.

beginning whether the AF coupling is a necessary prerequisite for the occurrence of the GMR phenomenon. The observation of a large GMR effect in Cu(Co) granular magnetic alloys as reported simultaneously by the groups of Berkowitz [179] and Chien [180] in 1992 was of great importance for settling this issue. A granular magnetic alloy consists of a NM metallic matrix into which small particles of a FM metal (or alloy) are embedded. Granular alloys are usually produced by first preparing a solid solution of a FM metal in a NM metal matrix, e.g., in the form of a film or thin layer by using non-equilibrium deposition techniques such as evaporation, sputtering or electrodeposition; or in a more bulky form by rapidly quenching the melt of an appropriate alloy. If the FM metal is not miscible with the base NM metal, the solid solution is metastable. The application of a heat treatment results in the precipitation of the insoluble FM metal, which then forms small particles in the NM matrix. By varying the starting alloy composition as well as the time and temperature of the heat treatment, one can control the size and separation (density) of the precipitated FM particles. The magnetic properties of these nanoscale precipitates are governed by finite size effects as discussed below.

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For a FM material, below a critical size that depends on both macroscopic sample shape and material parameters like the FM exchange coupling strength between atomic moments and the magnetocrystalline anisotropy constants, a single-domain state will be the preferred configuration due to the energy associated with the creation of a magnetic domain wall [160]. Reducing further the particle size at a given temperature, the FM particle will exhibit a so-called superparamagnetic (SPM) behavior if the thermal energy kT becomes larger than the total anisotropy energy KV of the particle where K is the constant of the resultant anisotropy of any sources and V is the particle volume [160]. Namely, in such a case the thermal agitation energy permanently overcomes the energy barrier separating the stable positions of the particle’s magnetization the orientation of which will, therefore, randomly fluctuate and its temporal average will be zero. An assembly of non-interacting SPM particles will show a hysteresis-free magnetization curve (both remanence and coercivity are zero). The magnetization process for SPM particles is described by the classical theory of Langevin paramagnetism [160]. Accordingly, the magnetization curve M(H) follows the Langevin function L(x) = cth x  1/x where x = lH/kT with l constituting the average magnetic moment of a SPM particle. The magnitude of l is typically in the range 100–1000 lB where lB is the Bohr magneton. The saturation of SPM particles usually requires high magnetic fields, quite often several tens of kOe. Consider now an assembly of SPM particles embedded in a NM metallic matrix in a manner that due to a sufficient separation distance between the particles there is no magnetic interaction between their magnetic moments (see Fig. 9). A conduction electron just leaving a SPM entity is polarized by the magnetic moment of the particle. If this electron travels in the NM matrix and arrives at a neighboring SPM particle in a time shorter than its spin-memory time (i.e., it arrives with its original spin orientation), then at the second SPM particle it will definitely undergo a spin-dependent scattering event since the orientations of the two SPM moments are uncorrelated. This spin-dependent scattering gives rise to a GMR contribution just as in the case of FM/NM multilayers when electrons travel via the NM spacer between the two FM layers of different magnetization orientations. We shall now discuss the relation between the field dependence of the magnetization and the magnetoresistance for granular magnetic alloys containing non-interacting SPM particles only. As noticed above, for an assembly of non-interacting SPM particles, the magnetization is proportional to the Langevin function [160]:

MðHÞ / LðlH=kTÞ:

ð7aÞ

Gittleman et al. [181] considered for the first time quantitatively the contribution to the magnetoresistance MR(H) due to an electron moving from one SPM particle to another one as pictured in Fig. 9. They showed that the probability for the electron to be scattered depends on the degree of correlation

Fig. 9. NM metallic matrix containing non-interacting SPM particles (circles). The arrows attached to each circle represent the random thermal fluctuation of the SPM magnetic moment orientations. The arrow connecting two SPM particles indicates the path of a conduction electron travelling between the particles. Due to the rapid and uncorrelated thermal fluctuation of the SPM moments, an electron polarized by the magnetic moment of the first particle will definitely encounter a different magnetic moment orientation when arriving at the second SPM particle and, therefore, will undergo a spin-dependent scattering there. This gives, then, rise to a GMR effect.

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123

M

M ~ L(μH/kT) 0

2

MR

MR ~ -[L(μH/kT)]

0

H Fig. 10. Field dependence of the magnetization (M) and magnetoresistance (MR) for an ideal granular magnetic alloy consisting of non-interacting SPM particles embedded in a NM metallic matrix.

between the magnetic moment orientations of neighboring SPM particles averaged over all configurations. Thus, MR(H) is proportional to hl1 l2i where l1 and l2 are the magnetic moment vectors of the initial and final SPM particles that define the path of the scattered electron and h

i denotes a thermal average. For non-interacting SPM particles, the externally applied magnetic field is the only cause for correlations between l1 and l2. Therefore, it follows that:

MRðHÞ / hl1 l2 i ¼ hl1 nH ihl2 nH i ¼ ½MðHÞ2 ¼ ½LðlH=kTÞ2 ;

ð7bÞ

where nH is a unit vector in the direction of the magnetic field. The same form of the dependence of MR(H) on the magnetic field was derived also by Zhang [182] and Wiser [183] and the considerations for the derivation of Eq. (7b) were taken from the latter work. The field dependence of the magnetization and magnetoresistance for such an ideal granular magnetic alloy containing non-interacting SPM particles only is compared in Fig. 10. Since according to Eq. (7b) we can also write MR(H) / [M(H)]2, the experimental GMR data are often plotted against the square of the magnetization taken at the same magnetic field. We can now make a parallelism between the GMR in FM/NM multilayers and in ideal granular magnetic alloys. In both cases, the GMR effect arises due to a spin-dependent scattering event at the end of an electron path ‘‘magnetic region A ? non-magnetic region ? magnetic region B”, provided the orientations of the magnetic moment (magnetization) in regions A and B are different at the moment when the electron ‘‘visits” both regions (during the time the electron needs to travel between two such adjacent regions, the magnetic moments of the latter ones can be regarded as fixed). For FM/NM multilayers, the magnetic regions A and B correspond to two adjacent FM layers, whereas for ideal granular magnetic alloys, to two neighboring SPM particles. There are, however, also important differences in this respect between the two kinds of magnetic nanostructures. In a granular alloy, rapid thermal fluctuations of the SPM moments ensure that, apart from full saturation of the SPM moments by a strong external magnetic field, there is with some finite probability a spin-dependent scattering (i.e., a GMR contribution) whenever an electron travels between two SPM particles. The scattering probability is governed by Eq. (7b) and is completely suppressed in the magnetically saturated state only. On the other hand, in FM/NM multilayers for which the magnetization orientation of each magnetic layer is stable at a constant field value, some other mechanisms are required to ensure that the magnetic moments of two neighboring FM layers (at least over areas laterally larger than the electron mean free path in the NM layers) are not aligned

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parallel to each other. Such a mechanism can be, for example, the AF exchange coupling between the magnetizations of adjacent layers discussed in Section 2.4.2. In the presence of an AF coupling, in zero magnetic field the magnetizations of the FM layers are alternately antiparallel aligned and the probability of spin-dependent scattering for electron pathways ‘‘FM layer A ? non-magnetic layer ? FM layer B” is high, resulting in high resistivity. In an increasing external magnetic field, the layer magnetizations are gradually aligned parallel to each other (along the field direction) and the probability of spin-dependent scattering reduces, diminishing also the multilayer resistivity; finally, the MR is saturated when all the layer magnetizations are aligned along the external field. A schematic evolution of the magnetization and magnetoresistance of an AF-coupled FM/NM multilayer was shown in Fig. 4. For such multilayers, a quantitative description of M(H) and MR(H) cannot be given in a general form since the actual magnetization process which then governs the field dependence of the magnetoresistance is controlled by the interplay of the strengths of the actual interlayer coupling and the magnetic anisotropies present, with the range varying from dominating AF coupling through a 90°-coupling to a completely uncoupled state. However, it was established experimentally for a strongly-coupled Fe/Cr/ Fe trilayer [184] that the GMR magnitude depends on the relative orientation of the layer magnetizations, which are governed by the interlayer exchange coupling, magnetic anisotropy and external magnetic field. The dependence of GMR on the relative orientation of the adjacent layer magnetization has been explicitly studied also theoretically [185]. In any case, the observation of GMR in granular magnetic alloys demonstrated that the occurrence of a GMR effect in FM/NM multilayers is not connected to the presence of an AF coupling although such a coupling, if present, can ensure a high degree of antiparallel alignment and can result in a high GMR. However, since the saturation field of SPM particles is usually fairly high, a large magnetoresistance in granular alloys can be reached in correspondingly high magnetic fields only. This fact, on the other hand, leads to a low field-sensitivity of such materials in the application-relevant range of magnetic fields to be detected. Up to now, we have discussed the case of ideal granular magnetic alloys only in which all the magnetic regions exhibit SPM behavior and they are non-interacting. It was, however, observed by Hickey et al. [186] that in some melt-quenched granular alloys the field-dependence of the magnetoresistance could not be described by Eq. (7b). This could be explained [183,186] by assuming a distribution of FM particle sizes so that at a given temperature some of the particles are in the SPM regime (smaller particles), while some others are so large that their magnetization is blocked against thermal agitation (kT < KV) and, therefore, they fall in the FM regime. This situation is sketched in Fig. 11. It is easy to see that with the simultaneous presence of both FM and SPM regions, the GMR may contain in general the following three contributions [183,186]:

Fig. 11. A NM metallic matrix containing SPM particles (circles) and FM particles (ellipsoids). The arrows attached to each circle represent the random thermal fluctuation of the SPM magnetic moment orientations. The magnetic moment of the larger FM regions have a fixed orientation specified here by the shape anisotropy. The arrow connecting an SPM particle and a FM particle indicates the path of a conduction electron travelling between the two particles. Due to the rapid thermal fluctuation of the SPM moments, an electron polarized by the magnetic moment of the SPM particle will definitely encounter a different magnetic moment orientation when arriving at the FM particle and, therefore, will undergo a spin-dependent scattering there (a reverse path will give the same result). This results, finally, in a GMR effect.

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(i) GMRSPM–SPM term (ideal granular magnetic alloy) corresponding to a spin dependent scattering event for an electron path ‘‘SPM region 1 ? NM region ? SPM region 2”. (ii) GMRFM–FM (GMRFM) term (ideal FM/NM multilayer) corresponding to a spin dependent scattering event for an electron path ‘‘FM region 1 ? NM region ? FM region 2”. (iii) GMRFM–SPM = GMRSPM–FM (GMRSPM) term (mixed magnetic nanostructure) corresponding to a spin dependent scattering event for an electron path ‘‘FM region ? NM region ? SPM region” or ‘‘SPM region ? NM region ? FM region”. With arguments similar to those used at the derivation of Eq. (7b), i.e., the GMRSPM–SPM term in the above notation, one can show [183] that the relation

GMRSPM / LðlH=kTÞ;

ð8Þ

holds. The reduction in the power of the Langevin function comes from the fact that now a thermal average of correlation should be taken between a randomly fluctuating magnetic moment (SPM region) and a magnetic moment with fixed orientation (FM region). It turned out [186] that in rapidly quenched Cu(Co) granular magnetic alloys it is sufficient to take into account the GMRSPM and GMRSPM-SPM terms only for describing the field and temperature dependence of the GMR (except for low magnetic fields, i.e., below the full alignment of the FM entities. The ratio of the GMRSPM and GMRSPM–SPM terms was found to vary with temperature for an Au(Fe) granular magnetic alloy [187]: at high temperatures the majority of the magnetic regions is in an SPM state (here the GMRSPM–SPM term dominates) whereas at low temperatures the larger magnetic particles become blocked so they are in an FM state, leading to a dominance of the GMRSPM term here. We shall see in Section 2.4.4 that the same model can be successfully applied also to the MR results on many electrodeposited multilayers although in these latter systems the GMRFM and GMRSPM terms will survive only. This indicates that magnetic nanostructures can have a variety of size distribution of the FM and SPM magnetic regions. Besides the magnetic field and the measurement temperature, even their mutual spatial arrangement can influence which GMR contributions will be either important or rather negligible. Based on the experience with these above discussed two kinds of magnetic nanostructure, the intricate question arises whether materials processing parameters can be tailored in a manner, for either FM/NM multilayers or granular magnetic alloys, that the three possible GMR terms (GMRFM, GMRSPM-SPM and GMRSPM) will have comparable magnitude. 2.4.4. GMR in multilayers containing SPM regions As it was exposed in the Introduction, the dissimilarities of the GMR features of PD and ED multilayers deserve particular attention. One major difference is that an oscillatory GMR behavior as a function of the spacer layer thickness has not yet been conclusively demonstrated for ED multilayers. This issue will be discussed later only (in Section 5), partly based on the information to be presented in the current section. This information relates to an understanding of the origin of the other difference between PD and ED multilayers, namely the often observed high MR saturation fields in the latter systems which is, at the same time, also one reason for the still lower magnetic field sensitivity of GMR in ED multilayers. This problem could be best approached by analyzing the field dependence of the magnetoresistance. Namely, the magnetic field dependence of the magnetoresistance, MR(H), in ED multilayers was found to be in most cases very different from that of PD multilayers with clear AF coupling (see Section 2.4.1) or from that observed in ideal granular magnetic alloys (see Section 2.4.3). For many electrodeposited multilayers, the MR(H) curves were reported, especially at small spacer layer thicknesses, to saturate at magnetic fields well beyond 10 kOe only. In contrast, for PD multilayers MR saturation against the AF-coupling can usually be achieved, at least for Cu spacer layers, below 10 kOe [177,188–190]. A key point towards progress in this field was the observation [87] that in some ED Co–Ni–Cu/Cu multilayers the field dependence of the magnetoresistance could be well fitted by a Langevin function. It is important to note that the same function was derived for the field dependence of the GMRSPM term (see Eq. (8)) in the Wiser–Hickey model [183,186] for the case of non-ideal granular magnetic alloys as discussed in Section 2.4.3. This term arises from electron transitions between

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SPM

SPM FM

SPM FM

SPM

FM

FM

Fig. 12. Schematic representation of the gradual transition from ideal granular alloys (left) containing small SPM regions only to ideal magnetic/non-magnetic multilayers (right) in which the magnetic layers exhibit FM behavior only. Reprinted from Ref. [137] with permission of Wiley-VCH Verlag GmbH (Weinheim, Germany).

FM and SPM regions. It was, therefore, quite obvious to extend this model to the case of non-ideal FM/ NM multilayers [103] by assuming that the magnetic layers consist of not only FM regions but of SPM regions as well. This is quite common since it has been reported for several magnetic/non-magnetic multilayer systems prepared by using a variety of techniques such as sputtering [191–195], MBE [196,197] or electrodeposition [88,103,104,132] that the magnetization may contain both FM and SPM contributions. One can visualize the magnetic layers in such multilayers as consisting of both FM and SPM regions whereby the SPM regions are magnetically decoupled from the FM regions. The actual physical appearance and location of these SPM regions in the multilayer structure has not yet been revealed and this evidently requires further studies. Nevertheless, evidence from the magnetic data clearly supports their possible presence in magnetic/non-magnetic multilayers. A pictorial representation of such SPM regions can be in the form of small magnetic islands [104] within the magnetic layers which are surrounded by some magnetically diluted or completely NM regions, ensuring magnetic decoupling from the FM regions. Another hypothesis can be that the SPM regions constitute an interfacial zone between the FM and NM layers [115]. In any case, we can also imagine that there is a continuous transition from ideal multilayers with only FM layers to ideal granular magnetic alloys containing only non-interacting SPM regions, with a variety of transitional states in between. The latter ones often better correspond to real granular alloys and multilayers as depicted in Fig. 12 on the basis of Ref. [137]. The field dependence of the magnetization for a magnetic/non-magnetic multilayer containing both SPM and FM regions can be described as follows. The FM regions have a relatively low saturation field Hs (in ED Co-Cu/Cu multilayers, for example, it was found [103] that technical saturation of the magnetization is reached at about Hs = 1.7 kOe). Therefore, for H > Hs we can write

MðHÞ ¼ M s ðFMÞ þ M s ðSPMÞ LðlH=kTÞ;

ð9Þ

where Ms(FM) and Ms(SPM) are the saturation components of the FM and SPM contributions, respectively, to the total saturation magnetization of the multilayer and l is the average SPM moment. According to the Wiser–Hickey model [183,186], the GMR in such a multilayer can be assumed to arise from spin-dependent scattering of electrons which travel through either (i) the non-magnetic spacer between two FM regions (GMRFM = GMRFM-FM) or (ii) between a FM region and a SPM region (GMRSPM = GMRSPM-FM = GMRFM-SPM) whichever region is the first or second along the electron path or (iii) between two SPM regions (GMRSPM–SPM). It was found for ED Co–Cu/Cu multilayers [103] for which the magnetization M(H) followed Eq. (9) that for H > Hs, the field dependence of the magnetoresistance MR(H) could be described by the Langevin function L(lH/kT):

MRðHÞ ¼ MRFM þ GMRSPM LðlH=kTÞ;

ð10Þ

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0

0

-2

-1

MR (%)

-4

MR (%)

LMR, experimental data MRFM(H) MRSPM(H) + MRFM(H=∞ )

-6 -8 -10

-9

-6

-3

0

H (kOe)

3

-3 -4

LMR, experimental data MRFM(H) MRSPM(H) + MRFM(H=∞ )

-12

-2

-5 6

9

-9

-6

-3

0

3

6

9

H (kOe)

Fig. 13. Decomposition of the magnetoresistance (MR) for the case when, after saturating the GMRFM contribution (H > Hs), the only prevailing MR contribution is the GMRSPM term. Measured longitudinal MR(H) data (symbols h) for two ED Co(3.4 nm)/ Cu(2.2 nm) multilayers and the decomposed FM (solid line) and SPM (dotted line) contributions to the observed magnetoresistance. The dominance of the SPM term (left) or that of the FM term (right) is clearly demonstrated.

whereby MRFM = AMR + GMRFM is a constant for H > Hs. Eq. (10) also implies that for these multilayers, from the above three possible GMR mechanisms (GMRFM, GMRSPM and GMRSPM-SPM), the first two terms only survive (GMRFM and GMRSPM) and the third one (GMRSPM–SPM) is negligible. If the FM regions are spatially sufficiently extended (at least in two dimensions such as in multilayers), an electron may undergo two subsequent spin-dependent scattering events in the same FM region. This gives rise to an anisotropic magnetoresistance (AMR) effect [160,164]. Beyond saturation (H > Hs), the GMRFM and the AMR terms are saturated and, hence, their contributions are constant, apart from a small linear term which is also present in bulk ferromagnets but which is negligible compared to the other effects in multilayer samples. Therefore, the contribution of the GMRFM and AMR terms cannot be separated from each other at H > Hs and their sum was denoted as a single MRFM term in Eq. (10). The relative weight of the MRFM and the GMRSPM terms as well as that of the two contributions in MRFM (AMR and GMRFM) do not simply depend on the volume fractions of the two kinds of magnetic region. This is because these weights are also determined by the scattering probabilities in the different regions and by geometric factors, i.e., the shape, relative position, spatial distribution and the morphology of each region. In ED Co–Cu/Cu multilayers, for example, the AMR term is typically 1% which corresponds well to the AMR of a Co-rich bulk Co–Cu alloy [103] whereas the GMRFM contribution can be as high as 10% [99,121]. The decomposition of the FM and SPM magnetoresistance terms on the basis of Eq. (10) is illustrated in Fig. 13 for two ED Co–Cu/Cu multilayers in which either the SPM (a) or the FM (b) contribution dominates the observed GMR. The decomposition procedure [103] starts with the fitting of the experimental MR(H) data in the high-field region (H > Hs) with a Langevin function L(lH/kT). This fitted Langevin function is then subtracted from the experimental data and this yields the MRFM term of Eq. (10). The procedure is carried out for both the longitudinal and transverse MR data, yielding usually very similar fit parameters including the average SPM magnetic moment l. If the difference between LMR and TMR experimental data (AMR) is small as is often the case, the MRFM term can be identified as the GMRFM contribution. The presence of a GMRSPM component well explains the observed large MR saturation field in many ED multilayers. It is evident from the foregoing discussion that the low-field sensitivity of a GMR multilayer remains small if it contains a large fraction of SPM regions since then the total magnetoresistance can reach saturation in fairly large magnetic fields only. It should be noted that, in addition to ED Co–Cu/Cu [103,115,121,128,132,138,139,141] and Ni-Cu/ Cu [104] multilayers, the presence of a GMRSPM term could also be identified in MBE-grown Co/Cu [198,199] and sputtered Co/Cu [200], Fe/In [201] and Ni83Fe17/Cu [192] multilayers. Therefore, the above decomposition method can be helpful in analyzing the strongly non-saturating behavior of GMR curves observed in any multilayer system.

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The oscillatory GMR arising from an oscillatory exchange coupling of the layer magnetizations can be associated with the GMRFM term only. Therefore, when searching for an oscillatory GMR in electrodeposited multilayers, evidently the GMRFM term should be separated from the total measured GMR or the GMRSPM term should be suppressed as much as possible by the preparation conditions. Although we do not yet have direct experimental evidence for the actual physical form and location of the SPM regions, we can already identify some factors controlling their appearance which can be summarized as follows. (i) It has been shown [99] that an improper choice of the deposition potential of the NM metal (typically Cu) can result in a significant increase of the GMRSPM fraction. The electrochemical processes responsible for this will be discussed in more detail in Section 3.6. (ii) In the case of immiscible elements such as Co and Cu [202], the tendency for the formation of SPM regions was found [121] to be larger with increasing amount of Cu in the magnetic layer, the latter always being a Co-rich Co–Cu alloy for multilayers of these elements when electrodeposited from a single bath. (iii) It has been demonstrated recently [200] that a large GMRSPM term may develop also for sputtered Co/Cu multilayers deposited on rough Ta substrates and a model was also provided by these authors to visualize this mechanism. Our recent results [128,139] corroborate this finding in that we have found that for Co–Cu/Cu multilayers electrodeposited on a rough substrate (mechanically polished Ti foil) the GMRSPM fraction was definitely larger than for multilayers deposited under identical conditions on a smooth substrate (Si wafer with evaporated Ta or Cr adhesive and Cu seed layers). In Section 5, a more detailed discussion supported with reported MR data will be presented to illustrate the ways how SPM regions can form during the preparation of ED magnetic/non-magnetic multilayers.

3. Summary of FM/NM multilayer film preparation by electrodeposition 3.1. Electrochemical background Electrochemistry [203–208] is a field of physical chemistry. It deals with condensed systems in which at least one phase is composed of, partly or entirely, ions that can be displaced upon applying an external electric field. While electrochemistry deals with all phenomena related to ionic conductivity as well, an electrochemical reaction always involves a charge transfer. In a heterogeneous electrochemical reaction, the charge transfer occurs at a phase boundary. For electrodeposition, one needs a solid metal or a semiconductor substrate on the surface of which the charge transfer takes place and where the deposit is accumulated. The literature of electrodeposition [209–213] makes use of both the scientific background of electrochemistry and the technical knowledge accumulated in experiments; hence, it is significantly empirical. The scientific background of electrochemical deposition is shortly summarized in this part, while other technical issues and empirical findings that are important for multilayer deposition are dealt with in other parts of Section 3. A commonly used term in electrochemistry is the notion of the electrode. According to the scientific approach, the electrode is the boundary layer between an ionic conductor and an electron conductor because this boundary exhibits all features that make the charge transfer between these phases possible. In contrast to this scientific approach, the metallic (or semiconductor) phase that is immersed in an electrolyte solution is also often termed as an electrode in the lab slang. Interested readers are advised to consult Ref. [214], which is the most comprehensive and up-to-date collection of definitions related to electrochemistry. It is of topmost importance to emphasize that the potential difference between the internal part of an electrolyte solution and the metal immersed in this electrolyte cannot be measured. For a measurement, one has to immerse another electron conductor phase to the electrolyte; hence, the measurement can be carried out by using two electrodes only. If we fix the electrode referring to which all

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potentials are measured, we can create a potential scale. One can see now that the electrochemical potential scale exhibits all features that any other potential scale does, namely: (i) the potential value itself depends on the reference point chosen; (ii) not the potential values themselves but the potential differences are indicative of the driving force of the charge transfer. The standpoint for referencing the potential in electrochemical systems is the standard hydrogen electrode (SHE). Although this electrode is seldom used in electrodeposition experiments for various practical reasons, other commonly used reference systems are all validated as compared to SHE. The reference electrodes used in studies related to multilayer electrodeposition will be discussed in Section 3.4. Electrodes can be classified according to various categories. If the nature and number of phases involved in the electrode reaction are taken, one can distinguish electrodes on the first kind, second kind and redox electrodes. Electrodes composed of a metal and the solution of its salt(s) belong to the family of the electrode of first kind, whereas the electrodes of second kind are good reference electrodes due to their excellent potential stability. If the direction of the current flowing through the electrode is important, one can speak about cathode and anode. On the cathode, the current is negative, which corresponds to the flow of the positive electricity from the electrolyte toward the electron conductor, while the opposite holds for the anode (regardless of the exact nature of the charge carrying species). If the basis of categorization is the number of electrode reactions, one can distinguish simple electrodes with a single electrode reaction or mixed electrodes with multiple electrode reactions. If an alloy is deposited, one always has to encounter a mixed electrode. Finally, if the role of an electrode in the cell is regarded, the general terms of working electrode, reference electrode and counterelectrode (or, alternatively, auxiliary electrode) are used. The working electrode is simply the electrode at which the reaction taking place is of the researcher’s interest. The reference electrode serves as the standpoint for the potential scale during the experiments and it carries no current. The counterelectrode serves for passing the charge through the cell. The dependence of the equilibrium potential of an electrode with a single electrode reaction on the concentration of the electroactive species is called the Nernst equation and it can be evaluated from thermodynamics. For a metal in equilibrium with its ion in the solution, it can be written as follows:

E ¼ E0 þ

RT cðMezþ Þ ; ln zF c0

ð11Þ

where E is the electrode potential, E° is the standard electrode potential, F is the Faraday constant (F = 95,486 C/mol), z is the charge of the metal ion in the electron charge unit, c(Mez+) is the metal ion concentration in the solution and c° = 1 mold m3 is the standard concentration. The latter term makes the argument in Eq. (11) dimensionless and also defines the E° values (for different choices of c°, different actual E° values would be obtained with the same order of the standard potentials for various electrode reactions). It can be seen that one order of magnitude concentration change of the metal ions in the electrolyte results in about 60 mV/z change in the equilibrium potential at room temperature; hence, the deposition sequence as a function of potential cannot be changed by simply modifying the metal ion concentrations in the electrolyte. Although the equilibrium potentials yield a guideline to determine in which direction the electrode reactions can go on, electrodeposition always takes place at a non-zero rate; hence, the appropriate kinetic equations have to be known. The steady-state current density – potential relationship is known as the polarization curve in the electrochemistry literature. The electrode potential can always be regarded as the driving force of the electrochemical processes, although the current can also be a regulated parameter from the viewpoint of the electronic control. For a one-step reversible electrochemical reaction, O + e = R (O and R are the oxidized and reduced form of the same species), the dependence of the current density i on the electrode potential can be written as

i ¼ kA cR exp





azFE RT

  ð1  aÞzFE  kC cO exp  ; RT

ð12Þ

where k is the rate constant of the anode or cathode reaction as shown by the lower index, c is the concentration of the appropriate reactant and a is the transfer parameter. The first term in Eq. (11) accounts for the anodic process, while the second term is related to the cathodic process of the same equilibrium reaction. Eq. (12) is known as the Butler–Volmer relationship [215].

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Electrochemical reactions are usually composed of several consecutive steps, and the rate of the electrode reaction is determined by the most hindered step. In case of electrodeposition, the following general scheme presents the consecutive processes: (i) transport of the electroactive species in the electrolyte solution; (ii) preliminary chemical reaction of the electroactive species; (iii) adsorption of the active species at the electrode; (iv) charge transfer step(s); (v) nucleation of the deposit at the electrode; (vi) diffusion of the adatom at the electrode surface towards the final lattice position. It is possible that some steps are missing or more steps are necessary to describe the deposition mechanism; therefore, all deposition processes are more or less unique. Amongst the steps listed above, the formation of the new solid phase gains an outstanding interest [216] in the case of electrodeposition. 3.2. Single-bath technique with pulse plating In the single-bath techniques, one has to fix all parameters related to the electrolyte composition before the experiment, and the modulated structure is to be obtained with the modulation of the deposition parameters electric in nature, i.e., current or potential. Regardless of the actual pulse sequence used, the deposition of a new layer always begins when a new pulse is started. (In a few cases, a pause is also applied between the pulses. The impact of such deposition mode is difficult to explain and has no theoretical importance.) While the modulation of the electric parameters is very simple for computer-controlled devices even at the millisecond time frame, the single-bath technique also has drawbacks imposed by the limitations of the predetermined electrolyte composition. These aspects of the technique are explained in the following sections. 3.2.1. Bath composition and its relation to the deposit composition One prerequisite of the achievement of the FM/NM layer structure with GMR is that besides the non-ferromagnetic behavior of the NM layer, this layer must be void of any ferromagnetic impurity in order to exclude any undesirable electron scattering in the NM layer. To fulfill this criterion, the spacer metal has to be a more noble element than the magnetic metals, and the nature of codeposition of the NM metal with all FM metals in the same bath has to be the so-called normal one [209]. Simply speaking, the normal codeposition means that the more noble NM metal can be deposited alone under the conditions where the FM metals would not deposit on their parent metal either. The diagnostic criteria of the normal codeposition can be approached in various ways, depending on whether one puts emphasis on the electrochemical control parameters or on atomic interactions: (i) If the deposition is carried out in the current control mode (galvanostatic conditions, G), the normal codeposition means that the NM metal is deposited alone as long as the mass transport of its ions from the bulk electrolyte to the cathode surface provides the flux required by the current maintained, regardless of the concentration of any other metal salt at the cathode. At the limiting current density, all NM metal ions arriving to the cathode surface are immediately discharged, and any further current increase can be achieved only by the occurrence of another electrochemical reaction, too. (About the concept of the limiting current, see Section 3.2.2). (ii) If the deposition is carried out in the potential control mode (potentiostatic conditions, P), the normal codeposition means that there is a potential range where the NM metal is deposited alone. This potential range is more positive than the deposition potential of the other metals. The magnetic metal(s) can be deposited at sufficiently negative potential only, and the onset of their deposition potential is identical to or more negative than the deposition potential of the same metal ions on their parent metal. The larger the energy barrier of the nucleation of the less noble metal on the more noble one, the larger is the shift of its deposition potential on the NM metal with respect to the deposition on the parent metal. (iii) If the normal codeposition is described in terms of the strength of the atomic interactions, it can be said that NM-FM interaction cannot be stronger than the NM-NM interaction, otherwise the alloying by the incorporation of the FM metal into the NM one can occur at much more positive potentials than the deposition potential of the FM metal ions on its parent metal.

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1.0

anomalous codeposition 0.8

yLN

0.6

0.4

normal codeposition

0.2

0.0 0.0

0.2

0.4

0.6

0.8

1.0

cLN/(cLN+cMN) Fig. 14. Schematic representation of the deposit composition as a function of the solution composition for two basic codeposition modes. Symbols: c is the ion concentration in the electrolyte and y is the molar fraction in the deposit. The denotation LN and MN in the lower index refer to the less noble and more noble metals, respectively. The arrows indicate how the functions change when either the current density decreases or the bath agitation becomes more intense.

The codeposition mode of two elements is often described with a composition diagram (see Fig. 14). For the construction of such diagrams, all technical parameters have to be fixed (temperature, hydrodynamic conditions, total metal ion concentration, current density); hence, the impact of the change in the ion concentration ratio in the electrolyte on the deposit composition can be obtained in the G mode. For normal codeposition, the alloy formation starts when the flux of the NM metal toward the cathode cannot account for all current passed. Of course, in this case the resulting deposition potential (which is not regulated here) must achieve the deposition potential of the FM metal. For most of the possible pairs of NM and FM metals for which GMR is expected, the codeposition mode is normal. NM metals like Cu, Ag, Au, Pd and Pt all can be deposited without the incorporation of Ni, Co or Fe. Hence, for the overwhelming majority of FM/NM metal pairs of practical importance it is automatically fulfilled that the NM metal can be deposited without a magnetic impurity within the NM metal. In contrast to the NM metal, the FM metal is always alloyed since the deposition of the NM metal cannot be suspended for the deposition time of the FM metal. The composition of the FM metal is directly determined by the rate of discharge of the ions of the bath components. In order to provide the deposition of an alloy with a sufficiently high Curie temperature, the FM layer should contain a high enough percentage of magnetic metal. This is achieved by choosing a large concentration for the ion of the FM metal in the electrolyte. The solubility of the salt of magnetic metals in aqueous electrolyte is seldom higher than 2 mol/l, and this limiting concentration also depends on the temperature and the concentration of other bath components. By increasing the concentration ratio of the FM and NM metal ions in the electrolyte, the deposit obtained under the same circumstances can contain a larger and larger percentage of the FM metal. The concentration ratio is usually set to between 7 and 200, and the FM metal content of the resulting alloy is typically between 80 and 99 at.%. Although the decrease in the concentration of the NM compound would also lead to the enrichment of the FM metal in the magnetic layer under identical deposition conditions, there are practical limitations on the concentration ratio of the FM and NM salts in the electrolyte:

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(i) If the concentration of the NM metal is reduced, the deposition time for a given layer thickness increases. Since the transport of the NM metal salt under the conditions of the NM layer deposition can be described with Fick’s first law, the deposition time is inversely proportional to the concentration. The necessity of obtaining a layer within a reasonable time frame sets the lower limit of the NM metal salt in the millimol/liter range. (ii) By decreasing the concentration of the NM metal, a larger volume of electrolyte solution contains the desired amount of the NM metal only. The minimum electrolyte volume V can be calculated with the help of the following formula: NdAq/M = cV where N is the number of bilayers, d is the thickness of the NM layer, A is the sample surface area, q is the density of the NM metal, M is its molar weight, and c is the electrolyte concentration. Moreover, not only does one need a large enough solution volume to contain the desired amount of material but the electrolyte composition should be practically unchanged during the deposition procedure in order to maintain identical experimental conditions throughout the sample preparation. This condition is fulfilled if the change in the salt concentration is not higher than a few percent. Hence, if the total deposit thickness is in the micrometer range, the surface area is in the cm2 range and an inert auxiliary electrode is used, the practical lower limit for the NM metal salt is again in the millimol/liter range, otherwise the electrolyte volume should be enlarged dramatically to provide unchanged conditions throughout the sample preparation procedure. The problem is similar if the auxiliary electrode is a sacrificial anode made of the NM metal. In this case, the electrolyte solution is not depleted during the experiment but enriched with respect to the NM metal salts since the deposition of the FM metal is also accompanied by the dissolution of the anode made of the NM metal. If the NM metal concentration is too low at the beginning, the dissolution of the sacrificial anode may lead to a substantial change in the solution composition. (iii) It is difficult to completely avoid all contaminations during the solution preparation and the deposition. For instance, the dissolved oxygen is always present in the electrolyte. At the potentials where the metals to be deposited are discharged, oxygen is also reactive and can be reduced at a measurable rate. The oxygen reduction can result in various products. Hydrogen peroxide is a strong oxidizer that may react with some components of the electrolyte in the bulk. The final product of oxygen reduction, i.e. hydroxide ions, can modify the pH in the vicinity of the cathode, hence changing the real deposition condition substantially. Since oxygen is hardly ever expelled from the solution prior to the deposition of the multilayer samples, its reduction should not exhibit a significant current contribution during the deposition of the NM metal. For this reason, it is advised that the NM metal salts are applied at least in a concentration of a few millimol/liter. If the NM layer in the multilayer is composed of an alloy of two NM elements, the impact of the electrolyte composition on the deposit composition is almost the same. In this case, both NM elements have to be codeposited with the FM element(s) in the normal way, and the concentration ratio of the FM metal salt and the sum of the NM metals salts has to be quite high. However, there is no restriction for the mutual deposition preferences of the two NM metals, and their codeposition features and their concentration ratio will determine the final composition of the NM layer. While the NM metal in the FM layer is often just a trace element, the FM layer itself may have several major magnetic components. The mutual deposition preference of the iron-group metals can be described with the so-called anomalous codeposition [209]. The name of the codeposition mode comes from the fact that the deposition preference is just the opposite that is expected from the formal indicator of the nobility (i.e., the standard potential) of the metals. While the order of standard potentials indicates the ease of deposition of a metal ion on the parent metal, this deposition order seems to be violated when some metals are codeposited. For instance, Co-rich deposits can be obtained from Ni-rich Ni–Co baths because of the preferred deposition of Co besides Ni. Similarly, baths developed for obtaining electrodeposited Fe–Co–Ni alloys contain the precursor compounds of the three metals in increasing concentrations as the standard potential of the metals increase. The typical composition curve for a metal pair exhibiting anomalous codeposition is also depicted in Fig. 14.

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There are various kinetic theories for the reason of the anomalous codeposition [217–221]. These theories mostly account for the interaction of the intermediates and the suppression of the deposition rate of the more noble metal in the presence of the ions of the less noble metal. It is worthwhile of mentioning that the mutual deposition preferences are mostly retained when more than two metals are being codeposited. For instance, in case of the deposition of Fe–Co–Ni alloys the deposition of Fe is preferred as opposed to the deposition of either Co or Ni, but the deposition of Co is preferred against the deposition of Ni only. Cu can be deposited in the normal way in the presence of either Co or Ni, and the deposition of Co is preferred against the deposition of Ni in the presence of Cu as well, as if Cu were not even present. The retention of deposition preferences facilitates the design of the bath composition on the basis of the behavior of metal pairs. 3.2.2. Deposition pulse modes, pulse combinations For understanding the advantages and drawbacks of the various deposition modes, one has to be aware of the polarization behavior of a typical bath used for depositing FM/NM multilayers. Fig. 15 shows a typical polarization curve for an electrolyte used to deposit FM/NM multilayers. It can be seen from Fig. 15 that a small cathodic polarization results in the deposition of the pure NM component. At the deposition potential of the NM element, the cathodic current starts to increase with potential in an exponential manner, according to the Butler–Volmer relationship (Eq. (12)). Since the transport of the electroactive ions in the electrolyte is not a function of the electrode potential and hence it cannot increase to the same extent as the driving force of the electrode process, the deposition of the NM element soon becomes diffusion limited. In this case, the vicinity of the cathode is so much depleted with respect to the ions of the NM element that its concentration extrapolated to the electrode surface is zero. In other words, the residence time of any intermediate of the NM ions adsorbed at the cathode tends to approach zero. The current density in the diffusion-limited case is determined by the slope of the concentration profile near the cathode via Fick’s first law (J = Doc/ox  Dc/dN, where J is the flux of the component, D is its diffusion coefficient, c is the concentration of the diffusing component and dN is the so-called Nernstian diffusion-layer thickness). When the potential is negative enough to achieve the deposition potential of the less noble FM element, its deposition may also start. The actual deposition potential of the less noble element can be more negative on the NM metal than on its parent metal because of an excess nucleation barrier. Once the deposition of the FM metal also takes place, the polarization curve is the sum of the two deposition

Fig. 15. Schematic presentation of the partial polarization curves of the more noble and less noble metals in relation to the overall polarization curve. Sections above the zero line: dissolution (Diss.), below the zero line: deposition (Dep.) Left: Dashed curve: more noble metal (MN), dotted curve: less noble metal (LN), both for concentrated electrolytes. The thin solid line shows the polarization curve of the MN metal if the ion concentration is low enough to achieve the diffusion-limited regime (DL). The thick solid line refers to a typical polarization curve for a single bath used for multilayer deposition (Ni–Cu, Co–Cu, Co–Ag etc.). Right: The arrows show how the molar fraction of the corresponding metals in the deposit can be estimated from the steadystate polarization curve: yLN = (iLN/zLN)/(iMN/zMN+iLN/zLN).

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processes, but only the deposition rate of the FM metal will be a function of the electrode potential. The latter process is usually not a diffusion-limited one due to the large concentration of the ions of the FM metal. Besides the processes of metal deposition, there are usually other unwanted side processes as well. The most typical one is the evolution of hydrogen, which can be described in acidic media with the following equation: H+ + e = ½ H2. The various aspects of this process will be discussed in detail in Section 3.5. What is important here for the elucidation of the polarization curves is that the hydrogen evolution modifies the hydrodynamic conditions around the cathode by stirring the electrolyte. This can increase the deposition rate of the NM element and prevents the occurrence of the mass transport limited deposition rate for the FM metal. Fig. 16 shows schematic polarization curves and illustrates the available range of the non-regulated parameter for both the G and the P modes (see Figs. 16a and b, respectively). We shall now discuss the deposition of both the more noble and the less noble metal with each control modes. If the NM metal is deposited in the G mode, the maximum current that can be applied is the diffusion-limited current, otherwise the less noble magnetic metal will contaminate the NM metal, which is detrimental for the GMR effect. However, the actual diffusion-limited current is difficult to estimate accurately even in the case when it is determined in a preliminary experiment with identical conditions. Minor changes in temperature, hydrodynamic conditions or ion concentration can all lead to a modification of the diffusion-limited current. If the deposition rate of the NM metal in the G mode is set to a value well below the diffusion-limited rate, the concentration of the NM metal ions close to the cathode will not be zero. This means that there is an excess metal ion concentration not fully exploited for the deposition of the NM layer. The ions not discharged by the current passed can take part in other processes not generating current. If the previously deposited FM metal is not yet fully covered with the NM metal, the ions of the more noble metal and the atoms of the less noble metal can react with each other, leading to a loss of a part of the FM layer and to an unwanted thickening of the NM layer. This is called the exchange reaction. For a pair like Co and Cu, this can be formulated as Cu2+ + Co = Cu + Co2+. In other words, the electrode potential can find its value in a regime where the dissolution of the less noble and the deposition of the more noble metals are favored, hence resulting in the above mentioned exchange reaction. If we consider the deposition of the NM metal simply from the viewpoint of the stability of the electrode potential, it can be said that the cathode potential is very uncertain in the G mode (see Fig. 16a). Therefore, it can be seen that the control of the deposition of the NM metal in the G mode is very difficult, if not impossible. The actual and nominal layer thicknesses and sample compositions are usually very different if the NM metal is deposited in the G mode.

Fig. 16. Relationship between the regulated and non-regulated electric parameter in various control modes: (a) current control (G mode), (b) potential control (P mode). The deposition of the FM and NM metals are indicated with different shades of color.

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In contrast, the deposition of the FM metal in the G mode is not really problematic. The ratio of the overall current to the diffusion-limited current of the NM metal can be used for a rough estimation of the composition of the FM layer (Fig. 16b). The estimate can be given more accurately if the current efficiency for the deposition of the FM layer is known. After the discussion of the G mode, we analyze now how the potential regulation (P mode) works for the deposition of each layer. If the NM layer is deposited in the P mode, the current will automatically set to the diffusion-limited current in the appropriate potential range, and the contamination of the NM layer with any FM element can be excluded. Some care has to be taken so that the potential is negative enough to avoid the dissolution of the previously deposited FM layer. The methods to optimize the deposition potential of the NM layer are discussed in Section 3.6. Figs. 15 and 16 show steady-state polarization curves, but the deposition of the multilayers requires pulse plating. In this case, the pulse change leads to the situation that the FM metal needs to be deposited onto the less noble FM metal. If the deposition potential of the NM metal makes it possible that the FM metal dissolves, this process will surely happen since the full coverage of the surface with the NM metal cannot be achieved instantaneously. In the P mode, the loss of the FM metal may also happen (similarly to the G mode), but the description of the phenomenon is different. In the P mode, the current is not fixed, but one has to take into account as many reactions as many may take place. Hence, we can speak about the independent deposition of the NM metal and the dissolution of the FM metal. Of course, the total current is taken now as the sum of the currents corresponding to the two processes. It is very important to understand this consequence of the change in the control parameter on the nomenclature. While the reaction of the more noble ions on the FM metal surface is taken into account via the exchange reaction in the G mode, one can only speak about the dissolution of the FM metal in the P mode. If the FM layer is deposited in the P mode, the situation is almost the same as in the G mode. The important difference is that one has to take into account that the deposition of the FM metal usually takes place with a short pulse during which steady-state deposition conditions cannot be achieved. This is why transient phenomena are of importance. If the potential is suddenly increased to the desired deposition potential of the NM metal, the electrolyte around the cathode becomes depleted, and the ratio of the deposition rate of the FM and NM metals will change during the pulse. Hence, the composition of the FM layer changes during the pulse. This problem does not occur if the FM metal is deposited in the G mode. We can summarize now the deposition opportunities as the arbitrary permutations of pulses of different nature. In the G/G mode, both the FM and NM components are deposited by the current control. (Hereinafter, the denotation of the deposition mode of the FM layer stands always first and is followed by that of the NM layer.) While the deposition of the FM layer should not be a problem, the deposition conditions of the NM layer cannot be adjusted precisely enough to keep the FM layer fully intact. Although the G/G mode is technically the simplest one since one needs only a current generator with a timer, it is worth of investing to use a more advanced setup. In the P/P mode, the deposition of both layers can be adjusted fairly well. Besides the conventional pulse methods with identical control parameters for both layers, a mixed deposition mode is also possible. In the G/P mode, the deposition conditions for both pulses can be adjusted the best way in the authors’ opinion. The application of the G/P mode was rationalized in several papers [99,137]. The G/P mode requires an instrument that can change the deposition mode really fast. While such instruments were scarce even a decade ago, nowadays more and more electrochemical instruments fulfill this criterion. If the reader has performed the homework well up to this point, it is needless to tell that the P/G mode would lead to the least benefit and the biggest problems in optimizing the deposition conditions, and hence it cannot be found in the practice.

3.2.3. Relationship with other chemically modulated structures Although this summary accounts for multilayer films only, it has to be emphasized that the principles of the electrodeposition of some other nanostructures are very similar to or even identical with those discussed above. The most closely related group of electrodeposited nanostructures is the multilayered nanowires.

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Nanowires are produced by using a template [150], which can be an etched plastic (mostly polycarbonate or polyamide) membrane or an anodic aluminum oxide. While the former always exhibits a disordered (random) pore system, the latter can be elaborated with a highly ordered hexagonal pore structure. The typical pore diameter is 50–500 nm for both cases. For etched plastic membranes, there is some freedom to adjust the pore diameter to pore distance ratio, whereas this parameter is practically fixed for anodic aluminum oxide templates. Regardless of the material and degree of ordering of the porous template, the common feature is that the aspect ratio is very high (up to 104), which makes it impossible to fill up the pores by sputtering or evaporation. However, electrodeposition offers a unique opportunity to form columnar deposits that fill up the pores from the bottom to the top. For electrodeposition, of course, one side of the template has to be covered with a conducting layer. Nanowires can also be produced with a chemical modulation along their length [151]. If the length of a section in the nanowire is much smaller than the wire diameter, one can speak about multilayered nanowires. In contrast, when the section length is larger than the wire diameter, the nanostructure can be called a multisegment nanowire. In case of both modulated nanowire structures, the pore diameter is small enough so that the growing metal column mostly forms a single crystal at the direction perpendicular to the wire axis, although a long nanowire can be built up from several crystals along the direction of the axis. This growth mode reduces the probability of the electron scattering events related to the crystal boundaries. Keeping the GMR effect in mind, one can easily recognize that for multilayered films the GMR measurement is possible only in the CIP geometry (cf. Fig. 2, left panel). However, it naturally follows that the resistance of a nanowire (or nanowire arrangement) can be measured with a test current flowing parallel to the wire axis; hence, one can obtain the GMR in the CPP geometry (cf. Fig. 2, right panel). The electrons traveling along the wire axis cross the layer boundaries with a much larger probability than electrons traveling parallel to the layer planes in a film. Therefore, the GMR effect observed in multilayered nanowires is higher than in multilayer films; or, alternatively, the layer thicknesses in the former case can be much larger to achieve the same effect than on multilayer films. The larger scale of layer thicknesses is an advantage from the point of view of the ease of sample preparation. On one hand, at 10 nm layer thickness the tolerance in the smoothness of the layer interfaces is much higher than for 1 nm thick layers and the danger of pinhole formation is negligible. On the other hand, the precision of the even layer formation in order to ensure the coupling of the magnetizations of the neighboring magnetic layers is no longer required. The principles for adjusting the electrolyte composition are the same for the deposition of multilayer films or nanowires with a modulated structure. However, it has to be recognized that the diffusion is more hindered within a pore than around a film exposed to the electrolyte at a large surface area. This is because the template thickness defines a zone where convection is impossible. Hence, the concentration of the NM metal ions can (and, from practical point of view, should) be higher in order to avoid very large deposition times. In contrast to the commonality of the electrolyte features in case of film and nanowire electrodeposition, the selection of the control mode is very restricted for multilayered nanowires. The application of the G mode is practically impossible here. At the beginning of the deposition, a diffusion field has to be built up within the electrolyte that fills the nanochannels; hence, the surface concentrations of the metal ions change significantly. When the pore is almost filled up with the wire deposited, the diffusion zone within the pore is shrunk, and the mass transport in the bulk electrolyte becomes more important. The change of the mass transport conditions during the wire formation is so huge that it is not possible to choose a fixed current that always can be passed with a good current efficiency. Therefore, the P mode is the only viable means to regulate the deposition of both types of layer, although the ohmic drop changes significantly from the beginning of the channel filling to the end due to the decrease of the electrolyte resistance within the channel. 3.3. Dual-bath technique Apparently the dual-bath technique offers more freedom for the sample production than the single-bath technique. As its name clearly tells, different electrolyte solutions are used in the dual-bath technique to deposit the layers of different composition. Therefore, the application of electrolyte

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components and the adjustment of their concentration can be independent of each other. Moreover, the magnetic layer can be void of the non-magnetic element. Should the variety of applicable electrolytes seem to be an unbeatable advantage as compared to the single-bath technique, the change of the electrolytes leads to larger problems than all gains provided by the independence of the deposition of the FM and NM layers. No wonder that the number of studies dealing with the dual-bath technique is negligible as compared to those applying the singlebath method. It is, therefore, worthwhile of presenting the difficulties arising in the case of various practical realizations of the dual-bath technique. In one type of the dual-bath systems, the cathode is transferred between two separated electrolyte tanks or electrolysis cells that both are equipped with an anode (and with a reference electrode when the cell is operated in the P mode). The cathodic deposition implies the reduction of the oxide formed during the sample transfer. However, the entire process comprising oxidation, stripping away of the oxide formed and nucleation of the new layer is expected to make the interface more diffuse than in the single-bath method. This surely leads to a reduced GMR effect. In spite of all these difficulties, successful multilayer deposition experiments have been reported and samples with fairly high GMR were obtained [75,120,127,140,143]. The exact details of the sample transfer methods, however, cannot be usually found in the papers published. It is assumed that the samples were rinsed between the immersion steps applying the various baths, but these details are not given. It could not be clarified either how much time the sample contacted the next electrolyte solution before the deposition could be started. The other type of multilayer deposition method from alternating baths is when the cathode is fixed in a conventional single electrochemical cell, and the composition of the electrolyte is modulated by using an alternating injection system. According to Hayashi et al. [107], the optimized electrolyte concentrations in this alternating injection system were much lower than in the single bath method, and 5 mM Co2+ concentration was sufficient to obtain multilayers with GMR. Due to the fact that the codeposition of the more noble element with the FM metal cannot take place since the two elements are never present in the cell at the same time, the deposition conditions of the FM metal can be chosen without considering the NM metal content of the magnetic layer. This is the reason why the difference in the deposition potential of the NM and FM metal could be as low as 200 mV [107], a value which would be unprecedentedly low for a single-bath system. In spite of the achievements with the double-bath systems, the single-bath technique remained the primary method for the deposition of multilayers with GMR. 3.4. Cell construction, electrodes and electrode arrangement Electrodeposition of multilayers with well-defined layer thicknesses requires an even current distribution at the cathode surface. The local current density at a particular area of the cathode cannot be controlled; instead, the average current density can be determined from the total current and the apparent surface area of the cathode. In order to achieve a uniform current distribution on the cathode, a few guidelines have to be obeyed. The parallel arrangement of the counterelectrode (anode) and the working electrode (cathode) is favorable for an even current distribution. When the counterelectrode is a platinum gauze (mesh) or spiral, the parallel arrangement is fulfilled if the overall electrode lengths and widths are about the same, the mean plane of the counterelectrode is parallel to the cathode and the feature size of the structured counterelectrode (wire diameter, spiral diameter etc.) is at least two orders of magnitude smaller than the diameter of the active cathode surface. When the criterion for parallel cell geometry is fulfilled, the current distribution in a stagnant solution is still influenced by the ratio of the anode-cathode distance to the electrode diameter. The larger this ratio, the larger the side effect at the edge of the electrodes. Therefore, deposits produced on a free-standing cathode will always exhibit some edge effect. The uneven layer thickness caused by the electric effects is enlarged by the easier accessibility of the edge of the cathode by the mass transport of the electrolyte components. A detailed analysis of these effects can be found for d.c.-plated NiP deposits in Ref. [222]. The edge effects originating from both the uneven mass transport and the electric field distribution have to be eliminated. This can be carried out in a stagnant solution by applying recessed electrode

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geometry where the cathode fills the entire cross section of the cell at least in the vicinity of the cathode. It is easy to recognize that once the depth of the recess is at least as large as the cathode diameter, both the electrode accessibility by the electrolyte component and the current distribution can be optimum. The problem of the electrode accessibility can be overcome by using a rotating disc electrode, although the problem of the uneven current distribution may still prevail here. The application of a rotating cylinder electrode can also be a solution for the problem of the current distribution, but the stirring efficiency will surely be uneven in this case due to the recess. Also, cell arrangements with rotating electrodes are unlikely to be used for technical applications where a predefined area of a workpiece has to be covered. Besides the mutual position of the anode and the cathode, the vertical or horizontal cathode position lead to different deposits. If the cathode is vertical and the electrolyte in the vicinity of the cathode is depleted during the deposition, the density of the electrolyte decreases. Depending on the viscosity of the electrolyte and the stirring rate, the decreased electrolyte density may result in a spontaneous convection by producing an electrolyte flow to the upward direction (buoyancy effect) [209]. In this case, the electrolyte flow itself changes the component distribution in the electrolyte around the cathode, and the bottom part of the deposit can grow faster because it is always flushed with some fresh electrolyte. When the cathode is in a horizontal position, the depletion due to the deposition also takes place, but the gravitational instability of the electrolyte and the resulting upward flow are produced at a much higher density difference only. The most important effects of cell shape are presented in Fig. 17. An uneven deposit thickness with varying composition and the variation of the MR within a single specimen was demonstrated [77,82] when inappropriate deposition geometry and cell construction were applied (cf. Fig. 17a and b). By introducing a tubular cell with an upward facing cathode that fills the total cross-section of the bottom part of the cell (Fig. 17c) [99,138], it was possible to eliminate all unwanted effects and to obtain laterally homogeneous multilayer deposits. There is some freedom in choosing the material of the counterelectrode. Inert electrodes are often sufficient. An electrode is usually called inert if the electrode reaction taking place thereon has no products that contaminate the bath used. The most often used members of this category are platinum, gold or various forms of carbon, the counterelectrode reaction being the oxygen evolution. Sometimes less demanding electrodes that passivate under anodic conditions can also be used, especially when an insoluble layer is produced on their surface upon oxidation besides the evolution of oxygen (Pb, Ti). However, the application of inert anodes can cause some problems, too. If the buffering capacity of the electrolyte used is fairly little, the oxygen evolution can change the electrolyte pH significantly. The anodes are capable of passivating by the production of the higher-valence oxide of the FM metal

(a)

(b)

(c)

Fig. 17. Impact of the shape and position of the cathode on the deposition conditions. (a) Current distribution between two identical free-standing electrodes. The current density at the edges is higher than at the center due to both the accessibility and the inclination of the current lines. (b) Demonstration of the buoyancy effect. The more and more depleted zones are indicated with darker and darker areas. The arrow shows the flow direction of the electrolyte. (c) Current distribution around a recessed cathode (bottom electrode).

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ions. Also, the production of dissolved Fe3+ is undesired because it can reach the cathode and modify the deposition mechanism entirely. These unwanted effects can be efficiently suppressed by separating the counterelectrode chamber by a membrane or a glass frit. Another concept for choosing the material of the counterelectrode is to use a sacrificial anode. In this case, the anode material has to be the same as the most noble component of the electrolyte; otherwise the spontaneous exchange reaction would lead to a depletion of the electrolyte. For instance, a Cu sacrificial anode for the deposition of Co/Cu multilayers is appropriate, but Co anode would lead to a spontaneous reduction of Cu even without a current. The drawback of the application of a sacrificial anode is that its dissolution increases the concentration of the NM metal ions since it dissolves not only during the deposition of the NM metal but during the deposition of the FM metal, too. Therefore, the cell volume should be adjusted to the expected amount of ions produced by the anode dissolution so that the bath concentration change does not modify the bulk electrolyte concentrations significantly. For the application of the G mode, a cell with two electrodes is sufficient. However, when either of the layers is deposited in the P mode, the application of a reference electrode is indispensable. Reference electrodes that are generally used in various electrochemical experiments are all equilibrium systems where referencing is based on the exact potential that is achieved by maintaining an electrochemical equilibrium. All other electrodes occasionally used for referencing whose potential is determined by a steady-state process are to be called quasireference electrodes, and their stability is much behind that of the equilibrium systems. Three basic reference electrodes are commonly used in the practice of multilayer electrodeposition. These reference electrodes with the abbreviation, the equilibrium electrode reaction and the corresponding electrode potential for the electrolyte saturated with a salt of the anion are listed in Table 1. The electrode potential of the reference electrodes is a function of the composition of the internal electrolytes; hence reference electrodes with a specific salt concentration exhibit different potentials. For the sake of uniformity, all potential values reported for multilayer preparation in the studies reviewed are transformed to the standard hydrogen electrode (SHE) scale in this work. Potential values in most cases are given with a 10 mV accuracy because the change in electrode potential within such an interval can cause an insignificant change only in the deposit properties. Although the selection of the reference electrode is usually not important, the mercury sulfate electrode is recommended if the electrolyte has to be void of chloride ions. The reference electrodes themselves are often connected to the cell via a salt bridge, i.e., by using an intermediate electrolyte to separate the reference electrode from the bath. This separation is essential when the bath contains ions whose reduction potential is higher than that of the reference electrode, otherwise the reference electrode is polarized by another redox system. The position of the reference electrode (or that of the junction of the reference electrode compartment) is crucial in the cell design for various reasons, and the requirements of a correct positioning are somewhat contradictory. First, the reference electrode should not disturb the even current distribution on the cathode surface. Secondly, the reference electrode has to be close to the working electrode in order to decrease the potential drop that arises due to the resistance of the electrolyte layer between the working and reference electrode and the current passing through this layer. This so-called uncompensated ohmic drop is usually smaller than the 1 mV accuracy of most electrochemical instruments when the current density is below the mA/cm2 range. However, this ohmic drop is rather large

Table 1 Properties of the most common reference electrodes used in experiments of GMR multilayer deposition. Name

Electrode reaction

Silver/silver chloride electrode (Ag/AgCl) Saturated calomel electrode (SCE) Saturated mercury sulfate electrode (SSE or SMSE)

AgCl + e = Ag + Cl Hg2Cl2 + 2e = 2Hg + 2 Cl Hg2 SO4 þ 2e ¼ 2Hg þ

SO2 4

Electrolyte saturated with . . .

Potential vs. SHE (V)

KCl KCl K2SO4

0.1976 0.2415 0.650

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when the magnetic layer is deposited at high current density. This is why the deposition potential of the magnetic layer has to be treated with caution since the values reported may be significantly influenced by the uncompensated solution resistance. The uncertainty of the ohmic drop is a further reason for the application of the G/P mode. 3.5. Role of additives and pH The family of bath components that are often termed as additives is difficult to describe with rigorous categories. In most cases, researchers use this term for bath components that are neither the source compound for the deposit-to-be nor a supporting electrolyte (which merely increases the solution conductivity without having any significant effect on the deposition mechanism and kinetics). The name ‘‘additive” refers to a bath component whose electrode reaction contributes negligibly to the total current, but it is still effective in regulating some important deposit property. Additives of importance in galvanotechnology belong to a large variety of compounds. They can be salts with as simple an anion as chloride ion, organic compounds with some heteroatom like S or N, aromatic compounds, complex-forming molecules with chelating capability, soluble polymers with polar groups and so on. Besides their chemical variety, the common features of these bath components are that (i) they bear one or several non-bonding electron pair(s) which (ii) can interact with either the electrode surface or the metal ions in the solution; therefore (iii) they are capable of substantially modifying the reaction kinetics and hence leading to a desired deposit property. Due to the highly empirical nature of our knowledge on additives, there is no principle which could help to predict the impact of an additive on a given deposit property. Moreover, the impact of an additive is usually not universal in the sense that an additive with favorable influence on the deposition of a particular metal might be even harmful to the deposition of another metal. Additives were developed for regulating the coating properties in galvanotechnology and the use of additives was mostly optimized for d.c. deposition. Besides the empirical categorization of additives [223], there are various attempts to calculate the optimum deposition conditions for obtaining smooth deposit surfaces [224–226] or to calculate the surface morphology for given deposition conditions [227–229]. These methods apply heavy mathematics demanding large computational capacity and work for one-component deposits only. Since the early works on ED multilayers adopted some baths originally developed for the purpose of galvanotechnology, we must deal with the additives as well. The primary role of the complexing agent is that they modify the speciation of the solution. The formation of a complex can be described with the following equation: Mez+ + k Ln = [MeLk](knz). The notation L stems from the word ‘‘ligand”. The complex formation leads to an equilibrium, and it is characterized with the stability constant. Depending on the number of the electron pairs by which the ligand binds to the central metal ion, one can speak about monodentate, bidentate, etc. ligand. Ligands with more than one electron pairs complexing a metal ion are also called chelates. The ligand usually replaces the complexing water molecule, which is a mobile ligand as compared to those that can replace it. The importance of the complex formation is that the fairly simple overall electrode reaction involving a metal ion

Mezþ þ ze ¼ Me;

ð13Þ

is modified as follows: n

½MeLk ðknzÞ þ ze ¼ Me þ kL :

ð14Þ

It cannot be emphasized strongly enough that the release of the ligand is the part of the electrode reaction, and the deposition of a complexed metal ion cannot be treated by means of the equation of complex formation equilibrium as if only the ‘‘free” uncomplexed metal ion were able to react at the electrode surface. This means that the standard potential of a complex metal electrode is more negative than that of the uncomplexed metal electrode. Hence, the formation of a complex compound yields an opportunity to tune the deposition potential of a metal. This effect has a great importance when the task is the deposition of an alloy. According to the general experience, it is favorable for the electrocrystallization of an alloy if the deposition potentials of its components are close to each

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other. The formation of complex compounds is a simple mean to bring closer the deposition potentials of metals that are otherwise miscible but the alloy formation would be difficult due to the big difference in the deposition potentials. The pursuit of galvanotechnology to ease the alloy formation, however, does not meet the requirement of ED multilayer fabrication. Especially when GMR is important, one has to strive to achieve a sharp interface between the FM and NM layer, and this does not require at all the formation of an alloy, whether or not the FM and NM metals are miscible. Therefore, the ability of codeposition as an alloy is not desired for the deposition of ED multilayers with GMR. Other additives used in galvanotechnology are classified as stress relievers, levelers and brighteners. These additives are usually applied in a very little concentration (a few millimol/liter as a maximum), as opposed to the metal salts and the complexing agents. They behave in a similar manner as catalyzers in the sense that their electrode reaction is negligible, however, they have a fundamental impact on the properties of the deposit. They are used for regulating the residual internal stress, hardness, tensile stress, corrosion resistance and the smoothness and brightness of the surface. Very little is known about the detail of how a particular additive exerts its effect. The common element in the mechanism assumed is that they are able to adsorb at the electrode surface. The preferred site of adsorption is the kink (or half-crystal) position [216]; i.e., the same site where the growth of the crystal by the incorporation of a new metal atom can take place by forming the largest number of bonds and also by reproducing a surface site with identical type. The incorporation of the adatom to the kink position minimizes the overall energy of the surface – adatom ensemble. Depending on which control parameter is regulated, the impact of the adsorption of an additive can be explained with slightly different approaches: (i) In case of the G mode, the adsorption of a foreign molecule (which will not incorporate into the deposit) at the preferred growth position of a crystal means that the current has to flow elsewhere, and the charge transfer will take place at surface positions of higher energy. Consequently, the cathode potential has to decrease (i.e., the potential has to be more negative) to provide the higher activation energy of the charge transfer. Since the adatoms produced cannot diffuse to the kink positions already occupied by the molecules of the additive, the surface concentration of the adatom increases, which leads to a high nucleation rate of new crystals. (ii) In the P mode, the driving force is constant, but the charge transfer becomes more hindered. This means that the current has to decrease if the potential is unchanged, and a more cathodic potential has to be applied to maintain the same current. Simply speaking, the polarization curves appear to be shifted to the cathodic direction. Regardless of the exact mechanism of the additive effects, the increase in the nucleation rate and decrease in the crystal growth rate results in the reduction of the crystallite size. This effect is well known in the electrochemistry literature [211,230]. The strength of the adsorption is often rated by the ‘‘inhibition intensity”, which is an ill-defined and qualitative notation of the additive effect. The problem with the notion of the inhibition intensity is that it is measured by the resulting sample properties rather than with an independent definition. Nanocrystalline deposits are often deposited in the presence of a large concentration of additive. The decrease in the crystallite size always leads to the increase in resistivity, simply because the electrons are frequently scattered at the grain boundaries where the atomic ordering is imperfect. Since the MR ratio is referred to the zero-field resistivity, a drastic increase in the latter quantity leads to a decrease in the MR ratio to the same extent. The commonality in the impact mechanism of the complexing and additive agents makes it possible to construct a single diagram that shows the consequences of their application. This is presented in Fig. 18. A direct comparison of the baths used is difficult since the primary aim of research has always been the investigation of the layer thickness dependence of the GMR properties, and once a bath worked, it was used throughout a particular study. Nevertheless, a few trends can be established. A great majority of the electrolyte compositions for multilayer deposition in the single-bath method is based on the sulfate of the metals. However, as the heritage of the galvanotechnology, several

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Application of complexing agents

Application of strongly adsorbing additives

Change in the mechanism of the deposition

Large charge transfer resistance

Increase in nucleation rate

Loss in preferred crystallographic orientation

Blocking the kink positions of the crystal

Accumulation of intermediates

Decrease in grain size

Periodic adsorption-desorption

Occurrence of capacitive current components Distortion of the applied signal

Formation of intermixed boundary layer of significant thickness

Increase in resistivity

Decrease in GMR

Fig. 18. The scheme by which the complexing agents and surface-active additives are thought to influence the deposition and GMR of ED multilayers.

studies applied citrate-type baths to produce Ni–Cu/Cu, Co–Cu/Cu and Co–Ni–Cu/Cu multilayers (literature references are listed in the corresponding parts of Section 5). These baths lead to deposits with lower GMR than baths without any complexing agent added. In parallel to the diminished GMR, the saturation fields for multilayer samples obtained with citrate-type baths are usually very high, indicating a significant SPM contribution, which is attributed to the small size of the magnetic entities. The small size of the magnetic entity can be a side effect of the small grain size, although there is no direct correlation between the grain structure and magnetic domain structure of the samples. The impact of the presence of some additive(s) in the bath was mentioned a few times as a collateral information in a study devoted to mostly physical measurements on FM/NM multilayers [17]. A few studies were published only that deal with the additive effect as the major goal of research [74,78]. The general experience was that the presence of the additive is harmful for the formation of the layer structure. There is an agreement in the literature that the structural imperfection caused by the additive results in the loss in GMR. The typical additives tested included surfactants used in galvanotechnology (proprietary names are known only [17]), common surfactants as sodium dodecilsulfate (SDS) [74], sulfur containing organic compounds [74] or sodium chloride [78]. The effect of the pH on the sample properties is usually dealt with as an independent issue. In practically all studies where the pH effect was discussed [28,99,102], the optimum pH was between 1 and 3, and it was mentioned that an increase in pH leads to the reduction of the GMR magnitude. The lower limit of the applicable pH is seldom mentioned. It is expected that solutions with very low pH are also inappropriate for obtaining a compact and well-structured multilayer deposit due to the unwanted effect of the intensive hydrogen evolution, especially in the high-current pulse. In a few cases, it was mentioned that the incorporation of the NM element into the magnetic layer also shows a pHdependence [102]. However, the analysis of the additive effect implies that the pH effect can be discussed within the same set of arguments. Transition metal ions in an aqueous electrolyte form complex with water molecules in the absence of other complexing agents. As the solution becomes more and more alkaline, the hydrated metal ion can behave as an acid:

½MeðH2 OÞn zþ þ OH ¼ ½MeðH2 OÞn1 OHðz1Þþ þ H2 O:

ð15Þ

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The reduction of the metal ions at the cathode is a multistep process. The reduction of a bivalent cation (such as Ni2+, Co2+, Fe2+, Cu2+) takes place via a monovalent intermediate. The relative hindrance of the first and second electron transfer is a function of the formation and decomposition rate of the intermediate and it is also influenced by the pH since the monovalent intermediate adsorbs at the surface together with a negative ion, mostly OH, forming a neutral species. In several reaction mechanisms, the higher pH promotes the formation of the partly reduced intermediate. This leads to the accumulation of adatoms at the cathode surface, having the same effect as that shown in the scheme in Fig. 18. It is not to be forgotten that the dissociation of the water molecule produces hydroxide ions that have non-bonding electron pairs by which they are capable of coordinating to the active sites of the growing crystal. This effect points to the same direction as other anionic species among the additives. With successive proton release from the metal hydroxide complex, the metal complex becomes neutral, and the further loss of water molecules leads to the precipitation of the metal hydroxide. This process takes place first at the location where the hydroxide ion concentration is the highest, or, in other words, where the pH is the highest. The increase in pH is related to the hydrogen evolution at the cathode. Once the metal hydroxide precipitates, the incorporation of non-metallic inclusions in the deposit leads to a structural disorder and an increase in electron scattering, both being detrimental for the GMR. The hydrolysis of the metal ions can be effectively controlled by weak acids, for instance, boric acid. The acidic dissociation of boric acid starts in the same pH range where the precipitation of most transition metal hydroxide begins; hence, boric acid is effective only when there is indeed a danger of a significant shift of pH at the vicinity of the cathode. Boric acid is commonly used because it is an efficient pH regulating agent without a significant influence on the deposit properties. 3.6. Control of electrochemical processes It was already discussed in Section 3.2.2 that, if the deposition conditions of the more noble metal are not properly chosen, the less noble FM metal can dissolve during the deposition of the NM metal. Regardless of the difference in the terminology (i.e., ‘‘exchange reaction” for the G mode and ‘‘dissolution” for the P mode), the result is the same: a part of the previously deposited FM layer is lost or replaced by the NM element. The deposition of the NM element with improper conditions leaves its fingerprint on the overall deposit composition and microstructure. If a particular layer of the multilayer stack is deposited with a charge Qi the quantity of the metal deposited can be obtained from Faraday’s law:

ni ¼

Q i gi ; zi F

ð16Þ

where g is current efficiency (for the other symbols, see the text below Eq. (11)). Here, Q can be obR tained as It if the current is constant or by the integral I(t)dt if the current varies and is monitored in time. A general solution for the extent of the metal exchange can be given if zNM = zFM (this is the case for the most common (Fe, Co, Ni)/Cu multilayers). In this case, the nominal molar fraction of the more noble NM metal in a two-component multilayer is:

yNM ¼

nNM Q þ y0NM gFM Q FM ¼ NM ; nNM þ nFM Q NM þ gFM Q FM

ð17Þ

where the current efficiency for the NM layer is taken unity as it was usually found and gFM < 1. The molar fraction of the NM metal in the FM layer is y0NM . If the exchange of the metals takes place, the molar fraction of the NM element in the deposit measured (yNM ) will be higher than the nominal value and it can be written as

y0NM ¼

Q NM þ y0NM gFM Q FM þ Q EXC ; Q NM þ gFM Q FM

ð18Þ

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and QEXC is the charge in a period which leads to the deposition of the NM metal instead of the FM metal. In other words, QEXC/zF is the amount of the more noble metal gained or that of the less noble metal lost within a repeat distance as compared to the nominal values. Considering that the molar weight is M = m/n, the density is q = m/V and that the total sample volume can also be obtained as V = Ad (A is the sample surface area and d is the layer thickness), one can obtain for the thickness of a layer that

d¼

Qg M : zFA q

ð19Þ

By substituting Q = QNM or Q = QNM + QEXC, one can get the nominal and actual NM layer thickness, respectively, and a similar procedure leads to the FM layer thickness. Since for Co, Ni and Cu the difference in both the molar weight and the density is little, the difference between the nominal and actual layer thickness will be the same within the experimental error for multilayers composed of these elements. Composition analysis after the deposition can be used for the estimation of the excess amount of NM element. However, there are some opportunities to control the deposition process in situ as well. Concerning the deposition control by following the electric parameters during the deposition, the G mode is a bad choice since the potential (which is not regulated but can be recorded) does not bear any quantitative information on the amounts of reacting components. In contrast, the P mode is appropriate to follow the dissolution of the previously deposited less noble metal and also to optimize the deposition potential by analyzing the current transients recorded for the low-current pulse. The following train of thoughts follows the guidelines given in Ref. [231]. Typical current transients for the NM layer deposition in the P mode are shown in Fig. 19. The pulse always starts with a sharp positive spike. Disregarding any instrumental artifact, this positive spike is related to the behavior of the electric double layer at the metal/electrolyte interface. The interfacial capacity of the double layer is around 20 lF/cm2 for strong electrolytes of about 1 mol/l concentration, independently of the nature of species present. Taking into account also that the typical order of magnitude of the electrode surface area is 1 cm2 and the uncompensated solution resistance is between 1

-4

6.0x10

E (Cu) / mV vs. SHE: -260 -340 -390 -405

-4

4.0x10

A

-4

i / Acm

-2

2.0x10

B 0.0 -4

-2.0x10

C

-4

-4.0x10

D 0.1

1

10

t/s Fig. 19. Potentiostatic current transients recorded during the deposition of the NM Cu layer in the P mode as a function of the Cu deposition potential. Sample: Co–Cu/Cu multilayer deposited by using a chloride-type bath in G/P mode. Curve A: long anodic transient with significant Co dissolution. Curve B: the transient is as short as possible and the current is never more negative than that corresponding to the diffusion-limited Cu deposition (no Co deposition in the Cu pulse; ideal case). Curve C: very negative current at the beginning of the Cu pulse, Co codeposition takes place with vanishing intensity. Curve D: the Co codeposition persists throughout the Cu pulse. Reprinted from Ref. [231] with permission of Elsevier.

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and 10 ohm, it can be seen that the time constant of the double layer charge and discharge process is s = RC < 1 ms. Nevertheless, a current transient of capacitive origin must always be present. If other processes are listed in the order of the time constants, the next is the adsorption (or desorption) of species. This current contribution is quite uncertain and depends on the electrolyte composition. After the full stabilization of the double layer and the decay of the adsorption/desorption process, all deviation from the steady-state diffusion-limited deposition of the NM layer is to be attributed to the dissolution of the previously deposited FM layer. The optimization method can be clearly seen from Fig. 19. The anodic transient must be as short as possible, and the cathodic current should not exceed the diffusion-limited Cu deposition current because this would correspond to the codeposition of the less noble FM metal with the more noble one. The current transient is also suitable for a rough estimation of the excess NM metal deposition and the FM metal loss if the dissolution rate of the FM metal is not very high. This latter condition has to be fulfilled so that the diffusion-limited deposition of the NM metal is undisturbed and its deposition rate is independent of the FM metal dissolution [99]. The difference of the actual and the diffusion-limited current has to be integrated, and the resulting charge corresponds to the charge contribution of all anodic processes. If the current transient is integrated and the linear section of the charge vs. time function is fitted with a linear function, the intercept with the y-axis yields the same charge [99]. If the dissolution rate of the FM layer is too high, the deposition rate of the NM metal is certainly influenced by the unwanted dissolution, and its deposition rate is not constant throughout the pulse; hence, the composition estimation is not possible [99]. Another means to follow the deposition process during the NM layer deposition is the application of an electrochemical quartz crystal microbalance (EQCM). EQCM can be used for both G and P modes (nevertheless, the application in the P mode was reported only [129]). The EQCM yields information on the weight of the deposit, and hence it is independent of the electric parameters. The disadvantage of the EQCM is that the difference in the molar weight per ion charge for typical FM and NM metals is very little, and hence the EQCM is fairly insensitive to detect the metal exchange. Also, the data acquisition frequency of the EQCM is usually smaller than that of the current recording. A very simple optimization method can be when a bath is prepared with all components except for the compound of the more noble metal and the rest potential of the less noble metal is measured [17]. Depending on the conditions, this potential can be either the equilibrium potential or the corrosion potential. If a good care is taken to record the equilibrium potential, the potential value is close to the optimum. A complete coincidence of the optimum deposition potential of the NM metal and the equilibrium potential of the FM metal cannot be expected since the concentration of the FM metal salt close to the cathode is smaller right after the high-current pulse than in the equilibrium. The loss of the magnetic layer by means of the exchange reaction (G mode) or by the FM metal dissolution (P mode) results in a layer boundary undulation and intermixing of the components. This is simply because of the fact that the NM metal cannot be deposited at the same position of the cathode where the previously deposited FM metal dissolves. Therefore, the multilayer structure for samples obtained with non-optimized conditions is less perfect. The largest the amount of the dissolved FM metal, the more likely that the FM layer will be discontinuous, and the entire deposit may become porous [232]. The dissolution of the FM metal during the low-current (or low-potential) pulse can be identified as one of the major reasons of the occurrence of the SPM character.

4. Structure of ED multilayer films Before presenting the GMR results reported for ED multilayers, it is necessary to have an overview of our current knowledge about the structure of such systems since the crystal structure, interface properties and microstructural features such as texture, grain size, lattice defects, etc. may have a strong influence on the magnetic and magnetotransport behavior. The direct way of getting information on multilayer structure is via transmission electron microscopy (TEM) whereas various X-ray diffraction (XRD) techniques are also widely used for deducing structural information. We shall first discuss separately some results obtained by these structural investigation methods and, then, consider

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how the structure can be influenced by the preparation conditions. On this latter issue, knowledge is still fairly poor and this offers place for significant improvement in GMR of ED multilayers if progress can be achieved in this field in the future.

4.1. TEM study of ED multilayer films For such studies, the multilayer sample should be thinned down to a thickness of about 10–20 nm in order to make it transparent for the electron beam. This is a rather tedious procedure and since TEM equipments are also not generally available, the TEM study can usually be carried out on a few selected ED multilayers within a large series only. TEM studies on ED Ni–Cu/Cu [29,72,233–242], Co–Ni–Cu/Cu [11,16,43,50,69], Co–Cu/Cu [24,33–35,41,77,78,80,94,98,116,128,141,149,243–246], Co–Fe–Cu/Cu [53], Co–Fe–Ni–Cu/Cu [85], Co–Ag/Ag [148,247], Co–Au/Au [248], Fe–Ni–Cu/Cu [249] and Co–Ru/Ru [83] multilayers have been performed in order to reveal some common microstructural features of such systems. The sample thinning can be performed in cross-sectional or in plane-view geometry as illustrated schematically in Fig. 20. For a cross-sectional TEM study, the multilayer is thinned along the direction perpendicular to the picture plane and the direction of the electron beam used for the TEM imaging is perpendicular to the picture plane. For plane-view TEM imaging, sample thinning is carried out by etching from both the final and the substrate side surface of the multilayer. In this case, the multilayer part with black stripes in the left panel remains for TEM study carried out with an electron beam perpendicular to the layer planes. The right panel shows the schematic cross-section of a multilayer with inclined layer planes with respect to the horizontal substrate. After sample thinning, the thin section between the two horizontal lines will remain for a plane-view TEM study. For samples with inclined layers, thinning parallel to the substrate leads to a TEM picture showing contour lines similar to a map. Fig. 21 shows cross-sectional (left panel) and plane-view (right panel) TEM images of ED Ni-Cu/Cu multilayers [72]. One can clearly observe the layered structure in the cross-sectional image. It should be noted, however, that one can not infer the actual layer thicknesses from such TEM pictures. The reason is that the atomic numbers of the FM and NM elements (e.g., Co and Cu or Ni and Cu which are accessible for the ED technique) are very close to each other and, thus, the electron scattering contrasts between these metal pairs are very weak. Therefore, TEM images like those shown in the left panel of Fig. 21 can only be obtained under conditions of strong defocus of the electron microscope and the image obtained does not reflect the actual thicknesses of the individual layers. Of course, the bilayer repeat length (K = dFM + dNM) can still be fairly correctly determined from the cross-sectional images as was shown in Refs. [29,128].

Fig. 20. Schematic illustration of the in-plane multilayer thinning procedure for TEM observations. The growth direction is vertical (upward) in both cases. The layer planes are parallel to the horizontal substrate on the left panel whereas they are inclined to the substrate on the right panel. The horizontal white and black (or grayish) stripes represent the individual magnetic or non-magnetic layers. Reproduced from Ref. [29] with permission, Ó Carl Hanser Verlag, München.

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Fig. 21. Left: cross-sectional TEM image of an ED Ni81Cu19(2.5 nm)/Cu(1.4 nm) multilayer sample. Right: plane-view TEM image of an ED Ni81Cu19(10 nm)/Cu(1.4 nm) multilayer sample. The layer thicknesses are nominal. Reprinted from Ref. [72] by permission of The Electrochemical Society.

Fig. 22. Left: low-magnification cross-sectional TEM image of an ED Ni–Cu(3.8 nm)/Cu(0.7 nm) multilayer. A column of several 100 nm width and extending through the total multilayer thickness can be clearly identified in the middle. Reproduced from Ref. [29] with permission, Ó Carl Hanser Verlag, München. Right: cross-sectional TEM picture of an ED Co–Cu(9.4 nm)/Cu(1.1 nm) multilayer close to the top of a columnar grain where the canting of the layer planes is the most characteristic. Layer thicknesses are nominal values. Reprinted from Ref. [243] with permission of Elsevier.

In order to understand the formation of the spiderweb-like microstructure shown by the planeview TEM image in the right panel of Fig. 21, we should first consider the low-magnification cross-sectional TEM image of an ED Ni–Cu/Cu multilayer (Fig. 22, left panel) where a columnar growth of the multilayer can be observed. The right panel of Fig. 22 shows the top region of a similar ED Co–Cu/Cu multilayer at higher magnification where a strong canting of the layer planes can be observed. In order to reduce the interfacial stress energy arising from the lattice mismatch between the two kinds of layers, the chemical modulation plane is canted whereas the crystal lattice planes remain parallel to the substrate. Cutting such a layered structure in a manner as sketched in the right panel of Fig. 20 for plane-view imaging, the pat-

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Fig. 23. Schematic illustration of the microstructure of ED multilayers as deduced from TEM studies. The thin dark and thick bright layers represent the idealized chemical modulation of the multilayer. The grey-shaded areas between the hexagonal columns correspond to grain boundaries. Reproduced from Ref. [29] with permission, Ó Carl Hanser Verlag, München.

tern shown in the right panel of Fig. 21 is obtained. The image well reveals the layered structure, the columnar growth and the hexagonal shape of the columns (grains). On the basis of such TEM studies, the schematic view shown in Fig. 23 can be considered as a representation of the actual microstructure of an ED multilayer which was deposited on a polycrystalline substrate. Since the magnitude of the GMR is strongly dependent on the thickness of the individual layers, especially on that of the NM spacer layer (see Section 2.4.2), it is important to have means for the determination of the constituent layer thicknesses. A possibility to establish directly the actual layer thicknesses is to carry out a cross-sectional elemental mapping with sufficient lateral resolution. The result of such a TEM analysis [128] is shown in Fig. 24 for an ED Co–Cu/Cu multilayer. Unlike the conventional cross-sectional TEM contrast images, the layer thicknesses elucidated from the elemental map can already be confronted with the result of the chemical analysis [99,128]. The deviation

Fig. 24. Compositional map taken by cross-sectional TEM analysis on the ED multilayer Si/Ta/Cu(20 nm)/[Co(3.4 nm)/ Cu(1.0 nm)]. The Co/Cu layer thicknesses specified are nominal; the actual values are Co(2.0 nm)/Cu(2.4 nm). Reprinted from Ref. [128] with permission of American Scientific Publishers.

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from nominal layer thicknesses due to the maladjustment of the deposition parameters can therefore be verified to be in full agreement with the overall composition analysis [99,128]. Since the magnetoresistance is defined as the ratio DR/Ro where Ro is the background resistance in zero external magnetic field, the quantity Ro should be as small as possible if we want to have large GMR. An important contribution to the background resistance may come from lattice defects such as, e.g., grain boundaries which have a larger resistivity than the crystal interior. Therefore, the multilayer if polycrystalline should consist of possibly large crystal grains. As discussed above, TEM studies revealed that ED multilayers grown on polycrystalline substrates consist of columns which, under favorable preparation conditions, extend through the total multilayer thickness even up to beyond 1 lm. It was demonstrated by cross-sectional TEM studies of ED Ni–Cu/Cu [29], Co–Cu/Cu [78,128] and Co–Ni–Cu/Cu [50] multilayers that the GMR magnitude is unambiguously correlated with the crystal (column) size. A significant GMR can only be obtained if the grains (columns) are sufficiently large. This is understood in a manner that the lateral and vertical extension of the columns should be much larger than the electron mean free path in the multilayer. This is fulfilled, e.g., for the columns shown in Figs. 21 and 22 which extend in both directions to several hundred nanometers. In such a case, the electrons ‘‘feel” as if travelling in a single crystal and grain boundary resistivity contribution to the total resistivity remains negligible. On the other hand, the loss of GMR in some ED Co–Ni–Cu/Cu multilayers [50] could be shown by cross-sectional TEM to be due to the fairly small grains sizes being as small as a few tens of nanometers only. It is noted finally that cross-sectional TEM with sufficiently high resolution including atomic resolution can enable us to directly observe other lattice defects such as stacking faults and twins [149,233] or crystal lattice plane undulations arising due to the lattice mismatch of the constituent layers [233]. 4.2. XRD study of ED multilayer films In view of the complexity of the TEM measurement itself and the tedious cross-sectional sample thinning procedure, the primary source of structural information about multilayers has been obtained by XRD. The typical lateral sample size for ED multilayers is usually of the order of 10 millimeters and such large samples can be conveniently investigated, without the need of any special sample preparation, by most conventional XRD equipments. Such measurements can quickly provide information on multilayer crystal structure, texture (preferred orientation) and grain size. A special feature of XRD is that, for multilayers of sufficiently good quality, so-called superlattice reflections [250] can appear in the XRD pattern which arise due to the periodic repetition of a bilayer unit. Michaelsen [251] presented a simple model for calculating the diffracted intensity (I) for a periodic multilayer film A/B. This intensity consists of three terms:

I ¼ IN ½IA þ IB þ IAB ;

ð20Þ

where IA and IB are the intensities of the constituent layer materials, IAB is a mixed term arising from the assumption of structural coherence along the thickness and IN is a prefactor related to the periodic repetition of the bilayer unit N times. It was shown [252,253] that the mixed term disappears in a noncoherent treatment. Accordingly, the necessary condition for the appearance of satellite reflections is a good structural coherence along the thickness (e.g., fairly uniform layer thicknesses, high smoothness of the surface of the growing layer as well as chemically and topologically sharp interfaces between constituent layers). Michaelsen [251] calculated the XRD intensity for a face-centered cubic (fcc) Co/Cu multilayer which is shown in Fig. 25 for an ideally perfect structure (upper curve) and by assuming some structural disorder leading to a broadening of the diffraction lines (lower curve). The large central peak corresponds to the main (1 1 1) reflection due to diffraction on atoms arranged in an fcc lattice. Lattice strains were not taken into account and an fcc structure was assumed also for Co. A growth of Co/Cu multilayers in a complete fcc stacking sequence has indeed been found for periodicities below about 10 nm in either PD [189,251,254] or ED [128,149,255] samples. The multilayer structure is characterized by the symmetrically spaced satellite reflections (also called superlattice reflections) which arise from a periodic repetition of the bilayer unit. Even the

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Fig. 25. XRD intensity around the (1 1 1) peak of an fcc [Co/Cu]20 multilayer with each type of layers (Co and Cu) containing 20 atomic layers (this sums up to about 8 nm bilayer period) as calculated on the basis of Eq. (20) (upper panel), and after convolution with a Gaussian distribution of width r = 0.1° (lower panel). Note the square-root intensity scale. Reprinted from Ref. [251] with permission of Taylor and Francis Ltd (http://www.informaworld.com).

first-order satellites have much smaller intensity than the main diffraction line and the intensity rapidly decreases for higher-order satellite peaks. It is evident from Fig. 25 that any structural disorder including loss of structural coherence along the thickness reduces the satellite peak intensities. In this manner, the presence and intensity of satellite reflections is often considered as a parameter characterizing the structural quality of a multilayer. From the position of the superlattice satellites, the bilayer repeat period can be determined in a manner as follows. First, we recall the well-known Bragg condition 2disinhi = nk [167] for the observation of constructive interference of diffracted X-rays with wavelength k if both the incident and diffracted beams are inclined at an angle hi to the crystal lattice planes (h,k,l)i which are at a distance di apart from each other and n is an integer number. If we look for the multilayer reflections due to the periodic repetitions of layers A and B with thicknesses dA and dB, respectively, the Bragg condition for the diffracted waves from the (h,k,l)i lattice planes of such an A/B multilayer is

2Ksinhi ¼ mk:

ð21aÞ

Here K = dA + dB is the bilayer length and m is an integer at which the multilayer Bragg condition is fulfilled for the angle hi and for coherently diffracting lattice planes at a distance K apart from each other. Condition (21a) leads to a Bragg peak at the same angle as occurring for diffracting waves being reflected coherently from neighboring atomic lattice planes at a distance of di. The first-order superlattice reflections around this peak appear at the conditions

2K sin hiþ1 ¼ ðm þ 1Þk;

ð21bÞ

2K sin hi1 ¼ ðm  1Þk:

ð21cÞ

and

By subtracting now Eq. (21c) from Eq. (21b), we can derive the expression

K ¼ k=ðsin hiþ1  sin hi1 Þ;

ð22Þ

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with the help of which the bilayer length K can be determined from the measured positions of the superlattice satellite peaks. Apparently, Lashmore and Dariel [256] were the first to show that ED multilayers can be grown in a quality that the XRD pattern exhibits (i) nonsplit Bragg peaks being at a position intermediate of the positions of the corresponding Bragg peaks of the constituent metals and (ii) multilayer satellites, both features providing unambiguous evidence regarding the growth of a sufficiently coherent superlattice structure. This is demonstrated in Fig. 26 for an ED Ni-Cu/Cu multilayer [256] in which two dominating crystal orientations can be observed. Fig. 27 shows the evolution of XRD patterns for a series of ED Co–Cu/Cu multilayers with varying Cu layer thicknesses [80]. The satellites (labeled as 1 and +1) around the (0 0 2) peak approach the main peak with increasing bilayer length (from top to bottom). A strong (0 0 2) texture of the multilayer can be observed here, with some amount of (1 1 1)-oriented crystallites as well. A small fraction of hexagonal close-packed (hcp) Co can be identified by the peak at the lowest angles; its amount, however, strongly decreases with increasing Cu layer thicknesses and it was found to disappear for dCu > 5 nm [80]. The presence of a hcp-Co fraction at low Cu layer thicknesses was reported [149,255] also for other ED Co–Cu/Cu multilayers and was explained by the discontinuous nature of the Cu layers at not sufficiently high coverages which enable the subsequent Co layers to grow on top of the previous Co layers. Beyond a certain Co thickness, there is no more epitaxial fcc constraint and the Co layer continues to grow with the equilibrium hcp structure. Along the same line, it is easy to understand that for larger Cu layer thicknesses the spacer layer continuity improves and the condition promoting the formation of hcp-Co crystallites strongly reduces. It should be noted that the kind and degree of textures shown in Figs. 26 and 27 are not specific to one or other multilayer systems but rather depend in a complex manner on the combination of preparation details (e.g., substrate orientation or deposition mode) as will be discussed in Section 4.3. The bilayer length derived from the XRD satellite peak positions (KXRD) for ED multilayers usually was found to be in good agreement with the nominal values from electrochemical preparation parameters although there seems to be some evidence [29,81,128,149,255] that the XRD data are systematically higher by about 10%. It is noted finally that various sophisticated XRD techniques are available to get more detailed structural information about multilayers which require some modeling and fitting of the experimental data. Wide-angle X-ray scattering (WAXS) can be employed for the simultaneous determination of the

Fig. 26. XRD pattern of an ED Ni-Cu/Cu multilayer prepared with a bilayer period of 5.3 nm on a single-crystal Cu substrate. Reprinted from Ref. [256] by permission of The Electrochemical Society.

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Fig. 27. XRD patterns of ED Co–Cu/Cu multilayers prepared with varying Cu layer thicknesses and constant magnetic layer thickness (1.8 nm) on Si(1 0 0)/Cu(20 nm) substrates. The positions of the various multilayer Bragg peaks [fcc-(0 0 2) and fcc 0)] and the satellite reflections (labeled with 1 and +1) are marked by the vertical arrows. In the figure, (1 1 1) and hcp-(1 01 layer thicknesses are specified in terms of monolayers (ML) whereby 1 ML is approximately 0.2 nm; thus the Cu layer thicknesses given correspond (from top to bottom) to about 1.1 nm, 2.2 nm, 3.3 nm and 4.4 nm. Reprinted from Ref. [80] by permission of The Electrochemical Society.

bilayer thickness and both the thickness and interplanar spacings of the individual component layers [250,257]. Small-angle X-ray scattering (SAXS) experiments offer the opportunity to characterize the continuity of the interfaces between constituent layers [200,258–262]. A further important technique to get information on the quality of interfaces (e.g., interface roughness and chemical intermixing) is X-ray reflectivity (XRR) [200,262]. A recent work [149] describes a thorough study of the formation of microstructure defects during the electrodeposition of Co/Cu multilayers by WAXS and SAXS experiments and modeling. 4.3. Influence of preparation conditions on structure in ED multilayer films The structural features of ED multilayers evidently depend on the deposition conditions although relatively little is known about the interrelation of structure and preparation conditions. Therefore, a few examples can only be given here in which cases this influence manifested quite clearly. One important aspect is the influence of substrate texture and crystallinity. Fig. 28a shows the XRD pattern for an ED Ni/Cu multilayer grown on a single-crystal Cu(1 0 0) substrate [81]. In this figure, the peak labeled as S(2 0 0) at about 50.5° belongs to the substrate and the peak labeled as ML(2 0 0) corresponds to the mean spacing of the atomic planes within the multilayer itself (main Bragg peak of the multilayer). These reflections from (2 0 0) planes of the multilayer and the Cu(1 0 0) substrate indicate that the multilayer structure is oriented in the [100] direction as its substrate. The satellite peaks show that the multilayer grown on a single-crystal substrate has good periodicity. The modulation wavelength (K) was calculated to be (4.28 ± 0.32) nm from the first-order satellite peaks of the XRD pattern in Fig. 28a which agrees fairly well with the nominal bilayer thickness of the multilayer. Fig. 28b

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Fig. 28. (a) XRD pattern of an ED multilayer [Ni(2.0 nm)/Cu(2.0 nm)]150 grown on a single-crystal Cu(1 0 0) substrate. Cu Ka radiation was used to record the pattern on the sample being on its substrate. The satellite peaks apparent in the XRD pattern up to third order are labeled as (+1) and (1), (+2) and (2) as well as (3); (b) XRD pattern of an ED multilayer [Ni(3.0 nm)/ Cu(2.0 nm)]50 grown on (1 0 0)-textured polycrystalline Cu, recorded after removing the substrate. Layer thicknesses are nominal; for more details, see text. Note the logarithmic intensity scale. Reprinted from Ref. [81] with permission of Springer Science + Business Media.

shows the XRD pattern of another ED Ni/Cu multilayer grown on a (1 0 0)-textured polycrystalline Cu substrate [81]. In this case, the main Bragg reflections of the multilayer observed are the peaks labeled as ML(1 1 1) and ML(2 0 0). The ML(2 0 0) peak is stronger than the ML(1 1 1) peak. This shows that the multilayer has the same strong (1 0 0) texture as its polycrystalline Cu substrate. A reduction of the structural quality of the multilayer grown on the polycrystalline substrate is indicated by the fact that only first-order satellite peaks appear around the ML(2 0 0) Bragg peak. The modulation wavelength was found to be (5.34 ± 0.37) nm which agrees well with the nominal bilayer thickness of 5 nm of the given multilayer. Another example of the substrate influence on multilayer structure is demonstrated by Fig. 29 [17]. These XRD patterns also reveal the deleterious effect of additives on the structural quality (satellite reflections and line broadening). The correlation between reduced structural quality and diminished GMR will be discussed in Section 5.2.3.3. It was also found that the deposition mode can affect the multilayer texture and structure. Polycrystalline Cu(1 0 0) and Cu(1 1 0) substrates were used for electrodepositing Co–Ni–Cu/Cu multilayers in G/G and G/P modes [50] and, depending on the combination of deposition mode and substrate texture, the multilayer texture was found to be either (1 0 0) or (1 1 1) by cross-sectional TEM investigations. At the same time, the width of the growing columns varied in a wide range (from 10 to 2000 nm) with deposition parameters; sometimes even grains had diameters of the order of a few tens of nanometers only in any direction. In any case, this study revealed that the size of the GMR or even the absence of GMR was strongly correlated with the observed column/grain size. Similar correlations were reported also for ED Ni–Cu/Cu multilayers [29] from a combined XRD and TEM study when the layer thicknesses were varied, resulting in various grain and column sizes. In a detailed structural study [128] by XRD on ED Co–Cu/Cu multilayers, the importance of substrate roughness was revealed: whereas a pronounced fcc-(1 1 1) texture was observed in all investigated multilayers produced by either P/P, G/P or G/G mode, there was a strong difference

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Fig. 29. Overview of X-ray diffraction data for ED Co–Cu/Cu multilayers grown on (l 0 0)-oriented Cu substrates (a–c) and on (1 1 1)-oriented Au substrates (d and e). In the respective plating solutions, an additive-free sulfate bath was used for curves (a) and (d), whereas thiourea was added to the bath for curves (b) and (e) and brigthener Triton X-100 for curve (c). The Au(1 1 1) peak position and the positions of the (1 1 1) and (2 0 0) main multilayer peaks have been marked by arrows. Reprinted from Ref. [17] with permission of Elsevier.

in the degree of texture and appearance or absence of satellite reflections, depending on deposition mode and substrate roughness. The results suggested that a smooth Si/Ta/Cu substrate with evaporated Ta and Cu layers which provide a well-defined fcc starting structure is a favorable environment for a coherent quasi-epitaxial growth of Co–Cu/Cu multilayers, as opposed to a rough (mechanically polished) Ti substrate where the initial nucleation process results in a less uniform texture. Furthermore, for Ti substrates on which there is no opportunity of an epitaxial multilayer growth, it was also found [243] that the first few bilayers of the ED multilayer usually show a very disordered structure with very small crystallites. After achieving a critical thickness that is a function of all deposition conditions (electrolyte concentrations, temperature, electric parameters of the deposition), the self-organizing nature of the deposition leads to the formation of a textured and layered deposit with fairly large crystals. This self-organizing phenomenon is so pronounced that, for total multilayer thicknesses above a certain value, even an amorphous ribbon substrate can be used to produce ED multilayers [121] with GMR properties comparable to the cases when using crystalline substrates. It was already noticed above that the observation of GMR in FM/NM multilayers imposes some requirement on the structural quality of multilayer. Usually, the observation of multilayer satellite reflections is considered as necessary for high GMR, but apparently there is no such a strict correlation. In a study of Kubota et al. [189] on PD Co/Cu multilayers, the usual high GMR and its oscillatory behavior with Cu spacer thickness was observed, although the XRD patterns of the same multilayers exhibited hardly visible satellites only. It was also observed by the present authors on a series of ED Co/Cu multilayers that whereas clear XRD satellites were observed for Cu thicknesses from 2.0 to 4.0 nm [255] and an enhancement of the GMR was obtained with increasing Cu thickness, the absence of satellites for dCu > 4.0 nm resulted in a fairly small decrease of the observed GMR only [147]. Evidently, more systematic studies are still needed in order to reveal which microstructural features are really decisive in affecting the GMR magnitude. The interface quality is certainly one of these factors and efforts should be made to get further information on the interfaces in ED multilayers by appropriate diffraction techniques.

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5. Critical evaluation of GMR results reported on ED multilayer films 5.1. Introduction 5.1.1. Overview of systems investigated and preparation methods applied A complete bibliography of papers which report on MR characteristics of ED multilayer films exhibiting GMR behavior is provided in the list of references [9–149] whereby Refs. [26,65,73,81,125,126,133,137] constitute former reviews of this research area. It should be noted, however, that it was only the review by Schwarzacher and Lashmore [26] which considered and discussed in detail the results of all earlier reported GMR studies on these specific materials. Other reviews have mainly concentrated on general aspects of preparation of ED multilayer films exhibiting GMR behavior [81,133,137] or briefly summarized the progress achieved and listed the papers in the field [125,126]. By considering that the review by Schwarzacher and Lashmore [26] was published in 1996 and more than 80% of the papers reporting on GMR results in ED multilayer films were published thereafter, it is the aim of the present review to give a critical evaluation of the results of all papers which reported results in this field [9–149]. In view of the progress in ED technology and in understanding both electrochemical processes governing multilayer formation and physical mechanisms responsible for GMR in imperfect multilayered structures, a reconsideration of reported results already discussed in the earlier review [26] is inevitable also in the present overview. Most of the element combinations which were shown to exhibit GMR in magnetic/non-magnetic multilayers prepared by physical deposition methods are accessible also to the electrodeposition technique. Up to now, there have been reports on the GMR characteristics of electrodeposited Co–Cu/Cu, Ni–Cu/Cu, Co–Ag/Ag, Co–Au/Au, Co–Ru/Ru, Co–Ni–Cu/Cu, Fe–Ni–Cu/Cu, Fe–Co–Cu/Cu, Co(–Cu)–Zn/Cu and Fe–Co–Ni–Cu/Cu multilayer films. The discussion of reported results will follow this sequence and detailed references to individual systems will be listed in the corresponding sections. The majority of GMR studies on ED multilayer films were carried out on samples prepared by twopulse plating from a single bath containing both the magnetic and non-magnetic species. There are a few reports only on GMR of ED multilayer films prepared (i) by using a single-bath technique with a pulse sequence consisting of more than two pulses including an off-time between the deposition pulses [22,48,110,144] or of a pulse train [131,135], (ii) by a dual-bath technique [71,75,93,106,117,120,127,140,143] and (iii) by injecting the magnetic and non-magnetic species separately into a supporting flow electrolyte [107]. As discussed in Section 3.2.1, in single-bath multilayers the magnetic layers unavoidably contain some amount of the non-magnetic spacer layer material. Therefore, for the sake of generality, the composition of the magnetic layer in multilayers obtained by the single-bath deposition always will be specified by including also the NM component (e.g., Co–Cu/Cu) even if the contamination of the magnetic layer can be suppressed down to below the 1 at.% level under appropriate deposition conditions (high ionic ratio of the magnetic and non-magnetic elements in the bath and the application of a high-current or high-potential pulse for the deposition of the magnetic layer). In discussing the results of individual reports from the literature, we have to invoke the knowledge of the current level of understanding as summarized in former sections. However, even if we restrict ourselves to multilayers with a given element combination, the richness of data unavoidably leads to a diversity of preparation conditions and, thus, to a diversity of GMR characteristics observed. These GMR data are sometimes even contradictory and reproducibility from work to work is often less than satisfactory. Therefore, in order not to get lost in the turmoil of details, for a given multilayer system we shall attempt to group the results according to various aspects of the deposition conditions and by considering all related reports of one laboratory simultaneously.

5.1.2. Guidelines in evaluating the experimental results on GMR in ED multilayer films Before starting the evaluation of specific multilayer systems, the basic viewpoints of assessment are listed here first. A very important issue is the pulse control mode. As discussed in Section 3.2.2, under the application of the G/G pulse combination, due to the so-called exchange reaction (spontaneous replacement of previously deposited Fe, Co or Ni atoms by Cu atoms from the electrolyte), the

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individual layer thicknesses will be definitely not those set on the computer controlling the current pulses but the magnetic layer will be thinner and the non-magnetic layer thicker than the preset values. For the P/P and G/P deposition modes, the actual layer thicknesses will be equal to the preset (nominal) values only if the deposition potential of the more noble (non-magnetic) component is properly optimized (see Section 3.6 and Refs. [129,231]) in order to avoid a dissolution of the previously deposited magnetic layer (another constraint for the optimization is to prevent the codeposition of the magnetic element into the spacer layer). The layer thickness changes due to the exchange reaction or magnetic metal dissolution can be estimated by direct chemical analysis of the overall multilayer sample [99,103,128]. A direct evidence for such layer thickness changes which can be as high as 1.4 nm derived from the chemical analysis data could also be given by measuring cross-sectional elemental maps for ED Co/Cu multilayers by a high-resolution TEM [128]. The correct values of the individual layer thicknesses become important when attempting to investigate the dependence of GMR on the thickness of the magnetic and non-magnetic layers. The latter dependence has been especially often addressed in GMR studies of ED multilayer films and the reports are rather contradictory as to whether an oscillatory GMR behavior exists in such systems, although this distinct feature is well-documented for many multilayer systems prepared by PD methods. This issue is further complicated by the fact that the field dependence of the magnetoresistance (namely, the frequently observed non-saturating behavior of the MR(H) curves up to magnetic fields as high as several 10 kOe) in ED multilayers indicates the presence of superparamagnetic regions (see Section 2.4.4) which give rise to a GMR contribution (GMRSPM). At the end of Section 2.4.4, some factors were mentioned which may lead to the formation of SPM regions during ED multilayer preparation and, thus, to a significant increase of the GMRSPM fraction. If these factors are not controlled to a degree that the GMRSPM contribution is strongly suppressed, then the MR(H) data should be decomposed to separate the GMRSPM term and to deduce the GMRFM term since this latter term is only related to the oscillatory GMR. This is because such an oscillation is ultimately connected with an alternating FM and AF exchange coupling between adjacent magnetic layers as a function of the NM spacer layer thickness whereas we have no definite evidence yet for the presence of an AF coupling in ED multilayers at all. These above discussed features are strongly dependent on the microstructure and defect structure of the multilayer. These latter, on the other hand, are governed by other details of the deposition process such as (i) the film growth texture (preferred crystallographic orientation) mainly controlled by the substrate material and its crystal orientation, (ii) the electrolyte composition and agitation, the pH and the presence of additives, (iii) the magnetic layer contamination by the non-magnetic element and (iv) the cathode position and orientation. The latter two features mentioned in item (iv) influence, together with the cell geometry, also the lateral film homogeneity over the cathode surface area. We should keep in mind all these aspects when evaluating reported data and should consider also other parameters available on the same multilayer, with special reference to structural data. It is to be noted finally that a sizeable GMR could often be observed even under non-optimal deposition conditions; in such cases, we should proceed with special care when assessing the GMR data. Anyway, it should be established that the GMR magnitude in most ED multilayer films has still remained well below the values reported for the corresponding PD multilayers. The probable reason for the inferior performance of ED multilayer films is that, in spite of the significant progress in the preparation technology (see Section 3), we still cannot control some fine details of the ED process. Nevertheless, relatively high GMR values could be achieved also in ED multilayer films. It should be made, however, clear already at the outset that if we intend to make a meaningful comparison with the GMR of corresponding PD multilayers, we need to establish as accurately as possible the actual individual layer thicknesses of the multilayers in question and also the magnetic field at which the given GMR value was achieved (or it would be even better to know the full MR(H) dependence). It may have already turned out from the preceding discussions that layer thicknesses and GMR magnitude (and, also, the field dependence of the GMR) are very strongly interrelated for various reasons. An overall view of the expected GMR behavior for various layer thickness ranges will be made later in Section 5.2.3.1 when discussing results on the magnetoresistance and magnetic properties of the most widely studied ED Co–Cu/Cu multilayer system.

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With the above description of evaluation guidelines, we are now armored to discuss the results of GMR studies reported on ED multilayer films on a general basis. The presentation and evaluation of the GMR results on ED multilayer films will be made by following a sequence of increasing number of element components. The discussion here in Section 5 will be limited to multilayer films comprised of numerous repetitions of a FM/NM bilayer unit whereas a summary of results on ED sandwich structures (such as, e.g., spin-valves or pseudo spin-valves) [46,47,58,64,73,76,79,91,140,143] will be presented in Section 6. 5.2. Co–Cu/Cu multilayer films 5.2.1. General overview of deposition conditions About half of the papers reported on a GMR study of ED multilayer films was dealing with Co–Cu/Cu multilayers [15,17,21,22,24,25,27,30–36,41–43,45,49,51,52,55,57,60–63,66,68,74,75,77,78,80,84,88, 90,92–95,98–101,103,107,109,113–116,120,121,123,124,127–136,139,141,144–147,149]. The driving force for this extended research effort was the fact that, among PD multilayers, the Co/Cu system exhibited the largest GMR being as high as about 50% at room temperature [177,178,189]. Unfortunately, the GMR magnitude of ED Co–Cu/Cu multilayer films has hardly reached even only 20% (apart from some scarce reports [15,120,127] the results of which could not yet be reproduced by other laboratories). Due to the differences in the electrochemical behavior of Co and Cu as well as the immiscibility of the two elements, the multilayer growth process appears to be especially unfavorable for achieving appropriate GMR characteristics in this system by electrodeposition. As will be discussed in detail in Section 5.2.3.1, the thicknesses of the magnetic and non-magnetic layers play a decisive role in determining the expected GMR behavior. On the other hand, the layer thicknesses are controlled by the deposition conditions, namely by the applied pulse amplitudes and lengths as well as by the pulse control modes of the individual pulses (G or P) as pointed out in Section 3. For the preparation of Co–Cu/Cu multilayers, G/G [21,27,31,32,43,60– 62,68,75,77,78,90,94,99,103,116,123,124,128,131,133,135], G/P [99,109,115,116,121,128,132,139, 147,149] or P/P [15,17,22,24,25,30,33–36,41,42,49,52,54, 55,57,63,66,74,80,84,88,92,93,95,98– 101,107,113,114,116,120,127–130,134,136,141,145,146] pulse combinations were applied. The deposition of ED Co–Cu/Cu multilayer films was carried out in most cases by the conventional two-pulse plating method utilizing square-shaped pulses from a single bath; exceptions from this practice will be discussed at the relevant places. As was emphasized above and in Section 3, the actual layer thicknesses can be considered equal to the preset values during deposition only if the exchange reaction or the Co-dissolution are successfully suppressed during the non-magnetic layer deposition cycle (in addition, in view of getting a good GMR, the Co incorporation into the Cu layer should also be avoided). These requirements can be simultaneously fulfilled only if the Cu layer is deposited under potentiostatic control at an optimized potential. In this respect, the choice of the Co deposition control mode (G or P) and the chosen current or potential amplitude is usually of negligible influence only, these parameters affect mainly the amount of Cu incorporated into the Co-layer. We intend to focus first on selecting those works where the above specified conditions are fulfilled and evaluate their GMR results. Since it has been well documented that a significant GMR can be achieved even under non-optimized deposition conditions, we shall consider also these results and compare them to those obtained on multilayers prepared under optimized conditions. 5.2.2. Optimization of deposition parameters It must be clear from Section 3.2.2 that in the case of pulse-plating from a single bath, the various electrochemical processes are controlled differently by the various deposition modes (G/G, G/P or P/P). Since in electrodeposition the main control parameter is the cathode potential, a proper control of the deposition process can only be achieved by controlling the cathode potential during the deposition of the more noble non-magnetic metal. This can be realized in the G/P and P/P modes whereas in the G/G mode, the cathode potential is not controlled. In this section, we shall restrict discussion to the optimization of the more noble layer deposition potential (ECu) in the G/P and P/P modes. A discussion of

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GMR results on ED Co–Cu multilayers prepared from a single bath in the G/G mode will follow in Section 5.2.4. As discussed in Section 3, the Cu deposition potential crucially influences the multilayer formation in ED Co–Cu/Cu multilayer films and the optimum value of the Cu deposition potential may somewhat depend on the bath composition. In the original reports, the Cu deposition potential values used were specified with respect to various reference electrodes (see Table 1). In order to have a more transparent overview of the Cu deposition potentials (ECu) applied for the preparation of ED multilayers with GMR, we have transformed these reported ECu values to standardized ECu(SHE) values by referencing to the saturated hydrogen electrode (SHE). The optimization of the Cu deposition potential was carried out mainly for the most investigated sulfate bath only. The optimization problem was thoroughly approached in Ref. [231] (see also Section 3.6) where it could be shown that the commonly applied cyclic voltammetry cannot provide a reliable basis for determining the optimum potential value for the more-noble component deposition by pulse plating but it should rather be established under pulsed conditions corresponding to the actual multilayer deposition. In Ref. [231], an analysis of the current transient during the more-noble component deposition, which occurs after switching from the high-potential/high-current pulse of the less-noble (magnetic) component, was suggested for the optimization process. By using this procedure, it was found [109] that for a bath containing CoSO4 and CuSO4 only, ECu(SHE) = 0.38 V is an optimum potential ensuring a simultaneous suppression of both Co dissolution and incorporation of Co into the Cu layer. A very close value [ECu(SHE) = 0.36 V] was obtained from a sulfate bath containing also some amount of H3BO3 and (NH4)2SO4 [255]. An alternative approach was elaborated by Ghosh et al. [129], which relies on the weight change of the substrate by using an electrochemical quartz crystal microbalance and this method yielded a similar value, ECu(SHE) = 0.36 V, as the optimum deposition potential for ED Co–Cu/Cu multilayer preparation from a pure sulfate bath. Of course, the optimization procedure described in these works can be transferred to any other bath used for multilayer deposition. In most studies of GMR in ED Co–Cu/Cu multilayer films with potentiostatic control of the Cu layer, the Cu deposition potential was established from considerations based on the cyclic voltammograms or the basis of choice was not specified at all. An exception is the work of Lenczowski et al. [17] who choose ECu(SHE) = 0.34 V as an optimum value by taking the experimentally determined equilibrium potential for the Co ¡ Co2þ þ 2e reaction of their sulfate bath without Cu2+ ions. The applied ECu(SHE) values reported in previous other studies range from 0.56 V to 0.01 V for the sulfate based electrolytes and from 0.16 V to +0.05 V for the sulfamate/sulfate electrolytes. In case the Cu deposition potential is chosen more positive than the optimum ECu value for the actual electrolyte, some degree of the dissolution of the previously deposited Co-rich magnetic layer occurs. Since the layer thickness is usually controlled on the basis of the amount of charge passing through the bath, the dissolved Co (negative charge) is balanced by an excess Cu (positive charge) to reach the desired total charge preset for achieving the nominal thickness. This results in a reduced Co layer thickness and in an increased Cu layer thickness. On the other hand, if the Cu deposition potential is chosen more negative than the optimum ECu value for the actual electrolyte, then incorporation of Co into the non-magnetic Cu layer occurs. Therefore, we shall consider GMR results on the basis of the applied Cu deposition potential since only in this case we can have the actual individual layer thicknesses equal to the nominal values (and, also to ensure that the Cu layers are free of Co impurities). In case the Cu deposition is not the optimum value, one can estimate the actual layer thicknesses from a measurement of the overall chemical composition [17,99,128,263]. In this manner, Lenczowski et al. [17] was able to show that although their Cu deposition potential of ECu(SHE) = 0.34 V enabled a slight dissolution of Co but this dissolution was less than 0.5 nm (i.e., about two monolayers) as it can be established from Fig. 1 of Ref. [17]. Three main types of bath were used for preparing ED Co–Cu/Cu multilayers films exhibiting GMR characteristics. Most ED Co–Cu/Cu multilayer films for GMR studies have been deposited from a bath containing CoSO4 and CuSO4 only, eventually with some other rather neutral (buffering) agents. There were also several studies by using such sulfate-based baths with some

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additives. There were a few reports only where Co-sulfamate Co(NH2SO3)2 was used instead of the Co-sulfate component and in one work, sulfamic acid (NH2SO3H) having an equivalent effect as the Co-sulfamate was added to the sulfate bath. The discussion of the reported GMR results will be carried out separately for multilayers prepared from these three main bath types in the G/P and P/P modes.

5.2.3. GMR of ED Co–Cu/Cu multilayer films prepared in P/P or G/P mode from a single bath 5.2.3.1. Deposition with optimized Cu deposition potential from sulfate bath. Table 2 summarizes the detailed preparation conditions from reports on ED Co–Cu/Cu multilayer films for GMR studies which were deposited from a bath containing CoSO4 and CuSO4 only, eventually with some other fairly neutral (buffering) agents (X) mainly for the purpose of controlling the bath pH value. For these simple electrolytes, ECu(SHE) = 0.36 V to 0.38 V was established as an optimum Cu deposition potential in Refs. [109,121,129,141,255]. Therefore, for these multilayers, the nominal layer thicknesses correspond well to the actual values. As noted above, Lenczowski et al. [17] has also chosen a close value for ECu and for this case the Co dissolution could be estimated to cause a layer thickness change of less than 0.5 nm. We shall, therefore, treat the results of all these works as obtained on ED Co–Cu/Cu multilayer films prepared under optimized conditions. The first well-documented work on GMR for ED Co–Cu multilayers from among the latter reports was carried out by Lenczowski et al. [17] who performed a systematic study as a function of both layer thicknesses. These authors used Si(1 0 0)/50nm–Au(1 1 1) and Si(1 0 0)/20nm–Cu(1 0 0) substrates whereby the metal layers on top of the Si wafer were obtained by evaporation and sputtering, respectively. The bilayer repeat number was fixed at 50 and the total multilayer thickness ranged from about 100 nm to 350 nm. The Cu-content in the magnetic layers was measured to be about 4 at.%. XRD revealed that the Co–Cu/Cu multilayers have an fcc structure and inherited the texture of the seed metal layers. The XRD patterns exhibited clearly resolved first-order satellite reflections for a (1 1 1)-ori-

Table 2 Deposition parameters of ED Co–Cu/Cu multilayer films and sandwich structures with GMR which were obtained under P/P or G/P control from a single sulfate bath containing buffering agents X only with components CoSO4 + CuSO4 + X where X stands for H3BO3, (NH4)2SO4 and/or NH4OH. The cathode potentials applied for deposition are referenced to the standard hydrogen electrode (SHE), see Table 1. In works above the horizontal dashed line (for both multilayers and sandwich structures), the Cu deposition potential used was fairly close to the optimum value. Authors and Refs. in [] Multilayer films Liu et al. [109,121] Bakonyi and cow. [147,149] Ghosh and cow. [129,141] Lenczowski et al. [17] Dinia and cow. [30,41]

X

pH

(NH4)2SO4 + H3BO3 H3BO3 H3BO3

2.7 4–4.5 4.0

Co2+/Cu2+ ionic ratio

ECo(SHE)/ECu(SHE) or jCo/ECu(SHE)

5–200 53 250 176 200

70.5 mA cm2/0.38 V 70 mA/cm2/0.36 V 1.16 V/0.36 V 0.96 V to 1.36 V/0.34 V 0.96 V/0.34 V

------------------------------------------------------------------------------------Zhang et al. [63,101] H3BO3 + NH4OH 5.6 190 0.86 V/0.32 V 3.5 190 0.77 V/0.31 V Shima and cow. [45,55,57,80,88] H3BO3 Weihnacht and cow. [99,115,128,132,139] 1.6–2.8 40 ECo, jCo/0.01 to 0.26 V Pandya and cow. [130,136] 4.0–5.4 40 0.96 V/0.26 V Jyoko and cow. [35] 100 0.90 V/0.20 V 2.0 50 or 100 1.56 V/floatinga/0.02 V Lashmore and Hua [22] H3BO3 Yan et al. [92] 200 Potentials not specified Sandwich structures 3.3 176 0.90 V/0.35 V Attenborough et al. [46,47,64] H3BO3 -----------------------------------------------------------------------------3.5 190 0.77 V/0.31 V Shima et al. [58] H3BO3 Pasa and cow. [76,79] H3BO3 4.0–4.5 176 0.91 V/0.16 V a

Between the Co and Cu deposition pulses, the electrodeposition cell circuit was opened for 2 s in each bilayer period.

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ented Co(1 nm)/Cu(4 nm) multilayer on Au and satellite reflections up to the third order for a (1 0 0)oriented Co(1.5 nm)/Cu(4 nm) multilayer on Cu (see Fig. 29). Liu et al. [109] investigated the layer thickness dependence of both the magnetic properties and the magnetoresistance of fcc (1 1 1)-oriented ED Co–Cu/Cu multilayers obtained on polished Ti sheets from which the multilayers having a total constant thickness of about 1700 nm could be peeled off mechanically. In this work, the Cu-content of the magnetic layer was about 5 at.%. XRD could not reveal the presence of satellite reflections on the few samples selected for structural investigations, probably due to the use of a rather rough substrate. In a subsequent work, Liu et al. [121] extended the investigations for ED Co–Cu/Cu multilayers with varying amount of Cu in the magnetic layers which was achieved by changing the Co2+/Cu2+ ionic ratio in the electrolytes. In both studies, a tubular cell suggested in Ref. [99] was used which ensured a lateral homogeneity of the deposits by avoiding so-called ‘‘edge effects”. An important characteristic of the works by Lenczowski et al. [17] and Liu et al. [109,121] was that the cathode was placed facing upward which had the beneficial effect of avoiding the variation of the GMR properties along the cathode surface area due to a spontaneous convection the effect on GMR of which was demonstrated [77,82] in the commonly used arrangement of a vertical cathode. Actually, the cathode position and orientation is usually not mentioned in the reports and, therefore, we shall assume in such cases that a vertical cathode with ‘‘open” deposition geometry not preventing edge effects was applied. This must have been the case in the works by Ghosh and coworkers as well [129,141] who reported on a detailed study of GMR and magnetic properties as a function of the Co layer thickness. These authors used a Si(1 1 1)/Ti(20 nm)/Cu(20 nm) substrate with sputtered Ti adhesive and Cu seed layers. The bilayer number was fixed at 50 so that the total multilayer thickness ranged from about 200 nm to 700 nm. Although not explicitly specified, the Cu content of the magnetic layer was certainly fairly low due to the high magnetic/non-magnetic ionic ratio in the bath. XRD could reveal satellite reflections for Co layers at least as thick as 8 nm only and a pronounced fcc(1 1 1) texture was observed. The present authors [147] investigated the Cu layer thickness dependence of the GMR behavior and magnetic properties of optimized ED Co–Cu/Cu multilayers obtained from a sulfate bath on Si(1 0 0)/ Cr(5 nm)/Cu(20 nm) substrates with evaporated metallic underlayers. The total multilayer thickness was constant at about 450 nm. It was found in a separate study [255] that the Cu-content of the magnetic layer was very low (ca. 0.6 at.%) and that there was a systematic evolution of the structure with Cu spacer layer thickness. For Cu spacers thinner than about 2 nm, a small hcp-Co fraction was also present and there were no satellite reflections of the dominating fcc-(1 1 1) multilayer XRD peak. The hcp-Co fraction can be taken as an indication for a discontinuity of the Cu layers, enabling the coherent growth of Co to a thickness where the fcc phase is not stable any longer and hcp-Co appears. On the other hand, above 2 nm Cu layer thickness the hcp-Co fraction disappeared and first order satellites could be identified up to about 4 nm Cu layer thickness. A recent more detailed XRD study [149] on similarly prepared ED Co/Cu multilayers revealed microstructural features supporting the structural information deduced in Ref. [255]. We shall now summarize the results of magnetoresistance and magnetic measurements reported in Refs. [17,109,121,129,141,147,149] on ED Co–Cu multilayers obtained from sulfate baths with optimized Cu-deposition potential. First, the results of Liu et al. [109] will be presented in detail. Fig. 30 shows the MR(H) and M(H) curves for two multilayers with thick (5.7 nm) and thin (1.1 nm) Co layers in both cases with thick Cu layers (dCu = 2.5 nm). For these multilayers, we can observe a GMR behavior in that both the LMR and the TMR components are negative for all fields investigated (cf. Fig. 3). A splitting of the MR(H) curves can be observed for both multilayers (Fig. 30a and b). The MR peak position values (Hp) correlate well with the coercive fields (Hc) deduced from the corresponding hysteresis loops of these samples (Fig. 30c). For the multilayer Co(5.7 nm)/Cu(2.5 nm), the MR peaks are relatively narrow and the MR value reached at around 1 kOe changed only very little up to 8 kOe. According to Fig. 30c, the magnetization certainly also reaches technical saturation by around 1 kOe for this sample. Furthermore, there is about 1% difference between the LMR and TMR values that can be ascribed to an AMR contribution (cf. Fig. 3) to the observed MR due to the relatively large Co layer thickness (5.7 nm) in this multilayer.

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Fig. 30. (a) Longitudinal and transverse MR(H) curves of an ED Co–Cu/Cu multilayer with thick magnetic and thick nonmagnetic layers. (b) Longitudinal and transverse MR(H) curves of an ED Co–Cu/Cu multilayer with thin magnetic and thick nonmagnetic layers. (c) Magnetic hysteresis loops of ED Co–Cu/Cu multilayers with thick and thin magnetic layers at the same thickness of the non-magnetic layers (dCu = 2.5 nm). Reprinted from Ref. [109] with permission of Elsevier.

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Fig. 31 shows the MR(H) and M(H) curves from Ref. [109] for two multilayers with thick (5.7 nm) and thin (1.1 nm) Co layers, in both cases with thin Cu layers (dCu = 1.1 nm). For the multilayer Co(5.7 nm)/Cu(1.1 nm) shown in Fig. 31a, we can observe an AMR behavior since LMR > 0 and TMR < 0 (cf. Fig. 3) whereas for the multilayer Co(1.1 nm)/Cu(1.1 nm) shown in Fig. 31b, a GMR behavior. However, the MR(H) curves of the latter multilayer do not saturate at 2 kOe and MR saturation was not reached even at 8 kOe. As Fig. 31c indicates, the multilayer with thick Co and thin Cu layers exhibits a very low coercive field of about 30 Oe. The results on the Co layer thickness dependence of the magnetoresistance MR(H = 2 kOe) and the coercive field for various Cu layer thicknesses in ED Co–Cu/Cu multilayers from the work of Liu et al. [109] are presented in Fig. 32 and the same data are shown in Fig. 33 on the Cu layer thickness dependence for various Co layer thicknesses. A fairly systematic trend can be observed and in each case, a monotonous change of the magnitude of the magnetoresistance and the coercive field is obtained. The Hc and Hp values show a good correlation with each other. Although in the study of Liu et al. [109] the MR(H) curves were not yet separated quantitatively into FM and SPM contributions according to the method described in Ref. [103], the shape of the measured MR(H) curves unambiguously indicated the relative importance of the GMRFM and GMRSPM terms. Therefore, it was possible to draw some important qualitative conclusions from this comprehensive magnetoresistance and magnetic study for the evolution of GMR and microstructure with layer thicknesses in ED Co–Cu/Cu multilayers which will be discussed now with the help of the schematic diagrams shown in Fig. 34. With reference to Section 2.4.1, we consider first a kind of an ideal FM/NM multilayer in which both the FM and the NM layers are continuous and of uniform thickness (upper right panel in Fig. 34). This condition can be relatively easily realized by any deposition technique provided the layer thicknesses are sufficiently large as for example the values indicated in the panel. Under such conditions, the magnetic layers usually exhibit regular FM behavior without any SPM contribution. Even at fairly large (say 10 nm) NM spacer layer thicknesses, there may still be some exchange coupling between adjacent magnetic layers. A dominating AF coupling (AP alignment in zero external magnetic field) gives rise to a conventional GMR contribution (GMRFM). Due to the large bilayer length, this contribution may be so small that it may be overwhelmed by the AMR effect in the thick magnetic layers so finally the observed magnetoresistance may show AMR features only. On the other hand, if the interlayer coupling is of FM type (P alignment in zero field), no GMR effect arises and we definitely get AMR only. In case the thick spacer layer does not mediate any coupling between adjacent magnetic layers, a possible magnetic state will be a distribution of layer magnetization orientations in the layer planes if each layer constitutes a single domain. The magnetization orientation of the individual layers is determined by the local anisotropies, mainly of magnetocrystalline origin, which may have several preferred orientations in the layer planes. In such a case, a partial AP alignment of adjacent layer magnetizations results in a GMRFM contribution. The field evolution of the magnetization and, thus, of the magnetoresistance is governed by the laws of rotational remagnetization as the layer magnetizations approach to saturation against the anisotropies present; some hysteresis of the MR(H) curves may appear. For uncoupled magnetic layers, a multidomain state of each layer can also be envisioned [190,265] and in zero magnetic field there may be here also adjacent domains of neighboring magnetic layers with non-aligned magnetizations. These non-aligned domain pairs will, of course, also contribute a GMRFM term and in such cases the field dependence MR(H) is governed through the magnetization process via domain wall motion. The MR(H) curves are then characterized by a distinct split-peak behavior whereby the peak positions (Hp) are symmetrically spaced around H = 0 and appear roughly around the coercive field (Hc). Along these lines, we may say that the MR(H) curves in Fig. 30a correspond approximately to the situation sketched in the upper right panel of Fig. 34 where allowance was made also for a very small SPM contribution. As a next step, we consider a multilayer in which the magnetic layer thickness is getting reduced from the previously discussed large value. It has been a general experience in magnetic/non-magnetic multilayer preparation that down to a critical thickness which is about 1 nm, the magnetic layer typically exhibits a predominantly FM state with some SPM regions also present; the amount of the latter may slightly increase with decreasing Co layer thickness. This implies that the magnetic layer material

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Fig. 31. (a) Longitudinal and transverse MR(H) curves of an ED Co–Cu/Cu multilayer with thick magnetic and thin non-magnetic layers. (b) Longitudinal and transverse MR(H) curves of an ED Co–Cu/Cu multilayer with thin magnetic and thin non-magnetic layers. (c) Magnetic hysteresis loops of ED Co–Cu/Cu multilayers with thick and thin magnetic layers at the same thickness of the non-magnetic layers (dCu = 1.1 nm). Reprinted from Ref. [109] with permission of Elsevier.

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Fig. 32. (a) Evolution of the MR (H = 2 kOe) of ED Co–Cu/Cu multilayers with magnetic layer thickness for both thin and thick non-magnetic layers. (b) Evolution of the coercive field Hc and MR peak position Hp of ED Co–Cu/Cu multilayers with magnetic layer thickness for both thin and thick non-magnetic layers. The triangle (D) indicates the coercive field of a single Co thin film electrodeposited on Cu [264]. Reprinted from Ref. [109] with permission of Elsevier.

deposited on top of the non-magnetic layer consists of large percolated FM areas and only a small fraction of the magnetic material is in the form of magnetically separated islands (SPM regions). This means that down to this thickness range the basic GMR mechanism does not change (large GMRFM and small GMRSPM terms). As to the magnetic characteristics, if the magnetization process proceeds by rotation, no change is expected, provided the anisotropies remain the same as for thicker magnetic layers. If the magnetization mechanism is via domain wall motion, the coercive force increases approximately inversely proportional to the layer thickness. This situation corresponds fairly well to the MR(H) curves shown in Fig. 30b and to the lower right panel in Fig. 34. In the latter case, allowance was made for a somewhat larger SPM contribution in agreement with the slightly higher saturation fields in Fig. 30b in comparison with Fig. 30a. We should also consider the case when reducing the spacer layer thickness while keeping the magnetic layer thickness at a high value where it is still completely ferromagnetic (cf. upper left panel in

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Fig. 33. (a) Evolution of the MR (H = 2 kOe) values of ED Co–Cu/Cu multilayers with non-magnetic layer thickness for magnetic layers with various thicknesses. (b) Evolution of the coercive field Hc and MR peak position Hp of ED Co–Cu/Cu multilayers with non-magnetic layer thickness for magnetic layers with various thicknesses. Reprinted from Ref. [109] with permission of Elsevier.

Fig. 34). Starting from a thick spacer layer, we may observe a GMR oscillation if an AF interlayer coupling dominates. In the absence of a significant AF coupling at any spacer thickness as discussed above, we may still have a GMRFM contribution due to non-aligned magnetizations in adjacent regions. This GMR prevails or remains unchanged until we finally reach a spacer thickness range where the nonmagnetic spacer becomes not completely continuous but will contain discontinuities in the form of tiny pin-holes, i.e., not being able to separate completely the adjacent magnetic layers. If this happens, a direct FM coupling develops between the adjacent magnetic layers via the pin-holes, should they be just a few atoms wide only. The P alignment of some regions of adjacent magnetic layers via pin-holes will definitely lead to a reduction of the observed GMR. With decreasing spacer thickness, the density of pin-holes can be expected to increase, leading finally to a reduced GMR towards thinner spacer layers. If the FM-coupled regions in the magnetic layers become dominant, the GMR may completely disappear and we can observe an AMR effect only, a characteristic of bulk ferromagnetic metals (cf. the MR(H) curve in Fig. 31a and the data points with LMR > 0 in Figs. 32a and 33a). This is a consequence of

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Hc = 29 Oe

MR

Hc = 130 Oe longitudinal

MR H GMRSPM

AMR H

dCo (nm)

Co

GMRFM

Cu

Co Cu Co (5.7) / Cu (2.5)

transverse

Co (5.7) / Cu (1.1)

Hc = 140 Oe

Hc = 276 Oe

MR

MR

H

H

GMRFM

Co

Cu

Co (1.1) / Cu (1.1)

GMRSPM

Co

GMRSPM

Cu

GMRFM

Co (1.1) / Cu (2.5)

dCu (nm) Fig. 34. Schematic view of the microstructural features and MR contributions for selected ED Co-Cu/Cu multilayers with possible combinations of thick and thin magnetic and non-magnetic layer thicknesses. Reprinted from Ref. [109] with permission of Elsevier.

the percolation between adjacent magnetic layers via the pinholes. The appearance of a pin-hole-mediated FM coupling is accompanied also by a reduction of the coercive field due to the approach to the bulk-like behavior, with Hc indeed reaching the bulk coercive force of the magnetic layer material for very thin spacer layers where already AMR dominates (cf. Fig. 31c). The critical spacer-layer thickness where the GMR converts to an AMR behavior certainly depends on the details of multilayer preparation and for ED multilayers, it is in the range of 1–2 nm. Whereas it is very hard to observe pin-holes directly by structural methods, XRD studies have already provided some indirect evidence for the presence of pin-holes in ED Co–Cu/Cu multilayers [149,255]. We should consider one more case which seems to occur when the spacer layer thickness is reduced for multilayers with fairly thin magnetic layers. Namely, as opposed to thin magnetic layers with thick spacer layers, a fragmentation of thin magnetic layers may occur in combination with thin non-magnetic layers. Apparently, the discontinuity of thin spacer layers induces a discontinuous growth of thin magnetic layers as well. As a result, a mixture of FM and SPM magnetic particles in a non-magnetic matrix is formed, with magnetic and MR characteristics somewhat similar to the combination of thick spacer layers and very thin magnetic layers as discussed above. All this means that for typically 1 nm thickness of both the magnetic and the non-magnetic layers, no well-defined layered structure occurs and the GMRSPM contribution will dominate (cf. Fig. 31b and lower left panel of Fig. 34). A basic conclusion that could be drawn from the work of Liu et al. [109] is that for ED Co–Cu/Cu multilayers with sufficiently thick (at least 1.5–2.5 nm) Cu spacer layers and with sufficiently thick Co layers (at least about 1 nm), the MR nearly saturates at magnetic fields around 2 kOe and this MR value can be mainly ascribed to the GMRFM term arising from spin-dependent scattering of electrons travelling through the spacer between to FM layers. For such a thick spacer layer, the GMR magnitude decreases with increasing Co layer thickness. For a constant Co layer thickness, the GMR increases with Cu layer thickness monotonously without any observable oscillatory behavior. A

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subsequent work by Liu et al. [121] in which the GMRFM term was properly established by removing the GMRSPM contribution gave a similar evolution of GMRFM with Cu layer thickness up to dCu = 3.5 nm. This is in good agreement with the former results of Lenczowski et al. [17] who observed a similarly low MR saturation field of about 2–3 kOe and a monotonous increase of the GMRFM term up to about 4.5 nm; however, beyond that thickness the GMR started to decrease again (Fig. 35) up to the highest Cu thickness investigated (6.5 nm). Recent works by the present authors [147,149] also revealed a monotonous GMR increase with spacer thickness up to about dCu = 4 nm. The spacer layer thickness dependence of GMRFM in ED Co–Cu/Cu multilayers will be discussed later in more detail in Section 5.2.7. Fig. 35 also reveals that for the Co–Cu/Cu multilayers studied by Lenczowski et al. [17], the GMR is larger for the case of Au(1 1 1) substrates in comparison with the corresponding multilayers on Cu(1 0 0) substrates although the qualitative dependence on Cu spacer thickness is very similar. Whereas XRD patterns [17] indicated better multilayer quality for multilayers on Cu(1 0 0) substrates (visible third-order satellites in contrast to visible first-order satellites only for multilayers on Au(1 1 1) substrates, cf. Fig. 29), the GMR performance is better for the Au(1 1 1) substrates (especially if we consider the differences in the shunting effects due to the different seed-layer thicknesses: 20 nm for Cu and 50 nm for Au). Apparently, the overall conditions (crystal orientation, lattice mismatch and surface roughness) for the growth of a more perfect multilayer are better fulfilled for the Au(1 1 1) substrates, at least as far as the magnitude of the GMR is concerned. Lenczowski et al. [17] have also found that the addition of the commonly applied brighteners thiourea and Triton X-100 to the sulfate bath had a deleterious effect on the multilayer structure and properties: (i) strong reduction or complete disappearance of XRD satellites (cf. Fig. 29), (ii) increase of zero-field electrical resistivity and (iii) reduction of the GMR at least by an order of magnitude as seen in Fig. 35b. The decrease of GMR for large Cu layer thicknesses as can be observed in Fig. 35 is a natural consequence of the increasing bilayer repeat period since then the number of magnetic/non-magnetic alternations per unit length along the layer thickness decreases and, in this manner, the primary source of spin-dependent scattering events responsible for the GMR is reduced.

Fig. 35. Dependence of the giant magnetoresistance (MR) on the thickness of the Cu spacer layer (dCu) for (a) (l 0 0)-oriented Co–Cu/Cu multilayers grown on Cu(1 0 0) substrates, and (b) (1 1 1)-oriented Co/Cu multilayers grown on Au(1 1 1) substrates. The magnetic layers had a thickness of 1.3 nm and the number of repeats was 50. The crosses in the bottom curve of (b) represent data of samples which were electrodeposited with thiourea added to the electrolyte. Note that these values have been multiplied by a factor of five. The full lines are guides to the eye. Reprinted from Ref. [17] with permission of Elsevier.

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The same mechanism governs the decrease of the GMRFM term with increasing Co layer thickness for ED Co–Cu/Cu multilayers as observed for Co thicknesses above about 1 nm (see Fig. 32a with data from Ref. [109]). Similar results were reported by Lenczowski et al. [17] and Chowdhury et al. [141] in this thickness range as shown in Figs. 36a and b. However, we can also see from Figs. 36a and b that, when reducing the Co layer thickness at a relatively thick Cu spacer, the monotonous GMR increase turns into a sudden decrease at about 1 nm in both latter studies. This can be ascribed to a fragmentation of the Co layers into SPM regions as explained by Lenczowski et al. [17] who observed that the decreasing GMR was also accompanied with a large drop of both the coercive field and the remanence magnetization of these multilayers, both factors speaking for the appearance of SPM regions due to the discontinuous nature of the extremely thin Co films. In line with this, Chowdhury et al. [141], after properly decomposing the total measured MR, obtained that, simultaneously with the reduction of the GMRFM term below 1 nm Co layer thickness,

Fig. 36a. Dependence of the giant magnetoresistance (MR) at room temperature on the magnetic Co layer thickness (dCo) for (l 0 0)-oriented ED [Co–Cu(dCo)/Cu(4 nm)]50 multilayers. The full and dashed-dotted lines are guides to the eye. Reprinted from Ref. [17] with permission of Elsevier.

Fig. 36b. The variation in the total GMR and its decomposed GMRFM and GMRSPM contributions are shown for ED [Co–Cu(dCo)/ Cu(4 nm)]50 multilayers as a function of the magnetic Co layer thickness. The lines are guides to the eye only. Reprinted from Ref. [141] with permission, copyright (2008) of the American Physical Society.

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the GMRSPM term exhibited a significant increase. Simultaneously, these authors have also observed [141] that both the coercive field and the MR peak position drastically decreased below 1 nm Co layer thickness. From the decomposition analysis, the SPM moment size could also be derived [141] the sharp drop of which indicated the SPM regions getting smaller and smaller with decreasing average Co layer thickness below about 1 nm. Dinia and coworkers [30,41] prepared ED Co–Cu/Cu multilayers in the P/P mode under the same conditions (bath composition and deposition potentials) as used by Lenczowski et al. [17]. A glass plate covered with a sputtered 50 nm thick Cu layer was used as substrate and XRD revealed an fcc-(1 1 1) texture with very faint first-order satellites for a [Co(6 nm)/Cu(4 nm)]25 multilayer for which the MR(H) curves were similar to those of Liu et al. [101] for comparable layer thicknesses (Fig. 30a). For a [Co(1.5 nm)/Cu(4 nm)]25 multilayer, Dinia and coworkers [30,41] obtained a total GMR of less than 1% only and the MR(H) curves exhibited a higher saturation field. This was an indication that in this multilayer already a fragmentation of the Co layers sets in, similarly to the observations of formers reports [17,141] about the decrease of GMR below 1 nm Co layer thickness, all this pointing toward a gradual transition from FM to SPM behavior of the magnetic layers. 5.2.3.2. Deposition with non-optimized Cu deposition potential from sulfate bath. In this section, we shall discuss the magnetoresistance results of reports on ED Co–Cu/Cu multilayers prepared under nonoptimized Cu deposition potential from a pure sulfate type bath by G/P or P/P method. Zhang and coworkers [63,101] electrodeposited Co–Cu/Cu multilayers on a Si(1 0 0)/Ti(10 nm)/ Cu(20 nm) substrate with the metal layers produced by sputtering. The electrolyte was deaerated for about 30 min before immersing the substrate and afterwards the cell was sealed and electrodeposition was carried out under stagnant conditions with ECu(SHE) = 0.32 V. The deposition started with a Cu layer at the chosen Cu potential for about 1–2 min in order to reach a steady state for Cu deposition and the multilayer deposition was started only afterwards. The authors observed, however, that there was an anodic current transient when switching the cathode potential from Co deposition to Cu deposition which evidenced Co dissolution during the Cu deposition cycle; this is actually expected since the Cu deposition potential applied was slightly more positive than the optimum potential for the sulfate electrolyte [ECu(SHE) = 0.34 V to 0.36 V]. XRD revealed a strong fcc-(1 0 0) texture, i.e., that of the Si substrate. The XRD pattern reported for a multilayer [Co(3 nm)/Cu(2 nm)]100 indicated the presence of first order satellites; however, the bilayer period was determined from the satellite peak position to be 2 nm whereas the nominal bilayer period was 5 nm. For the same multilayer which exhibited the largest GMR (4%), a nearly linear decrease of the resistance was observed up to 3 kOe (the maximum field applied) where saturation was not yet achieved. The authors have not specified if the LMR or TMR component was measured. Shima and coworkers [45,55,57,80,88] carried out an extensive study on the preparation, structure and GMR behavior of ED multilayers obtained by P/P method with ECu(SHE) = 0.31 V from a sulfate bath which was deaerated with argon immediately prior to deposition. A Si(0 0 1)/Cu(20 nm) substrate was used, and the Si wafer was etched with HF before evaporating a copper seed layer on top of it [80]. A Cu plate was attached to the back of the Si wafer via an eutectic InGa ohmic contact layer and this Cu plate served for applying the deposition current through the Si wafer itself. XRD revealed that the Cu seed layer exhibited a pronounced (1 0 0) texture which was completely inherited by the fcc Co–Cu/Cu multilayer as well. If the thickness of both layers was as high as typically 1 nm, clear first-order XRD satellites could be observed, indicating good structural coherence of the multilayers (see Fig. 27). The bilayer lengths derived from the XRD satellite peak position were in very good agreement with the nominal values [45,88]. Although an electrolyte with high ionic ratio (Co2+/ Cu2+ = 190) was used, the applied Co deposition potential ECo(SHE) = 0.77 V may not have been sufficiently negative to prevent significant Cu codeposition, i.e., one cannot exclude a high Cu content in the magnetic layer. For comparison, it is noted that for sulfate electrolytes the breakpoint of the cyclic voltammogram curve marking the onset of Co deposition in the cathodic scan is at E(SHE)  0.55 V [99,109] which is only by 0.22 V apart from the ECo value used by Shima et al. [80,88]. Furthermore, for a very similar electrolyte and a close value of the ionic ratio (176), Lenczowski et al. [17] measured 4 at.% Cu in dc-plated deposits obtained at even more negative deposition potentials ECo(SHE) = 0.96 V to 1.36 V. From an atomic absorption analysis, Shima and coworkers [88] tried

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to correct the nominal layer thicknesses obtained from the coulometric control parameters during deposition and ascribing all the observed deviation to a reduced current efficiency which reduction was especially strong for the Co layer. By considering that their applied Cu deposition potential (ECu(SHE) = 0.31 V) was somewhat more positive than the above specified optimum ECu(SHE) value (about 0.36 V to 0.38 V), a partial dissolution of the Co layer during Cu deposition can also be responsible for the apparent reduction of current efficiency. Thus, we need to conclude that the actual layer thicknesses can be even smaller for Co and larger for Cu than the ‘‘corrected” thicknesses given by these authors but we cannot estimate these changes. Magnetic measurements on [Co(dCo)/Cu(3.4 nm)]100 multilayers with dCo below about 1 nm as specified in the papers exhibited characteristics typical of SPM entities (absence of remanence and negligible hysteresis). The corresponding MR(H) curves also did not have a hysteresis (splitting) and the shape of MR(H) curves with the apparently high saturation fields indicated the dominance of the GMRSPM term in the total observed magnetoresistance [57,88]. For sufficiently thick Co layers such as typically dCo > 1.5 nm, the MR(H) curves were studied [80] for Co/Cu multilayers with varying Cu layer thicknesses. The MR(H) curves are shown in Fig. 37 for a series of ED [Co(1.8 nm)/Cu(dCu)]100 multilayers with dCu ranging from 1.1 nm to 4.4 nm [80]. From the fact that the observed longitudinal magnetoresistance decreases with increasing magnetic field, we can conclude that a GMR effect is the dominant mechanism for each Cu layer thicknesses. On the other hand, we can also infer from Fig. 37 that the GMR is fairly low for dCu = 1.1 and 2.2 nm and the nearly linear MR variation indicates very high saturation fields. Apparently, the magnetic layers are still fairly thin (probably thinner than specified, see above) and, with reference to Fig. 31b, for not sufficiently thick Cu layers the magnetic layers do not yet develop during growth a continuous coverage on the Cu surface and will not be completely ferromagnetic here; therefore, the GMR is dominated by a SPM contribution. Although for dCu = 3.3 and 4.4 nm the GMR is already significantly higher and there is also a splitting of the MR(H) curves, the FM contribution still delivers only about half of the total GMR at the maximum field applied whereas the high field slope indicates that actually the magnetoresistance will saturate in fairly high magnetic fields only. This implies that even for such large Cu layer thicknesses there is a large SPM contribution to the GMR. The origin of this behavior is not clear since the XRD patterns indicated a good quality of the multilayer structure and cross-sec-

Fig. 37. Longitudinal magnetoresistance curves of fcc-(1 0 0) oriented ED [Co(1.8 nm)/Cu(dCu)]100 multilayers with dCu = 1.1 nm (a), 2.2 nm (b), 3.3 nm (c) and 4.4 nm (d). For converting the monolayer (ML) thicknesses specified in the original paper, it was assumed here that 1 ML corresponds to about 0.2 nm. Reprinted from Ref. [80] by permission of The Electrochemical Society.

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tional TEM also revealed the layered structure with a columnar growth where the columns were as wide as 50–100 nm. A possible explanation for the strong SPM contribution to the GMR may be the high Cu content of the magnetic layers as was suggested above (see also Section 5.2.6). Nevertheless, Shima and coworkers [55,80] obtained a monotonous increase of the observed GMR as measured at their maximum field strength with Cu layer thickness up to about dCu = 5.5 nm. Weihnacht and coworkers [99,115,128,132,139] electrodeposited Co–Cu/Cu multilayers with typically 300 bilayer repeats on mechanically polished polycrystalline Ti sheets or on Si/Ta(5 nm)/Cu substrates with Cu seed layer thicknesses of 20, 50 or 100 nm (both metal layers were produced by evaporation). A tubular type cell was applied to avoid edge effects and the cathode facing upwards was placed to the bottom of the cell [99]. Mainly the G/P mode with ECu(SHE) = 0.01 V was applied but some multilayers were prepared also in the P/P and G/G mode or at a different Cu deposition potential. The magnetic layer contained typically 5 at.% Cu. A detailed structural study [128] was carried out on these multilayers. The TEM results were summarized in Section 4.1 (cf. Fig. 24) and the XRD results in Sections 4.2 and 4.3. Since the Cu deposition potential for the multilayers studied by Weihnacht and coworkers [99,115,128,132,139] was not optimized, in order to get an estimate for the concomitant layer thickness changes due to the significant Co-dissolution during the Cu deposition pulse, a chemical analysis of the overall multilayer composition was carried out [99,128] from which the actual layer thicknesses could be determined. At a Cu deposition potential of ECu(SHE) = 0.01 V, the Co layer thickness reduced from the nominal preset value of 3.4 nm to 2.0 nm and the Cu layer thickness increased from the nominal value of 1.0 nm to 2.4 nm. A cross-sectional elemental mapping analysis (Fig. 24) by high-resolution TEM on the same sample [128] clearly confirmed these layer thickness changes. For the GMR studies, both the LMR and TMR components were measured in magnetic fields up to 8 kOe whereby the multilayers were either mechanically stripped from the Ti sheets or were left on their Si/Ta/Cu substrates. First, we discuss some results reported in Ref. [99] in which work especially important aspects include (i) a study of the influence of Cu deposition potentials ranging from ECu(SHE) = 0.01 V to 0.26 V on the GMR and a (ii) comparison of the influence of different deposition pulse modes on GMR at identical nominal layer thicknesses. Fig. 38 shows the MR curves measured on three samples obtained in the P/P mode with Cu deposition potential values of ECu(SHE) = 0.01 V, 0.21 V and 0.26 V whereas keeping the Co deposition potential constant. It can be seen that in spite of the same nominal Co- and Cu-layer thicknesses in each case (3.4 nm and 1 nm, respectively), the magnitude of the GMR changes by a factor of three with Cu-deposition potential and the change is even non-monotonic. This strange behavior finds its explanation in the deposition potential dependent degree of the dissolution of the Co layer during the Cu deposition pulse and the excess Cu deposition in order to maintain the preset charge balance. This governs furthermore the actual layer thicknesses, layer thickness fluctuations and, in extreme cases, even the continuous or discontinuous nature of the layers. In the various deposition modes, these features can be differently controlled and it turned out [99] that the deposition parameters can be tuned for the G/P, P/P and G/G modes in a way as to provide practically the same GMR behavior in each case (Fig. 39). It should be kept in mind, however, that whereas the nominal (preset) layer thicknesses were 3.4 nm for the magnetic layer and 1 nm for the non-magnetic layer in each mode, the very similar GMR curves could be achieved only since for the three samples the actual layer thicknesses as determined from chemical analysis [128] were also very similar (between 1.9 and 2.5 for both layer types). However, this example represents a rather unique situation and it cannot be expected that any specific MR(H) parameters obtained by a selected deposition mode can be definitely achieved also by another mode. A particular insight into the multilayer formation under non-optimized Cu deposition potential can be obtained by considering Fig. 40 [99] which shows the measured MR(H) curves of two samples prepared in the G/P mode at very different Cu deposition potentials while keeping the same nominal (preset) layer thickness values (Co–Cu: 3.4 nm, Cu: 0.2 nm) together with the MR(H) curves of a dc-plated Co–Cu alloy prepared at the same current density as used for depositing the magnetic layer in both multilayers. The dc-plated sample containing about 5 at.% Cu displays a characteristic AMR behavior of a bulk ferromagnetic in that LMR > 0 and TMR < 0. It can be inferred from Fig. 40 that the multilayer deposited at ECu(SHE) = 0.26 V does not exhibit a GMR effect (LMR > 0) but rather an AMR effect is

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1

transverse MR, ΔR /R0 (% )

C u: -0.25 V C u: -0.45 V C u: -0.50 V

P /P mode 300 ML C o/C u on T i

0

2

C o: -1.9 V , -10.6 mC /cm

-1

2

C u: -2.8 mC /cm

-2 -3 -4 -5 -6 -7 -8 -10

-8

-6

-4

-2

0

2

4

6

8

10

H (kOe) Fig. 38. Influence of Cu deposition potential on the transverse magnetoresistance of electrodeposited Co–Cu/Cu multilayers of 300 bilayer repeats prepared in the P/P mode under conditions as specified. The nominal layer thicknesses were kept constant at 3.4 nm (Co–Cu) and 1 nm (Cu). Reprinted from Ref. [99] with permission of The Electrochemical Society.

1

transverse MR, ΔR /R0 (% )

0 -1

P /P mode C o: -1.9 V , -10.6 mC /cm^ 2 C u: -0.25 V , -2.8 mC /cm^ 2 G /P mode C o: -35 mA/cm^ 2, 0.3 s C u: -0.25 V , -2.8 mC /cm^ 2 G /G mode C o: -35 mA/cm^ 2, 0.3 s C u: -0.41 mA/cm2, 6.9 s

Comparison of P/P, G/P and G/G modes 300 ML Co/Cu on Ti

-2 -3 -4 -5 -6 -7 -8 -10

-8

-6

-4

-2

0 2 H (kOe)

4

6

8

10

Fig. 39. Comparison of the transverse magnetoresistance curves of electrodeposited Co–Cu/Cu multilayers prepared at various conditions in either of the P/P, G/G, and G/P modes. Reprinted from Ref. [99] with permission of The Electrochemical Society.

predominant in this multilayer similarly to the dc-plated, bulk Co–Cu alloy. This is because at this particular Cu deposition potential, the Co dissolution process is suppressed to a large extent and the actual Cu layer that forms after each magnetic layer remains very thin (0.9 nm) due to the low preset 0.2 nm Cu layer thickness. The presence of a large density of pinholes in such a thin Cu layer results in a FM coupling between adjacent magnetic layers, yielding an explanation for the observed AMR behavior. On the other hand, at ECu(SHE) = 0.01 V the Co dissolution is very strong and to keep the preset charge balance, the Cu layer grows to an actual thickness of 1.6 nm. Both the LMR and TMR components are negative in this second multilayer which contains sufficiently thick and certainly mostly continuous Cu layers, leading to a dominant GMR effect. This result confirms the assumption that the effective Cu layer thickness is larger due to a significant dissolution of Co and its replacement by Cu at the beginning of each Cu-deposition cycle in this multilayer.

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1

L

0

MR, ΔR /R0 (%)

-1

L T T

-2

ML Co/Cu (Cu: -0.50 V) -3

d.c.-plating:

G/P mode 300 ML Co/Cu on Ti

2

-4

Co(Cu): -35 mA/cm , 116 s

2

Co: -35 mA/cm , 0.3 s

-5

2

Cu: -0.6 mC/cm = 0.2 nm

-6 -7

L

-8

ML Co/Cu (Cu: -0.25 V) -9 -10

-8

-6

-4

-2

0

2

4

T 6

8

10

H (kOe) Fig. 40. Influence of Cu deposition potential on the longitudinal (triangles) and transverse (squares) magnetoresistance of electrodeposited Co–Cu/Cu multilayers (data points connected with solid lines) of 300 bilayer repeats prepared in the G/P mode under conditions as specified. The nominal layer thicknesses were kept constant at 3.4 nm (Co–Cu) and 0.2 nm (Cu). Note that whereas the multilayer obtained with ECu(SHE) = 0.01 V exhibits a clear GMR behavior (both the LMR and TMR components are negative for the whole range of magnetic fields), the MR curves of the multilayer obtained with ECu(SHE) = 0.26 V indicate an AMR behavior (LMR > 0, TMR < 0). The dashed lines connect the data points obtained for a d.c.-plated Co–Cu alloy deposited under identical conditions as the magnetic layer in the multilayer structure, and they indicate an AMR behavior. Reprinted from Ref. [99] with permission of The Electrochemical Society.

The work of Weihnacht et al. [99] also revealed that in the G/P mode, large Co deposition current density and large Cu layer thicknesses favors high GMR. The MR curves of the multilayer prepared along these lines exhibited the largest GMR (Fig. 41). A GMR of 10% is achieved at about 1 kOe, with a maximum sensitivity of 0.04%/Oe at low fields. A more quantitative characterization of the Co dissolution process could be obtained from the GMR data reported by Weihnacht et al. [99] when applying the method elaborated later [103] for decomposing the GMR into a FM and a SPM contribution. This is demonstrated in Fig. 42 where the MR(H) curves are shown for two multilayers V8 and V2 with the nominal thicknesses Co(1.7 nm)/Cu(1.0 nm) and Co(3.4 nm)/Cu(1.0 nm), respectively. According to the results of chemical analysis [128], the actual layer thicknesses are Co(0.7 nm)/Cu(2.0 nm) and Co(2.0 nm)/Cu(2.4 nm), respectively. The applied Cu deposition potential was ECu(SHE) = 0.01 V for both samples and under these conditions the Co dissolution is very strong during the Cu deposition cycle, explaining the large reduction of the Co layer thickness. Fig. 42 clearly reveals that for sample V8 with an actual Co layer thickness of 0.7 nm the GMRSPM component dominates the total measured magnetoresistance. This is because at such a small Co layer thickness (below 1 nm), a fragmentation of the Co layer may already occurs and SPM regions can form as discussed above in Section 5.2.3.1. On the other hand, the formation of SPM regions during the Co layer dissolution is much more modest for sample V2 (the actual Co layer thickness being 2.0 nm) as revealed in the bottom of Fig. 42. Concomitantly to this qualitative picture, the SPM fitting procedure [128] yielded an average SPM moment of 2140 lB for sample V8 with very thin, certainly discontinuous Co layers and 4400 lB for sample V2 with thick, fairly continuous Co layers. This difference in the size of the SPM regions in the two multilayers completely corresponds to what we can expect during the dissolution of the two type of Co layers with their different original thicknesses. A quantitative SPM analysis of the GMR data of ED Co–Cu/Cu multilayer films has provided conclusive evidence also for the role of substrate roughness [128]. Two multilayers [Co(2.0 nm)/ Cu(3.9 nm)]300 were deposited under identical conditions on either rough Ti sheets or on smooth Si/Ta/Cu substrates. The GMRFM and GMRSPM values of the multilayer on rough substrate were 7.1%

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0

-2

G /P mode 300 ML Co/Cu on Ti

longitudinal

2

C o: -106 mA/cm, 0.1 s

transverse

2

MR, Δ R /R0 (%)

C u: -0.25 V , -5.9 mC /cm -4

-6

-8

-10

-12 -10

-8

-6

-4

-2

0

2

4

6

8

10

H (kOe) Fig. 41. Longitudinal and transverse magnetoresistance curves of an electrodeposited Co–Cu/Cu multilayer prepared at optimized conditions in the G/P mode. Reprinted from Ref. [99] with permission of The Electrochemical Society.

and 2.2%, respectively, for the LMR component. The corresponding values for the nominally identical multilayer on smooth substrate were 9.4% and 0%. Very similar values were obtained also for the TMR component. The rough substrate unambiguously favors the formation of SPM regions (for the multilayer on Ti sheet, the average SPM moment size was about 5000 lB whereas for the multilayer on the smooth substrate, an estimate of the possible SPM moment size yielded values in excess of 30,000 lB which can already be considered as having a FM behavior). A mechanism was proposed by Ishiji and Hashizume [200] to explain their similar GMR results on their sputtered Co/Cu multilayers deposited on substrates with various roughnesses. An important conclusion to be drawn from the studies of Weihnacht and coworkers [99,115,128,132,139] on ED Co–Cu/Cu multilayers is that one can achieve fairly good structural quality in terms of XRD satellite reflections [128] and GMR characteristics (cf. Fig. 41) even at non-optimized Cu deposition potentials under certain conditions. These studies have, however, also revealed that particular caution should be exerted when evaluating GMR results on such multilayers as shown by the example in Fig. 43. A series of ED Co–Cu/Cu multilayers [Co(2 nm/Cu(dCu)]300 was produced in the G/P mode at ECu(SHE) = 0.01 V, i.e. with strong Co-dissolution [99]. The total magnetoresistance measured in a magnetic field of 8 kOe shows a minimum in GMR as a function of the Cu layer thickness (Fig. 43) which would be hard to explain having these data only. On the other hand, carrying out a quantitative SPM analysis [139], it turned out that the observed behavior of the total magnetoresistance is merely the result of an interplay between the different dependence of the GMRFM and GMRSPM components on the Cu layer thickness. It can be seen in Fig. 43 that by separating out the GMRSPM contribution from the measured data, the GMRFM term shows a similar behavior as presented in Section 5.2.3.1 for ED Co–Cu/Cu multilayers deposited under optimized Cu deposition potential. Interestingly, the SPM region size was found to be fairly constant throughout this series at around 5000 lB [139]. By making some assumptions, the typical SPM region size extending throughout the whole magnetic layer thickness (SPM islands magnetically decoupled from the magnetic layers) could be estimated to be (2 nm 3.5 nm 3.5 nm) [139]. The fact that it remained constant with dCu is in agreement with what we can expect on the basis of the formation of SPM regions. Namely, as discussed above, they form due to a partial dissolution of the magnetic layers that takes place at the beginning of the Cu pulse, and the size of the SPM regions does not change after the complete coverage of the Co layer with Cu is achieved. The uniform SPM cluster size is due to the constant nominal (and actual) magnetic layer thickness at the same Cu deposition potential

MR, ΔR/R0 (%)

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V V8, LMR ( tCo = 1.7 nm )

5%

V2, LMR ( t Co = 3.4 nm )

V8, LMR, FM

V8, LMR, SPM V2, LMR, SPM

V2, LMR, FM -9

-6

-3

0

3

6

9

H (kOe) Fig. 42. Top: room-temperature LMR curves for multilayer samples V2 and V8. Middle: Decomposition of the SPM and FM contribution of the LMR curve for sample V8. Bottom: decomposition of the SPM and FM contribution of the LMR curve for sample V2. Layer thicknesses specified here are nominal; for actual layer thicknesses, see text. Reprinted from Ref. [128] with permission of American Scientific Publishers.

Fig. 43. Evolution of the total (saturation) magnetoresistance and the GMRFM and GMRSPM contributions with Cu layer thickness for a multilayer series. The magnetic layer thickness was 2.0 nm. Lines are only a guide for the eye. Reprinted from Ref. [139] with permission of Elsevier.

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applied throughout the multilayer series. Therefore, the constancy of the SPM region size is consistent with the model of the cluster formation and actual appearance of such regions as discussed beforehand. However, the probability of the SPM region formation can, indeed, be a function of the Cu layer thickness [121]. This can probably be explained by the leveling effect provided by the fairly thick Cu layers as discussed in some more detail in Ref. [115]. Pandya and coworkers [130] electrodeposited Co–Cu/Cu multilayers with ECu(SHE) = 0.26 V on glass plates covered by indium-tin oxide (ITO) with small bilayer numbers (2–20) and large individual layer thicknesses (dCu  8 nm and dCu  6 nm). The value of the Cu deposition potential hints for a significant Co dissolution during the Cu deposition cycle. An fcc-(1 1 1) texture of the multilayers was found by XRD but no satellite reflections could be seen. The root-mean-square surface roughness estimated by a surface profiler was found to increase nearly linearly from about 13 nm of the ITO-coated glass substrate to about 40 nm for the thickest multilayers with bilayer repeats of 20, corresponding to a total multilayer thickness of about 280 nm. The TMR component of the room-temperature magnetoresistance was only measured and the resistance decreased nearly linearly with increasing H. The MR value reached 1.08% (15 bilayers) and 1.5% (20 bilayers) at the maximum field strength (9 kOe) applied. The non-saturating MR behavior indicates the presence of an SPM contribution. Pandya and coworkers [136] have also investigated similarly prepared ED Co–Cu/Cu multilayer films with 50 bilayers and 15% GMR was achieved at 6 kOe with a nearly linear MR(H) variation beyond H = 2 kOe. The significant improvement in GMR was apparently due to the larger bilayer number only and this suggests the formation of a much more compact multilayer deposit at more than doubled total thickness on the fairly rough ITO substrate the surface roughness of which was established by AFM. A further reduction of the MR(H) saturation field, down to a complete saturation at about 1 kOe, was obtained by depositing an identical multilayer stack on a Cu-coated Si wafer substrate of much lower surface roughness although the saturation GMR was reduced to about 1/3 of the value on the glass/ITO substrate. The very small GMR saturation field indicates that the GMR contains an FM contribution only for the smooth Si/Cu substrate. This is in agreement with earlier findings [128,200] in that surface roughness plays a decisive role in the formation of SPM regions. On the other hand, it was also found in Ref. [136] that the Si/Cu/[Co-Cu/Cu]50 structure had a resistance half of that of the identical multilayer on glass/ITO substrate, corresponding to expectation by considering the surface roughness differences of the two substrates. These results on the resistance and the GMR might also suggest that, under identical deposition conditions, a better defined multilayer structure forms on the smooth substrate. This latter multilayer has, however, a small GMRFM contribution only whereas the rougher substrate leads to the formation of a more disordered multilayer structure with a high density of SPM regions which, on the other hand, give rise to a larger total GMR. Jyoko et al. [35] electrodeposited Co–Cu/Cu multilayers at ECu(SHE) = 0.20 V on fcc-(1 1 1) Cu in an atmosphere of purified nitrogen. An XRD study of [Co(2 nm)/Cu(dCu)]100 multilayers with dCu = 1.9 nm or 3.8 nm revealed an fcc-(1 1 1) texture but also the clear presence of hcp-Co peaks the intensity of which decreased with increasing Cu layer thickness; on the other hand, satellite peaks could not be observed at neither Cu layer thickness. The authors noticed a tendency for Co-dissolution at the beginning of the Cu deposition pulse; this can actually be expected from the fairly positive Cu deposition potential applied in comparison with the optimum value. As to the magnetoresistance, the TMR component was only measured and the shape of the reported MR(H) curves for dCu = 1.9 nm indicated the dominance of a GMRSPM contribution with a total GMR value of about 2.2% at 10 kOe where no saturation was yet achieved. Since the Co layer deposited was 2 nm only, due to the strong Co dissolution this thickness is significantly reduced and this may give rise to the formation of SPM (see Ref. [99]). For dCu = 3.8 nm, the GMR is slightly higher and it shows a tendency of reaching a high value already at around 2 kOe and approaching saturation at 10 kOe. This indicates a smaller SPM fraction and a more continuous growth of the magnetic layers. Lashmore and Hua [22] electrodeposited Co–Cu/Cu multilayers at ECu(SHE) = 0.02 V on annealed polycrystalline (pc) Cu sheets or on ITO-coated glass with either stirring or not stirring the electrolyte. A peculiarity of this work was the application of a 2 s long open-circuit period of the cell between each Co and Cu pulses. Since the open-circuit period followed the Co pulse, it seems very likely that a significant amount of Co was dissolved while Cu was deposited under the open-circuit condition. (The introduction of an open-circuit period here is likely to be a wrong adaptation of the pulse sequence

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177

developed for Ni–Cu compositionally modulated alloys [234,256] where it might have served as a tool to reduce the hydrogen evolution on the freshly deposited Ni surface, hence improving the sharpness of the Ni/Cu interface.) The rather positive Cu deposition potential with respect to the optimum value is expected to lead to a very strong dissolution of Co during the Cu deposition pulse. The nominal layer thicknesses were dCo = 3 nm and dCu = 1 nm or 2 nm with bilayer repeats around 300 (the total multilayer thicknesses ranged from 1.2 lm to 1.6 lm). XRD indicated a strong fcc-(1 0 0) texture with clear first-order satellites, especially for the lower Cu-content in the magnetic layer. These authors reported MR(H) curves, without specifying whether the LMR or the TMR component was measured, for various combinations of the layer thicknesses, Co-contents in the magnetic layers, substrates used and the presence or absence of stirring. In general, for thick Cu layers (nominal thickness: 2.0 nm) a more pronounced hysteresis occurs with a larger fraction of the GMRFM component saturating at about 0.5 kOe, the maximum magnetic field applied. For thin Cu layers (nominal thickness: 1.0 nm), apart from a small GMRFM component, usually a hysteresis-free, non-saturating MR behavior was observed, indicating the presence of a dominating SPM contribution similarly to the case in Ref. [80]. The measured total GMR at the maximum fields applied was typically around 5% or, in some cases, around 2%. A comparison of [Co(2 nm)/Cu(2 nm)]50 multilayers electrodeposited on ITO and pc-Cu yielded a hysteresis-free non-saturating GMR with a maximum value of 3% on the former substrate and an MR(H) curve with hysteresis and saturating with 6% on the latter one, in both cases in a magnetic field of 0.4 kOe. These authors [22] have also investigated the influence of annealing on GMR. However, since not the MR(H) curves just the DR/R values were reported as a function of the annealing temperature, it is hard to assess the actual changes appearing upon annealing but the GMR was found to decrease generally with annealing. Yan et al. [92] reported a GMR of 4% at room temperature and at 10 kOe for ED Co–Cu/Cu multilayers obtained in the P/P mode from a pure sulfate electrolyte but, unfortunately, the type of reference electrode was not specified. The peculiarity of this work was the use of a conducting polymer cathode in view of possible application of the multilayer as magnetoresistive sensor on a flexible substrate. Cziráki and Tichy [116] attempted to correlate structural properties and GMR magnitude by using magnetoresistance data reported for ED Co–Cu/Cu multilayers [99,109] which were discussed in this and in the previous section. A well-defined correlation was found [116] in that the GMR increased first with the average fcc multilayer lattice parameter obtained from the main Bragg peak position on the same samples and, after a sharp maximum, it decreased again. It should be pointed out that this apparent correlation does not mean an intrinsic dependence of GMR on the lattice parameter since the latter is just an overall quantity determined primarily by the ratio of the layer thicknesses and the lattice parameters of the constituent metals. The true quantities controlling the GMR are the Co and Cu layer thicknesses which are internal parameters in the GMR vs. lattice parameter plot. The correct physical picture we have to keep in mind is that with increasing Cu layer thickness first the multilayer lattice parameter increases and, independently, the GMR also increases (see Figs. 33a and 35). A decrease of the Co layer thickness results in changes in the same sense, again independently, both for the multilayer lattice parameter and the GMR. Beyond a certain high (low) value of the Cu (Co) layer thickness, the GMR starts to decrease (see Figs. 35, 36a and b, respectively), whereas the multilayer lattice parameter continues to increase, the latter two facts explaining the higher-end part of the apparent GMR vs. lattice parameter correlation. On passing this section, we should point out again that whereas fairly good multilayer quality and GMR characteristics can be achieved with either optimized or non-optimized Cu deposition potential in the G/P and P/P single-bath deposition modes, the actual layer thicknesses agree with the nominal values only for the case of optimized Cu deposition potential (by assuming the most commonly used charge-control method for setting the layer thickness during deposition). If a non-optimized Cu deposition potential is applied, the actual Co layer thickness will always be smaller than the preset nominal value and the reverse is true for the Cu layer thickness. In such a case, the actual layer thicknesses should be determined from independent measurements (e.g., overall composition analysis of the multilayer, cross-sectional TEM). The same holds true, due to the unavoidable exchange reaction, also for the case of multilayers prepared by G/G deposition mode to be discussed in Section 5.2.4. On the other hand, the actual bilayer repeat period will remain the same as the nominal repeat period for all three deposition modes.

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The notice in the previous paragraph had to be made at this point for the reason that an optimization of the Cu deposition potential has been carried out for the pure sulfate bath only. For the other types of bath (sulfate baths with some additives or the sulfamate/sulfate type baths) to be discussed in the following two sections, no such optimization was carried out for ED Co–Cu/Cu multilayers prepared for GMR studies and the reported Cu deposition potential values were chosen on the basis of cyclic voltammograms which have, however, a limited validity of use only as was emphasized in Ref. [231] (see also Section 3.6). The limited validity of use derives from the fact that the onset potential of Co dissolution as observed on the potentiodynamic curves seriously depends, besides the bath composition, on several experimental parameters such as sweep rate, cathodic limit, ion concentrations, pH and also on the number of cycles recorded [231]. 5.2.3.3. Deposition from sulfate bath with additives. Various additives to the sulfate bath were used to obtain ED Co–Cu/Cu multilayer films with GMR behavior and the detailed deposition parameters are listed in Table 3. The additives included alkali chlorides, trisodium citrate and chromic acid (CrO3). The reported GMR results will be discussed according to this grouping in the present section whereas the influence of other common galvanic bath additives will be dealt with in Section 5.2.4.4. Whereas no data on the optimized Cu deposition potential are available as yet for the sulfate bath with chromate addition, such data can be found in Ref. [231] for some sulfate baths with chloride or Na-citrate addition. The optimum ECu values varied in a range of about 0.05 V when comparing the pure sulfate bath, the sulfate/chloride and the sulfate/citrate baths; however, evidently the optimum should be established for each actual bath composition individually. It can be inferred from Table 3 that most of the applied ECu(SHE) values differ by more than 0.05 V from the expected optimum values. For those which fell towards more negative potentials, incorporation of Co into the Cu layer can be expected, whereas for deviations towards less negative ECu values, the Co dissolution process may have occurred. Chassaing and coworkers [24,25,49,74] electrodeposited Co–Cu/Cu multilayer films in the P/P mode from a sulfate electrolyte with CoCl2 additive at ECu(SHE) values ranging from 0.01 V to 0.36 V. The substrates used were either a glass plate covered with a sputtered Au layer of typically 15 nm thickness or ITO. In both cases, an fcc-(1 1 1) texture of the Co–Cu/Cu multilayer films was revealed by XRD, on the Au substrate even first-order satellite reflections were visible. Magnetoresistance data were reported for 4.2 K and 77 K only. Most of the MR(H) curves presented did not show a saturation up to magnetic fields of 5–10 kOe, indicating the presence of a fairly large contribution of SPM regions to the GMR. In some case this may have been due to the very small layer thicknesses, e.g., for a multilayer [Co(0.3 nm)/Cu(0.3 nm)]50 which corresponds to the case depicted in the lower left panel of Fig. 34. Whereas fairly large GMR values (10–15% at 6 kOe) were obtained at 4.2 K for layer thicknesses above 1 nm in a wide range of the Cu deposition potential (ECu(SHE) varying from Table 3 Deposition parameters of ED Co–Cu/Cu multilayer films with GMR which were obtained under P/P control from a single sulfate bath with additives and containing the bath components CoSO4 + CuSO4 + X + A where X and A denote buffering agents (H3BO3) and the various additives (e.g., NaCl), respectively. The cathode potentials applied for deposition are referenced to the standard hydrogen electrode (SHE), see Table 1. Authors and Refs. in []

X+A

pH

Co2+/Cu2+ ionic ratio

ECo(SHE)/ECu(SHE)

Sulfate/chloride bath Chassaing and cow. [24,25] Chassaing et al. [49,74]

CoCl2 + NaCl + H3BO3 CoCl2 + NaCl + H3BO3

3.0 1.7

141 141

0.81 V/0.26 V 0.76 V to 0.91 V/0.01 V to 0.36 V

Sulfate/citrate bath Gomez et al. [84,95] Pandya and cow. [130,134,146] Uhlemann et al. [98] Yao and cow. [93,100,114]

Na-citrate Na-citrate (+NaCl) Na-citrate + NaCl Na-citrate (+NaCl)

4.7 4.0–6.0 5.4 5.0–6.0

88–175 40 7.5 22

0.90 V/0.25 V 1.26 V/0.56 V 0.80 V/0.30 V 0.81 V to 0.86 V/0.26 V

Sulfate/chromate bath Jyoko and cow. [33–36,42,52]

CrO3 + H3BO3

100

0.90 V/0.20 V

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0.06 V to 0.26 V) [49] and even 25% GMR was achieved at 4.2 K in one case [74], in lack of reported magnetoresistance data above 77 K, we cannot conclude if there can be any measurable GMR in these multilayer films at room temperature. Although Chassaing and coworkers specify in Ref. [24] that the magnetic layers in the multilayer contain less than 4 at.% Cu, the cyclic voltammogram curves they reported [24,25,74] indicate that the applied Co deposition potential was by only 0.2–0.35 V more negative than the onset of Co deposition. Therefore, the strong SPM contribution to the GMR in their multilayers may have been also due to the much higher actual Cu content in the magnetic layers (see also Section 5.2.6.). Chassaing [74] also reported that whereas 25% GMR was obtained at 6 kOe and 4.2 K from the sulfate/chloride bath, the addition of saccharin and the agent SDS caused a decrease of GMR to 10% and the MR(H) curve shape became more dominated by a non-saturating magnetoresistance component. Thus, these additives evidently contributed to the formation of even more SPM regions in the multilayers during deposition. Na-citrate addition to the sulfate bath was used by Gomez et al. [84,95], Pandya and coworkers [130,134,146], Uhlemann et al. [98] and Yao and coworkers [93,100,114] to electrodeposit Co–Cu/ Cu multilayers with GMR behavior. Gomez et al. [84,95] prepared ED Co–Cu/Cu multilayers with an fcc-(1 1 1) texture on glass/ ITO(25 nm) substrates at a Cu deposition potential of ECu(SHE) = 0.25 V. The MR(H) curves reported for 300 K and 27 K showed no saturation in magnetic fields up to 10 kOe where the GMR value was 2% and 6%, respectively. The MR(H) curves reported indicate a predominantly SPM contribution to the GMR for any layer thickness combination studied. Pandya and coworkers [130,134,146] prepared ED Co–Cu/Cu multilayer films on ITO-coated glass substrate at ECu(SHE) = 0.56 V which is much more negative than the one they used for their multilayers from the citrate-free bath [130,136] and one may suspect the codeposition of Co into the nonmagnetic layer here. The XRD patterns revealed the same fcc-(1 1 1) texture of the multilayers and the absence of satellite reflections as found for the citrate-free bath. The presence of Na-citrate reduced the saturation field of the room-temperature MR(H) curves with respect to the citrate-free sulfate bath, but from the field evolution of the magnetoresistance one should conclude for the presence of substantial SPM contribution to the GMR also for most multilayers deposited from the sulfate/citrate bath. Multilayers with unusually thick layers were investigated such as, e.g., [Co(22 nm)/Cu(6 nm)]50 and the GMR value measured at a fixed magnetic field was found to increase monotonously as a function of the Co layer thickness from 8 nm to 22 nm. This finding is very hard to understand in the light of opposite former results discussed above (cf. Figs. 32a, 36a and b) for the Co-layer thickness dependence of the GMR. Another feature in the works of Pandya and coworkers [134,146] was that the LMR(H) and TMR(H) curves were often practically coinciding. This is a surprising fact since at such large Co layer thicknesses a significant AMR (=LMR  TMR) would be expected as a result of bulk scattering within the thick Co layers. Although GMR values as high as 10–12% were obtained by these authors in most of their multilayers, the absence of an AMR component of the measured MR(H) curves questions the presence of a well-defined layered structure in these multilayers for which no evidence from structural studies has been presented anyway. Uhlemann et al. [98] prepared ED Co–Cu/Cu multilayer films [Co(2.5 nm)/Cu(1.9 nm)]30 on Si(1 0 0)/Permalloy(20 nm)/Cu(20 nm) substrates. The applied Cu deposition potential (ECu(SHE) = 0.30 V) might have been fairly close to the optimum. On the other hand, the magnetic layer certainly had a large Cu content since the ionic ratio Co2+/Cu2+ = 7.5 was very low and the applied Co deposition potential ECo(SHE) = 0.80 V was only by about 0.2 V more negative than the onset of the Co deposition, according to the reported cyclic voltammogram curve. No MR(H) curves were reported and the room-temperature GMR measured at 1.5 kOe was about 3% only. The main interest in this work was to study the change of GMR as a function of annealing temperature (annealing was carried out for 1 h in vacuum). The GMR increased continuously from 2.7% (as-deposited state) to 3.3% (annealing at around 300 °C) and decreased rapidly for higher annealing temperatures. An XRD study of the fcc-(1 1 1) textured multilayers indicated a relaxation around 280 °C (increase of lattice parameter and grain size). The slight improvement of the GMR can be partly due to the relaxation leading to a more perfect multilayer structure and due to a decomposition occurring at the interfaces as a consequence of the immiscibility of Co and Cu. The degradation of the GMR at high annealing

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temperatures can be ascribed to a gradual loss of superlattice structure due to the increased atomic mobility. Yao and coworkers [93,100,114] electrodeposited Co–Cu/Cu multilayer films in P/P mode on Si substrate. XRD revealed first-order superlattice satellites around the fcc-(1 1 1) peak for bilayer repeats at and below 20 nm. The repeat periods determined from the positions of these satellites were by 5–12% only higher than the nominal values, in agreement with other similar observations [128,149,255]. The MR(H) curve presented [93] for a Co(1.4 nm)/Cu(3.5 nm) multilayer exhibited saturation at about 2.5 kOe with a saturation GMR value of 7.5%. For a series of multilayer films with the same layer thicknesses as specified before, the GMR was found to increase from 3% at 30 bilayers up to a saturation value of 10–11% when the bilayer number was between 250 and 300 [100]. For another series with fixed Co layer thicknesses, the GMR exhibited a similar overall evolution with Cu layer thickness as shown in Fig. 35a and reached a maximum GMR of 7%. Although there were large fluctuations in the GMR magnitude, in contrast to the interpretation of Yao et al. [114], these cannot be considered as manifestation of an oscillatory GMR mainly because the positions of the observed maxima and min-

Fig. 44. Room-temperature transverse magnetoresistance DR/R vs. magnetic field curves for two series of electrodeposited [Co(2.0 nm)/Cu(tCu)]300 multilayers: (a) results for multilayers deposited from a pure sulfate electrolyte; (b) results for multilayers deposited from a sulfate/chromate bath. The insets show the dependence of saturation magnetoresistance on Cu spacer layer thickness. Reprinted from Ref. [35] with permission of The Electrochemical Society.

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ima do not coincide with the well-known positions in this multilayer system (see Fig. 8) and the actual Cu layer thicknesses are anyway certainly larger than the specified nominal values due to the exchange reaction taking place at the applied Cu deposition potential ECu(SHE) = 0.26 V being more positive than the optimum value. Jyoko and coworkers [33–36,42,52] electrodeposited Co–Cu/Cu multilayer films from a sulfate bath with chromate (CrO3) addition. With reference to Section 5.2.3.2 where we discussed the results of Jyoko and coworkers on Co–Cu/Cu multilayers deposited from a pure sulfate bath [35], it is noted here that, at the same applied Cu deposition potential ECu(SHE) = 0.20 V where the Co dissolution process was very strong, the addition of chromate helped effectively suppress the formation of a hcp-Co phase already for dCu = 1 nm and the development of a fcc-(1 1 1) multilayer structure. On the other hand, XRD studies did not reveal the presence of satellite reflections either with chromate addition. The MR(H) curves are compared in Fig. 44 for some of their multilayers obtained from baths without and with chromate. Although the GMR at 10 kOe was as high as 11% for a [Co(2.0 nm)/Cu(0.9 nm)]300 multilayer which value is about five times larger than the GMR of a similar multilayer [Co(2.0 nm)/ Cu(1.9 nm)]300 obtained from a chromate-free bath, the shape of the MR(H) curve indicates a strong Co dissolution also from the chromate bath as revealed by the non-saturating MR(H) curve with the dominance of the GMRSPM term. The chromate bath became, however, effective at larger Cu layer thicknesses since for a multilayer [Co(2.0 nm)/Cu(3.25 nm)]300 the GMR reached sharply a value of about 15% at around 2 kOe already with a slight further change only up to 10 kOe. A possible explanation for the beneficial effect of the chromate addition to the bath could be that presence of chromate prevents the Cu incorporation into the magnetic layer although the authors have not presented cyclic voltammogram curves which could support or deny the validity of this hypothesis. 5.2.3.4. Deposition from sulfamate/sulfate bath. The preparation details for ED Co–Cu/Cu multilayers with GMR which were deposited in the P/P mode from baths where Co-sulfamate Co(NH2SO3)2 was used instead of the Co-sulfate component or sulfamic acid (NH2SO3H) having an equivalent effect as the Co-sulfamate was added to the sulfate bath are summarized in Table 4. We do not have any estimate on what is the optimum Cu deposition potential in this type of baths. Bird and Schlesinger [15] electrodeposited Co–Cu/Cu multilayer films in P/P mode from a sulfamate/sulfate bath with an unusually large ionic ratio Co2+/Cu2+  8000 which certainly causes the Cu deposition pulse to be extremely long. The Co layer thickness was kept constant at 3.2 nm and the Cu layer thickness was varied from 0.5 to 8 nm with bilayer repeat numbers between 800 and 6000. The total multilayer thickness was at least several micrometers and in many cases it may have exceeded even 20 lm. The LMR component of the magnetoresistance was measured at room temperature and saturation was claimed to have reached for H < 4.5 kOe but no MR(H) curves were presented in this brief report. The saturation GMR values were displayed for ten different Cu layer thicknesses up to dCu  4.3 nm and were interpreted by the authors as following the expected oscillatory (RKKY-type) behavior as a function of the Cu layer thickness. A maximum room-temperature GMR of 55%, equal to that reported for sputtered Co/Cu multilayers [177,178,189], was obtained for dCu around 0.75 nm, with the second and third GMR maximum being around 2 nm and 3.5 nm, respectively. Unfortunately, in lack of sufficient details about the preparation conditions and the magnetoresistance measure-

Table 4 Deposition parameters of ED Co–Cu/Cu multilayer films with GMR which were obtained under P/P control from a single sulfamate/ sulfate type bath with buffering agents X (H3BO3), supporting electrolytes (Na2SO4) and additive A. The cathode potentials applied for deposition are referenced to the standard hydrogen electrode (SHE), see Table 1. Authors and Refs. in []

X+A

pH

Co2+/Cu2+ ionic ratio

ECo(SHE)/ECu(SHE)

Co-sulfamate + CuSO4 Bird and Schlesinger [15] Sasaki et al. [113] Fedosyuk et al. [66]

Antipitting agent + H3BO3 H3BO3 H3BO3

3.8 4.0

8000 62 45

0.84 V/+0.05 V 1.06 V/0.16 V 1.76 V/0.26 V

CoSO4 + CuSO4 Safak et al. [145]

NH2SO3H + Na2SO4 + H3BO3

2.0–3.0

37

1.06 V/0.16 V

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ments, especially the shape of the MR(H) curves, one cannot properly assess the validity of these results which could not be even approached in most subsequent studies. Sasaki et al. [113] electrodeposited Co–Cu/Cu multilayers in P/P mode from a sulfamate/sulfate electrolyte on Si(0 0 1)/Cu(0 0 1) substrates where the 40 nm thick Cu seed layers were obtained by evaporation. For a [Co79Cu21(0.9 nm)/Cu(4.1 nm)]100 multilayer, XRD diffraction indicated the same fcc-(1 0 0) texture as that of the substrate and even satellite reflections were revealed from which a bilayer length of 5.2 nm was deduced, in good agreement with the nominal value (5 nm). In spite of the good multilayer structure [a scanning electron microscopy (SEM) study also indicated a fairly smooth and uniform surface], the MR(H) curve was very sharp and did not seem to exhibit a hysteresis at room temperature, indicating the dominance of an SPM contribution to the GMR reaching nearly 20% at H = 14 kOe. Fedosyuk et al. [66] electrodeposited Co–Cu/Cu multilayers in P/P mode from a sulfamate/sulfate electrolyte on GaAs wafers by contacting the back of the wafer for deposition. The room-temperature MR(H) curves for a [Co(2 nm)/Cu(1.5 nm)]5 multilayer gave 2% and 3% GMR for the LMR and TMR components, respectively. Although saturation was achieved for about H = 5 kOe, the sharp MR peaks and the field evolution of the magnetoresistance indicate the dominance of an SPM contribution to the GMR. Safak et al. [145] electrodeposited Co–Cu/Cu multilayers in P/P mode from a sulfate electrolyte with sulfamic acid added on polished Ti sheets at various pH values. The total thickness of the multilayers was fixed at around 5 lm and the multilayers were mechanically stripped from their substrates. XRD revealed a strong fcc-(1 1 1) texture but no satellite reflections were observed. For a pH of 3.0, the room-temperature MR(H) curves of Co(3 nm)/Cu(1 nm or 2 nm) multilayers indicated a dominant SPM contribution: the magnetoresistance did not show a saturation up to 8 kOe; the total GMR at the maximum field was about 9% and 3.5% for the thinner and thicker Cu layer, respectively. Chemical analysis showed that at higher pH (3.0) the multilayer contains more Cu than at low pH (2.0) and this is probably due to differences in the Cu content of the magnetic layers. A comparison of two multilayers grown at pH 2.0 and 3.0 gave much larger GMR (ca. 15%) with smaller saturation field for low pH and about 4% GMR with higher saturation field for high pH. This difference can be explained on the basis of differences in the Cu-content for the two multilayers as will be discussed separately in Section 5.2.6. 5.2.4. GMR of ED Co–Cu/Cu multilayer films prepared in G/G mode from a single bath 5.2.4.1. Deposition from sulfate bath. The present authors [77] have investigated the MR characteristics of ED Co–Cu/Cu multilayers prepared in G/G mode from a single sulfate electrolyte with various deposition parameters on polished Ti sheets. The bilayer number was varied from 700 to 1500, with a total multilayer thickness ranging from 8 to 18 lm. The magnetic layer composition was about Co95Cu5 as obtained from chemical analysis data. A well-defined fcc-(1 1 1) textured multilayer structure was obtained as revealed by the presence of first-order XRD satellite reflections and cross-sectional TEM images. A room-temperature GMR magnitude of 8–9% was measured at H = 8 kOe on these multilayers after mechanically peeling them from the Ti substrates. The MR(H) curves did not saturate and this feature indicates the presence of a dominating GMRSPM contribution. The total GMR was found to decrease slightly with increasing bilayer number, hinting at some degradation of the multilayer quality for very large total deposit thicknesses. A lateral variation of the GMR over the cathode area could be demonstrated by using 1–2 mm wide strips cut from the 20 by 20 mm large deposits. This variation arises from the buoyancy-induced spontaneous upward convection along the cathode surface which was in vertical position during deposition. A particularly interesting result from this work is shown in Fig. 45 where the MR(H) curves are displayed for three multilayers for which the Cu deposition current density was only different (0.25 mA/cm2, 0 and +0.25 mA/cm2). A negative current density corresponds to the normal cathodic current (deposition) and this yielded the largest GMR. Surprisingly, a zero or even positive (anodic) Cu ‘‘deposition” current yielded a GMR which was reduced by 1% only with respect to the cathodic current case. The occurrence of a fairly large GMR effect as observed here for zero or anodic current is due to the formation of a sufficiently thick Cu layer separating properly the magnetic layers to the extent that a direct FM coupling between them is prevented. This is a clear indication for the importance of the

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-2

ΔR/ R0 (%) / longitudinal /

0

i (Cu) = -0.25 mAcm i (Cu) = 0 -2 i (Cu) = 0.25 mAcm

-1

-2

-3

-4

-5 -8

-6

-4

-2

0

2

4

6

8

H (kOe) Fig. 45. Longitudinal room-temperature magnetoresistance curves of three selected ED Co–Cu/Cu multilayer films produced with different Cu deposition current densities i(Cu) as indicated in the legend; further deposition conditions: i(Co) = 32.5 mA/ cm2, t(Co) = 0.65 s, t(Cu) = 5.0 s. Reprinted from Ref. [77] with permission of The Electrochemical Society.

exchange reaction unavoidably taking place during the Cu deposition pulse in the G/G mode until the previous Co layer is completely covered by Cu. The exchange reaction consists of a spontaneous dissolution of the previously deposited Co atoms which are replaced by Cu atoms from the electrolyte and this reaction operates with zero net charge balance. This process is apparently so effective that even for anodic current at the cathode, i.e., when the Faradaic process provides charges for Co dissolution only, a sufficiently thick Cu layer can build up during the Cu deposition pulse. Of course, the strong Co dissolution of the previously deposited Co-rich layer promotes the formation of SPM regions similarly to the G/P and P/P deposition modes with Cu deposition potentials more positive than the optimum value (see Section 5.2.3.2). The first demonstration [103] of the decomposition of the FM and SPM contributions to both the magnetization and magnetoresistance was performed by using results on ED Co–Cu/Cu multilayer films prepared in the G/G mode with anodic current at the cathode. As to the effectiveness of the exchange reaction, for an ED Co–Cu/Cu multilayer film prepared in the G/G mode, for which the MR(H) curve was presented in Fig. 39 above, a chemical analysis revealed [128] that the actual layer thicknesses are dCo = 2.5 nm and dCu = 1.9 nm. Since the nominal thicknesses were dCo = 3.4 nm and dCu = 1.0 nm, this means that in this case the excess Cu layer thickness due to the exchange reaction is as high as 0.9 nm, i.e., the actual Cu layer thickness is almost twice the nominal value. The deleterious effect of NaCl additive on the structure and properties of ED Co–Cu/Cu multilayer films obtained from a sulfate bath in the G/G mode has also been studied by the present authors [78]. The addition of NaCl decreased the current efficiency of the multilayer deposition and resulted in multilayers with lower GMR and higher electrical resistivity than in the case of the chloride-free bath. The grain size of the deposit and the degree of orientation also decreased if the additive was present. It was concluded that the adsorption of chloride, the formation of copper(I) intermediate, the change in the deposition mechanism, the increase in nucleation rate, and the occurrence of a non-Faradaic current transient all contribute to the enhancement of the structural disorder that leads to the loss in GMR. Some general conclusions have also been formulated about the role of additives in electrodeposition baths, especially concerning the differences in d.c. plating and pulse-plating. There was an attempt [94] to control the GMR of ED Co–Cu/Cu multilayer films in the G/G mode by introducing a capacitance applied parallel to the electrolytic cell. The extra capacitance varying from 0.05 mF to 6.7 mF was expected to influence the sharpness of the interfaces by changing the rise and

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fall times of the deposition pulses. The GMR values measured at 8 kOe were around 5% but their evolution with the capacitance value attached could not be considered as a systematic change. This is probably due to the overwhelming influence of the exchange reaction controlling the overall layer deposition process in the G/G mode. More recently, Podlaha and coworkers [131,133] have also investigated ED Co–Cu/Cu multilayer films obtained in G/G mode from a sulfate bath at a pH of 3 with ionic ratio Co2+/Cu2+ = 100. The multilayers consisting of 1000 bilayers were deposited on Si(1 0 0)/Ti(1 0 0 nm)/Cu(20 nm) substrates. At a nominal Co layer thickness of 2.5 nm and for thick Cu layers (nominal value around 3–3.5 nm), a GMR of about 5% was achieved at a saturation field less than 3 kOe. An attempt was also made [131] to control the GMR magnitude by the application of a pulse train instead of a single Co deposition pulse. With increasing cycle number from 1 to 8 within the Co pulse, an improvement of the GMR to 7% was achieved while retaining the relatively low saturation field. For 10 and 20 cycles within a Co pulse, the GMR dropped to about 4% and a similar non-monotonous GMR change within the same limiting values was obtained when varying the duty cycle within the Co pulse. In view of the pertinent exchange reaction in the G/G mode, it is very hard to rationalize the observed changes.

5.2.4.2. Deposition from sulfate/citrate bath. Ueda and coworkers [27,51,60–62,68,90] electrodeposited Co–Cu/Cu multilayer films with typically 50 bilayers from a sulfate/citrate bath containing also NaCl in G/G mode on glass substrates with evaporated Cu seed layers. The ionic ratio in the bath was fairly low (Co2+/Cu2+ = 19), resulting in a high Cu content in the magnetic layer, typically ranging from 10 to 30 at.%. An XRD study [90] revealed an fcc-(1 1 1) texture of the multilayer structure but no satellite reflections could be observed. The room-temperature MR(H) curves exhibited in most cases a high saturation field, indicating the dominance of SPM regions which form both due to the high Cu content in the magnetic layers and the often very thin (nominally 1.5 nm) magnetic layers which are further thinned by the exchange reaction during the Cu deposition pulse. A non-monotonous variation of the magnetoresistance measured at 15 kOe was found as a function of the nominal Cu layer thickness, with a maximum between 1 and 2 nm (slightly shifting with varying Cu content in the magnetic layer) and a second maximum around 3.5 nm. The GMR was found to decrease with increasing Cu content of the magnetic layers. At the first GMR maximum, the actual position of which is certainly at higher values than specified by the nominal thicknesses, the MR(H) curve indicates a fully SPM behavior without a significant GMRFM contribution whereas the second maximum may correspond to that seen also in Fig. 35. Therefore, the observed GMR dependence on Cu layer thickness cannot be considered as corresponding to the conventional oscillatory GMR behavior. Ueda and coworkers [51,60–62] have also investigated the influence of pulse rise and fall times on the GMR behavior by using trapezoidal current pulses for deposition but in most cases a strong GMR reduction was observed due to the less sharp interfaces between magnetic and non-magnetic layers when using non-rectangular deposition pulses. Pradhan et al. [144] applied a rather unusual pulse control in that Co–Cu/Cu multilayers were deposited from a sulfate/citrate bath by using the repetition of a three-pulse combination UCo/UCu/ Uo where UCo = 2.6 V, UCu = 2.0 V and Uo = 0 V are the voltages between the anode and the cathode during the respective pulses. Obviously, under such a control, the deposition potentials are not defined in a controlled manner during the pulses and the deposition conditions may be rather similar to a G/G control mode. (Here, the 0 V potential difference between the cathode and the anode does not mean a zero current since the electrode materials are different.) Deposition was carried out on Si(1 0 0) wafers covered with an electroless Cu layer at the back for contacting and with various combinations of Co and Cu as well as Co/Cu bilayers applied as seed layers on the front side for the purpose of creating a conducting substrate. The main objective of this study was to establish a fully electrochemical technology for multilayer electrodeposition on Si. Due to the relatively poor adherence of Co and Cu directly on Si, a maximum of 12 and 20 bilayers could be deposited on the Si wafer without spontaneous peeling with vertical and upward-looking horizontal cathode position, respectively. The GMR was 1% or 2% only for the vertical cathode position whereas about 6% GMR could be achieved at 10 kOe for the thicker multilayers obtained on the horizontal cathodes, although the MR(H) curves even in the latter case indicated the dominance of an SPM contribution to the GMR.

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5.2.4.3. Deposition from sulfamate/sulfate bath. Kainuma et al. [43] electrodeposited Co–Cu/Cu multilayer films with 150 bilayers from a sulfamate/sulfate bath in G/G mode on evaporated Cu layer substrate. No MR(H) curves were presented and the TMR component of the GMR was measured at 15 kOe on strip samples for Cu layer thicknesses between 2 and 5 nm without specifying the magnetic layer thickness. The GMR values varied between 3% and 7.5% with a maximum at intermediate Cu layer thicknesses and the position of the maximum slightly shifted with magnetic layer deposition current density. Fedosyuk and coworkers [21,31,32] electrodeposited Co–Cu/Cu multilayer films from a sulfamate/ sulfate bath in G/G mode on Cu, Al or Ni-P substrates. An fcc-(1 1 1) texture was revealed by XRD and from the satellite positions the authors determined a bilayer length in relatively good agreement with the nominal values. The room-temperature LMR(H) curve for a Co(0.2 nm)/Cu(1.5 nm) multilayer was nearly linear up to 12 kOe where the GMR was about 0.3%, indicating a fully SPM behavior due to the granular nature of the very thin Co layers. For a Co(2.5 nm)/Cu(1.5 nm) multilayer, the LMR(H) curve indicated a fully AMR behavior (pinholes in the non-magnetic layer provide FM coupling between the continuous magnetic layers). A GMRFM component of about 1% with hysteresis was measured for a Co(14 nm)/Cu(15 nm) multilayer and saturation was achieved at around 2–3 kOe. A hysteresis was observed and the low GMR value was due to the very high bilayer repeat length (29 nm). 5.2.4.4. Deposition from a sulfate bath with additives. Zhang et al. [123,124] investigated the GMR of ED Co–Cu/Cu multilayers prepared from an electrolyte containing Co-sulfate, Cu-sulfate, H3BO3, saccharin and Triton X-100 at a pH of 5.2. Deposition was carried out on mechanically polished Cu disk electrodes rotating at a speed of 400 rpm. The deposition current densities for the magnetic and non-magnetic layers were 353 mA/cm2 and 1.8 mA/cm2, respectively, and the multilayers contained 2000 bilayer repeats. No structural characterization of the multilayers was reported. Only the transverse MR component was measured up to 90 kOe magnetic field. First, two Co–Cu/Cu multilayers with 30 and 2 wt.% Cu in the magnetic layer were prepared, without specifying the layer thicknesses (the deposition pulse lengths were only given to be 12.5 ms and 1.54 s, respectively). The TMR(H) curves indicated a small decrease of the resistivity, not saturating even at 90 kOe where the TMR value was 1.5% and 3.5% for the two multilayers, respectively. This behavior is certainly due to a GMRSPM term and indicates the presence of a large fraction of SPM regions in the magnetic layer for both multilayers. A second set of two multilayers with the same magnetic layer compositions was also prepared by using shorter pulse lengths and, thus, lower thicknesses for both layers. The measured TMR component became even smaller (dropped to 0.9% and 1.3%, respectively, at 90 kOe) than before although there was also a change in the shape of the MR(H) curves. For thinner individual layers, a sharp magnetoresistance decrease was observed and after a break at around a few kOe, the MR(H) decreased linearly further up to the maximum applied field. In the absence of further information on the actual layer thicknesses, it is hard to assess more about the MR behavior of these multilayers. 5.2.5. GMR of ED Co/Cu multilayer films prepared by dual-bath technique Unlike the case of multilayers from a single bath, for ED multilayers prepared by dual-bath technique, the magnetic layer is not contaminated by the spacer layer material since it is deposited from a separate bath void of the ions of the non-magnetic spacer element. Cao and coworkers [120,127] electrodeposited Co/Cu multilayer films by dual-bath technique in P/P mode on Si(1 1 1) substrate. Sulfate bath containing H3BO3 was used for both Co and Cu deposition which were carried out in separate compartments at cathode potentials ECo(SHE) = 0.86 V and ECu(SHE) = 0.26 V, respectively. The XRD pattern reported for a multilayer with 40 bilayer repeats revealed an fcc-(1 1 1) texture, with the intensity of the (2 0 0) peak being about 1/4 of the (1 1 1) intensity. There were clear first-order satellites around the main (1 1 1) peak and from their positions a bilayer repeat length of 11.2 nm was established. Unfortunately, neither the nominal bilayer lengths nor the individual layer thicknesses were specified for any multilayer on which results are reported in these works. The magnetization curves, M(H), of multilayers with 20, 40, 60 and 80 bilayers exhibited typical FM hysteresis loops with coercivities around 50 Oe. Magnetic saturation was achieved in a few kOe magnetic fields and the relative remanence was typically 0.7–0.8. The reported MR(H) curve for a multilayer with 40 bilayer repeats shows complete saturation around 3 kOe with a GMR of about 50%.

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As a function of the bilayer repeat number, the GMR was found to exhibit a maximum at 60 bilayers where the GMR was 90% and it decreased rapidly for 40 and 80 bilayers. A sharp maximum was found also as a function of both the Co and Cu layer thicknesses; however, one cannot locate the position of these maxima since the layer thicknesses were specified in terms of the deposition pulse lengths. The results of this work are very hard to reconcile in lack of sufficient details of sample preparation and layer thicknesses. Nevertheless, the claimed 90% GMR obtained hereby using the conservative definition supersedes the well-known result of 50–60% GMR at room temperature for sputtered Co/Cu multilayer films [177,178,189]. For this reason, one has to consider the results of Cao and coworkers [120,127] with a similar caution to those of Bird and Schlesinger [15] since the findings of these two groups could not be reproduced by any other laboratories. Myung et al. [75] investigated in detail the GMR parameters of ED Co/Cu multilayer films prepared by dual-bath technique from unstirred electrolytes under galvanostatic control of the deposition of both the Co and the Cu layers. The substrate was not specified, the deposition current density was typically 0.5 mA/cm2. The multilayers mostly consisted of at least 200 bilayers. Since the individual layer thicknesses were fairly large in most cases (Co: 5–25 nm; Cu: 1–9 nm), the total deposit thickness was usually in the range 1–5 lm. Due to the G control mode of the Cu deposition, the exchange reaction could always take place after bringing the lastly deposited Co layer in contact with the Cu deposition bath. Therefore, the actual Cu layer thickness was definitely higher than the specified nominal value (of course, the reverse is true for the Co layers). The electrical resistivity in zero external magnetic field was usually much larger (in some cases even above 100 lX cm) than the values reported by Lenczowski et al. [17] for ED Co–Cu/Cu multilayers with comparable layer thicknesses. For Co deposition, either an acidic sulfate bath (CoSO4 + H3BO3) or an alkaline sulfate/pyrophosphate bath (CoSO4 + Na4P2O7 + NH4OH) was used. Similarly, for Cu deposition, either an acidic sulfate bath (CuSO4 + H2SO4 + NaCl + commercial brightener) or an alkaline sulfate/pyrophosphate bath (CuSO4 + Na4P2O7 + NH4OH) was used. The solution pH was adjusted (by adding H2SO4, H3PO4, NaOH or NH4OH) to <1 for the sulfate Cu bath, 3–4 for the sulfate Co bath and 8.5–9 for both alkaline baths. The current efficiency was determined to be 100% for the sulfate Cu bath, 48% for the alkaline Cu bath, 69% for the sulfate Co bath and 55% for alkaline Co bath. The M(H) curves reported for a few multilayer with about 5 nm thick Co and Cu layers exhibited hysteresis with a relative remanence of typically 0.7 and with saturation apparently achieved in the highest magnetic field applied (2.5 kOe). The corresponding MR(H) curves showed also a hysteresis with split peaks, with the peak positions corresponding to the coercive fields. MR saturation was achieved in some cases around 1 kOe whereas other multilayers did not show saturation even at 2.5 kOe. The magnetoresistance, without specifying if the LMR or TMR component was measured, was investigated for multilayers prepared with several combinations of the two kinds of Co and Cu baths as a function of the Co and Cu layer thicknesses. Data for the GMR magnitude were only presented for this series with GMR values typically between 2% and 10% and exhibited a maximum around 5 nm of either Co or Cu layer thickness. Although no direct structural studies were carried out, a correlation between structural quality and GMR magnitude can be inferred from data showing that high GMR values (10–13%) could be obtained only for multilayers with zero-field resistivity around 25 lX cm whereas the GMR diminished sharply for both smaller and higher resistivities, yielding to a low GMR of 2–3% only for multilayers with a resistivity around 100 lX cm and above. An interesting approach to the dual-bath technique was presented by Hayashi et al. [107] who used a supporting flow electrolyte (H3BO3 aqueous solution) without metal ions into which alternately a Co-containing and a Cu-containing sulfate electrolyte was injected and electrodeposition was carried out by potentiostatically controlled pulses. The deposition potentials for Co and Cu were ECo(SHE) = 0.90 V and ECu(SHE) = 0.70 V, respectively. A Si wafer with evaporated Au seed layer was used as substrate to electrodeposit Co/Cu multilayer films with 15 bilayer repeats and a lateral size of 10 mm by 10 mm. XRD revealed an fcc-(1 1 1) texture and no satellite reflections could be observed. M(H) and MR(H) curves with in-plane magnetic fields were reported without specifying the relative orientation of current and magnetic field for [Co(2.0nm)/Cu(dCu)]15 multilayers with dCu = 1.0 nm, 1.4 nm and 2.0 nm. The M(H) curves were nearly identical for the three multilayers, exhibiting a sigmoidal-type shape with low coercivity, low relative remanence (not more than 10%) and saturation field around 5 kOe. The shape of the M(H) curves was considered by the authors as

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being typical for an AF-coupled multilayer. The MR(H) curve for dCu = 1.0 nm indicated saturation around 3–5 kOe with a maximum GMR of about 6% at 5 kOe. The MR(H) curves for dCu = 1.4 nm and 2.0 nm did not reach saturation at 5 kOe (for dCu = 1.4 nm, the resistivity changed nearly linearly with magnetic field). Since for the two latter Cu layer thicknesses the GMR magnitude was 1.5% and 5%, respectively, Hayashi et al. [107] interpreted the observed GMR variation with Cu layer thickness from 1.0 to 2.0 nm that it was due to an oscillatory exchange coupling. However, one cannot accept this explanation since, then, for dCu = 1.4 nm the lower GMR ought to have been accompanied with an FM type magnetization curve with a fairly high relative remanence what was evidently not the case and the magnetoresistance to exhibit an AMR behavior. The latter feature cannot be established at all in lack of information whether the LMR or the TMR component was measured. Instead, the MR(H) curve of the multilayer with dCu = 1.4 nm indicates definitely a dominant SPM contribution to the GMR and this cannot be excluded also for the other two multilayers. The observed M(H) curves for each multilayer can also be interpreted as due to an SPM behavior. Also, the MR(H) curve at the second AF maximum (dCu = 2.0 nm) saturates in a magnetic field typically by an order of magnitude smaller than at the first maximum (dCu = 1.0 nm) for Co/Cu multilayers with oscillatory exchange coupling and GMR [177,189] whereas the relation of the saturation fields was just the opposite here [107]. 5.2.6. Influence of Cu content in the magnetic layer on GMR in ED Co-Cu/Cu multilayer films A detailed systematic study of GMR was carried out by Liu et al. [121] on ED Co–Cu/Cu multilayer films prepared by single-bath plating in G/P mode as a function of the Cu2+ ion concentration in the electrolyte and the thickness of the Cu layers. An electrolyte containing 1 M CoSO4 and 0.005–0.2 M CuSO4 was used. The magnetic layer was deposited at 70.5 mA/cm2 current density for a duration providing a thickness of 3.4 nm. The Cu layer was deposited at an optimized potential of ECu(SHE) = 0.38 V with thicknesses ranging from 0.4 to 3.5 nm. The bilayer number was varied between 265 and 460 in order to maintain a total deposit thickness of about 1.7 lm. From a knowledge of the individual layer thicknesses and by measuring the overall multilayer composition with electron probe microanalysis, the magnetic layer composition was determined. Depending on the Cu2+ ion concentration in the bath, the Cu content of the magnetic layer could be controlled from a few percent up to about 50 at.%. The multilayer deposits were removed from their amorphous ribbon substrates by mechanical stripping for the magnetoresistance studies. Both the LMR and TMR components were measured on strip samples up to 8 kOe at room temperature. The shape of the MR(H) curves was similar to those shown in Fig. 42. A decomposition of the MR(H) curves was carried out to separate the GMRFM and GMRSPM contributions. The dependence of the saturation magnetoresistance, the two GMR contributions and the MR(H) peak positions for the FM component on the Cu2+ concentration in the bath and on the Cu layer thickness of the multilayer is shown in Fig. 46. The decomposition analysis reveals a distinct behavior of the GMRFM and GMRSPM contributions with varying Cu contents in the magnetic layers. The GMRFM term hardly depends on the Cu content of the magnetic layer whereas its dependence on the Cu layer thickness is very similar to that shown in Fig. 35 for corresponding layer thicknesses (monotonous increase). On the other hand, the GMRSPM term becomes larger with increasing Cu contents in the magnetic layers. These observations can be accounted for by the limited solubility of Cu in fcc-Co at room temperature [202]. According to the data in Fig. 46c, the GMRSPM term remains fairly small (typically 1–2%) for low Cu2+ ion concentrations corresponding to less than 10 at.% Cu in the magnetic layers. Beyond this limit, fcc-Co apparently cannot maintain the excess Cu atoms in solid solution and regions with much higher Cu contents will appear which may be even non-magnetic at room temperature. This may lead to the appearance of Co-rich regions (e.g., in the form of islands) magnetically decoupled from the rest of the magnetic layer of FM character and these decoupled regions exhibit SPM behavior. Finally, conduction electron paths between FM and SPM entities result in spin-dependent scattering processes which contribute to a GMRSPM (GMRFM–SPM  GMRSPM–FM) term. Since the FM parts of the magnetic layer keep a Cu content below about 10 at.%, their magnetic properties are not influenced noticeably by the overall Cu content in the magnetic layer. The evolution of Hp with Cu layer thickness is, therefore, similar to that shown in Fig. 33b and shows a kind of saturation for the largest Cu layer thicknesses. Here, Hp (and also Hc) is governed by the thickness of the magnetic layer only which are well separated for such thicknesses of the Cu layers.

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c 16 14

LMR sat (%)

12 10

2+

-3

c(Cu ) / moldm 0.005 0.0125 0.025 0.05 0.10 0.20

GMR SPM (%, from LMR)

a

8 6 4 2 0 0.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

2+

12 10 8

c (Cu )/ mol dm-3 0.005 0.0125 0.025 0.05 0.10 0.20

6 4 2 0 0.0

0.5

1.0

8 6

3.5

150 2+

c(Cu ) / M 0.005 0.0125 0.025 0.050 0.100 0.200

100 50

0 1.0

3.0

200

2

0.5

2.5

250

4

0.0

2.0

d

2+

c (Cu ) / moldm-3 0.005 0.0125 0.025 0.05 0.10 0.20

HP / Oe

GMR FM (%, from LMR)

b 10

1.5

d (Cu) / nm

d (Cu) / nm

1.5

2.0

2.5

d (Cu) / nm

3.0

3.5

0 0.0

0.5

1.0

1.5

2.0

2.5

3.0

long. trans.

3.5

d (Cu) / nm

Fig. 46. Results of the decomposition of the MR curves for ED Co–Cu/Cu multilayers prepared with various Cu2+ ion concentrations in the bath as a function of the Cu layer thickness. (a) Saturation magnetoresistance LMRsat obtained as the sum of the saturation values of the FM and the SPM contributions; (b) saturation value of the GMRFM component (LMR); (c) saturation value of the GMRSPM component (LMR); (d) peak position Hp of the GMRFM components (both LMR and TMR). The solid lines serve as a guide for the eye only. Reprinted from Ref. [121] with permission of The Electrochemical Society.

The average magnetic moments of the SPM regions were also determined from the analysis of the MR(H) curves and the values varied between 1000 lB and 6500 lB. Their evolution with overall Cu content in the magnetic layer could be rationalized along the same line as discussed above. When the Cu content is low in the magnetic layer, the Cu atoms are fairly homogeneously distributed in the Co matrix and can be enriched at a few places only. Therefore, there are chances for the formation of relatively large SPM regions only. With increasing Cu content, more and more less-magnetic regions will occur in the magnetic layers and this results in smaller SPM regions as observed. Lenczowski et al. [17] also investigated the influence of Cu2+ ion concentration in the bath on the GMR in ED [Co(1.5 nm/Cu(5.0 nm)]50 multilayer films grown in P/P mode on Au substrates (for more preparation details, see Section 5.2.3.1). The Co2+ ion concentration was fixed at 1.2 M and the Cu2+ ion concentration varied from 2 mM to about 70 mM (this corresponds to a range from 600 to 17, respectively, for the ionic ratio Co2+/Cu2+). A very definite GMR maximum of about 13% was observed in the Cu2+ ion concentration range from 5 mM to 10 mM, with the GMR diminishing to about 8% at both ends of the Cu2+ ion concentration range investigated. The authors explained this by arguing that towards high Cu2+ ion concentrations, the amount of Cu incorporation in the magnetic layer is very high. Concomitantly, XRD indicated the loss of texture and a partial suppression of the satellite peak intensities. In addition, a SEM study revealed the presence of segregated Cu crystallites. Although MR(H) curves were not reported for their multilayers in this series, we may think of the same behavior as discussed above in connection with the multilayers of Liu et al. [121], i.e., the dominance of an SPM

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contribution to the GMR for high Cu2+ ion concentrations in the magnetic layers. The decrease of GMR for very low Cu2+ ion concentrations in the bath was attributed by Lenczowski et al. [17] to an increasingly difficult control of Cu layer deposition due to the very low growth rates (0.01 nm/s) and that, under such circumstances, for example, the presence of uncontrolled amount of oxygen in the electrolyte may also play a detrimental role in layer formation. All this may then lead to a reduction of interface quality and to an increase of the interface roughness and surface roughness, the latter feature indeed observed by SEM. In the study of Safak et al. [145] already discussed in Section 5.2.3.4, it was found that at higher pH values (3.0) the overall Cu content measured in the multilayer and, consequently, also in the magnetic layer was found to be higher than for low pH (2.0). This difference was accompanied with a reduced GMR value (by a factor of 3.5) at the same individual and total layer thicknesses. From the reported MR(H) curves [145], one can assess that the GMRSPM contribution became also more dominant for the multilayers with high pH. This is again in good agreement with the conclusions deduced from the work of Liu et al. [121]. Finally, the results of a related work by Kubinski and Holloway [266] on sputtered Co1xCux(1.5 nm)/Cu(2.0 nm) multilayer films with 0 6 x 6 0.55 should be mentioned. The GMR decreased roughly linearly from 16% to 8% when x was increased from 0 to 0.48. At around x = 0.5, the GMR sharply dropped to zero. The magnetoresistance hysteresis as measured by the MR(H) curve peak position (Hp) similarly reduced to zero for x  0.5. The variation of the MR(H) curve shape indicated that for the largest Cu contents in the Co layers, a GMRSPM term may dominate as is also the case of ED Co–Cu/Cu multilayers with high Cu contents in the magnetic layers. The relatively high Cu content at which the GMR and Hp drops to zero in the sputtered multilayers is certainly due to the fact that sputtering is a much more non-equilibrium process in comparison with electrodeposition. This means that by sputtering, a homogeneous solid solution of Cu in fcc-Co can be produced up to higher x values than by electrodeposition in which process the phase separation of the immiscible elements Co and Cu sets in at lower Cu contents of the magnetic layer. 5.2.7. Dependence of GMR in ED Co-Cu/Cu multilayer films on Cu layer thickness: evidence for a nonoscillatory GMR behavior and for the absence of coupling It was discussed in Section 2.4.2 that in most FM/NM multilayers prepared by PD techniques an oscillatory behavior of the GMR can be observed as a function of the NM spacer thickness (see Fig. 8). This oscillation is ultimately connected with the alternation of the interlayer exchange coupling as the spacer thickness varies. For Co/Cu multilayers, the coupling is AF at about 1, 2 and 3 nm Cu layer thickness whereas the coupling is FM at about 1.5 and 2.5 nm. For spacer layers yielding an AF coupling, the zero-field state has a high resistance since the AF coupling forces the adjacent magnetic layer magnetizations to align antiparallel to each other. For spacer thicknesses resulting in FM coupling, all the layer magnetizations are aligned parallel even in zero field and in such a case, the application of an external magnetic field cannot improve the magnetization alignment; therefore, no GMR effect will occur. It must be clear from the above description that the oscillatory GMR is related to the GMRFM term only which arises from electron scattering events along an electron path ‘‘FM layer 1 ? NM spacer layer ? FM layer 2”. In fairly perfect multilayers, the GMRFM term is the sole contribution to the GMR and most PD multilayers exhibit this behavior. It happens, however, quite often that the multilayer is not perfect in the sense that the magnetic layer contains also some regions exhibiting SPM behavior which give a GMRSPM contribution to the observed magnetoresistance. As mentioned in Section 2.4.4, ED multilayers are especially prone to the formation of SPM regions under certain deposition conditions. Therefore, if we are looking for an oscillatory GMR behavior in ED multilayers we must ascertain that we can either suppress the GMRSPM term by appropriate control of the deposition process or unambiguously separate the GMRFM term from the measured MR data. We have selected those reports on the GMR of ED Co–Cu/Cu multilayers where it can be established that the GMR data correspond to the GMRFM term as discussed above. This was the case with the results of Lenczowski et al. [17] and Li et al. [131] and with three works of the present authors [121,139,147] and all these data are displayed in Fig. 47. Although the magnitude of GMR varies from study to study (probably due to differences in actual layer thicknesses, preferred texture, substrate

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10

GMRs (%)

8

6

[147]

4

[121] [139] 2

[17] [131]

0 0

1

2

3

4

5

6

dCu (nm) Fig. 47. Evolution of the saturation value of the GMRFM component in ED Co–Cu/Cu multilayers with Cu layer thickness dCu. The numbers in [] indicate literature data references. It is noted that, according to a study reported in Ref. [128], in ED Co–Cu/Cu multilayers obtained under galvanostatic control, the actual layer thicknesses for the samples from Ref. [131] can be by as much as 1 nm higher than the original nominal thicknesses specified. Reprinted from Ref. [147] with permission, copyright (2009) of The American Physical Society.

material and other details of the electrodeposition process), the general trend is that (i) a clear GMR effect develops above a certain Cu layer thickness of about 1 nm only, (ii) the GMR magnitude increases monotonically with dCu and (iii) a saturation or maximum occurs for Cu layer thicknesses around and above about 4 nm. A few further data of Lenczowski et al. [17] and Liu et al. [109] omitted from Fig. 47 show a qualitatively similar behavior (see Figs. 35b and 33a, respectively). A monotonic GMR increase was observed also by Shima et al. [80] up to about 5 nm Cu layer thickness but their MR(H) curves indicate that saturation has not been achieved up to the maximum magnetic field applied (see Fig. 37) and, therefore, their GMR values may contain an SPM contribution as well. The GMR results of Kainuma et al. [43] and Myung et al. [75] also reveal a maximum-like behavior in the same range as for the data displayed in Fig. 47 but these authors have not shown MR(H) curves at all and in this manner we have no information on the eventual importance of a GMRSPM term. The general conclusion is that the GMRFM term in electrodeposited Co/Cu multilayers does not exhibit an oscillatory behavior as a function of the Cu layer thickness in those cases where we can unambiguously identify the appropriate GMRFM contribution characteristic for FM/NM multilayers. On the other hand, there are several papers reporting on an ‘‘oscillation” of the GMR magnitude for electrodeposited Co/Cu multilayers. As to the results by Jyoko et al. [34] for dCu = 0.9 nm, they report an MR(H) curve (see Fig. 44b) which unambiguously reveals a dominant SPM contribution. Therefore, the high GMR value at this Cu layer thickness cannot originate from an AF-coupled state and this conclusion is further supported by the M(H) curve reported for this sample [34] since it shows a large remanence whereas the AF-coupled states should exhibit a low remanence. For a higher dCu value (3.25 nm), their multilayer displays an MR(H) curve (cf. Fig. 44b) typical for ED Co–Cu/Cu multilayers with such high Cu layer thicknesses (split MR peaks, low saturation field) and have a similar evolution of the GMR magnitude as shown in Fig. 47. The results of Ueda et al. [60] show the same features: the non-saturating MR(H) curves obtained for Cu thicknesses around their first observed GMR maximum (dCu  1.5 nm) are dominated by an SPM term whereas split MR peaks with low saturation fields appear for Cu layer thicknesses around 3–4 nm. It should also be noted that their first GMR maximum appears roughly at a Cu layer thickness where usually FM coupling is observed for physically deposited Co/Cu multilayers. Actually, their Cu layer thicknesses are probably even higher than the specified

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values as a consequence of the applied galvanostatic deposition technique since a significant exchange reaction takes place during the Cu deposition pulse as pointed out, e.g., in Refs. [99,128]. Based on the argumentation presented above, we can conclude that in the above described two works [34,60] no GMR oscillation corresponding to that observed in physically deposited Co/Cu multilayers can be identified. The first report of GMR oscillation in electrodeposited multilayers was by Bird and Schlesinger [15] who even fitted their ‘‘oscillatory” GMR data for Co/Cu and Ni/Cu by an RKKY function. However, no details including MR(H) curves were presented in that short communication. Furthermore, the GMR magnitude was reported to be as high as in the corresponding sputtered counterparts which results have, however, never been reproduced by other researchers. For this reason, we have to treat these results with appropriate caution. There are still two further reports [114,127] which also claim to have observed GMR oscillations in electrodeposited Co/Cu multilayers. However, in lack of sufficient details about sample preparation and MR measurements, we cannot conclude about the reliability of these data. If the SPM region fraction is fairly low (at most a few percent of the total magnetic material), then the magnetic and magnetotransport behavior of the FM/NM multilayer system will depend on whether the spacer layer material is continuous or it contains a high density of pin-holes. In the following, we shall discuss electrodeposited Co/Cu multilayers with a negligible SPM fraction only. Below a thickness of typically 1 nm, the Cu spacer layers in electrodeposited multilayers usually contain a large density of pin-holes which provide a direct FM exchange coupling between adjacent layer magnetizations. In such a case, bulk-like FM behavior with an AMR effect occurs due to a percolation of the magnetic layers via the pin-holes in the Cu layers. The same effect is observed if the Cu layer thickness fluctuation is significant and at some regions the very small spacer thickness enables a direct FM exchange coupling. With increasing average thickness, the continuity of the Cu layers increases and the reduced density of pin-holes as well as the improvement of Cu layer thickness uniformity weakens the FM exchange coupling between the magnetic layers which, thus, become gradually uncoupled. The uncoupled layer magnetizations will be randomly aligned and electron transitions between nonaligned adjacent layers can yield a larger and larger GMR effect as observed. At sufficiently large Cu layer thicknesses (around 3–4 nm), when the magnetic layers become completely uncoupled, there is no more increase in the randomness of the magnetization alignments and the GMR reaches saturation in parallel with the saturation of the coercive force, with the value of the latter becoming characteristic for thin individual magnetic layers. Since the relative remanence of the magnetization of both AMR and GMR multilayers was found to be practically the same [17,147], we have to conclude the absence of a significant AF coupling in the latter ones. Beyond a certain spacer layer thickness, we have to expect a reduction of the GMR due to a decrease of the number of the magnetic/non-magnetic interfaces per unit thickness (dilution effect). A decrease of GMR is also expected when exceeding the Cu layer thickness through which the spin-memory is no longer preserved for the conduction electrons what is another pre-requisite for the observation of the GMR. It should be noted that Shima et al. [80] suggested a surface roughness model for electrodeposited Co/Cu multilayers in order to explain the absence of oscillatory behavior. In this model, the authors have assumed that a Néel-type ‘‘orange-peel” coupling of magnetostatic origin provides a strong enough FM coupling to overcome the existing AF exchange coupling between adjacent magnetic layers. Even if this mechanism can explain the absence of GMR maxima at the expected positions of the AF coupling, in the intermediate Cu thickness regions the addition of the magnetostatic FM coupling to the FM exchange coupling would then provide a strong resulting FM coupling, i.e., a diminished GMR. By contrast, the GMRFM data displayed in Fig. 47 from various reports show a uniquely monotonous increase of the GMR magnitude (at least up to the maximum beyond 3 nm Cu thickness), thus questioning the validity of the model of Shima et al. [80]. The explanation proposed for explaining the evolution of the magnetoresistance in ED Co–Cu/Cu multilayers with spacer layer thickness [147] is not restricted to systems produced by this method. Parkin and coworkers used very similar arguments for explaining the occurrence of GMR in lack of AF coupling [190] in sputtered Co/Cu multilayers or the increase of the GMR with Cu layer thickness in some MBE-grown Co/Cu multilayers [190].

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5.2.8. Temperature dependence of GMR in ED Co–Cu/Cu multilayer films When summarizing the results of GMR studies on ED Co–Cu/Cu multilayer films in previous sections, it was already noticed that the GMR increases towards lower temperatures. For example, Chassaing [49] reported that GMR values as high as 25% could be obtained at T = 4.2 K which diminished to 15% at T = 77 K (no GMR data were given for room temperature in this work). As discussed in Section 2.1, the magnetoresistance is defined as the ratio DR/Ro (see Eq. (1)) where Ro is the zero-field resistance and DR is the resistance change upon the application of an external magnetic field. For sputtered [Co(1.5 nm)/Cu(dCu)]30 multilayers with dCu = 0.9 nm and 1.8 nm, Kubota et al. [267] reported a strong increase (by a factor of about 1.75) of the GMR from room temperature to the liquid helium range. At the same time, the value of Ro reduced with decreasing temperature by a factor of about 0.85 only. Therefore, the increase of the GMR towards lower temperatures is mainly due to the variation of the term DR. It was pointed out by Kubota et al. [267] that the observed temperature dependence of GMR for Co/Cu multilayers can be quantitatively well explained by the reduction of spin-fluctuations at low temperatures as suggested by the theory of Hasegawa [268]. The present authors [132] performed a detailed study of the temperature dependence of GMR for ED Co–Cu/Cu multilayers between 300 K and 12 K. Two particular multilayers were selected for this study: one with dominating GMRFM (sample V6) and another one with dominating GMRSPM contribution (sample V4). The main objective of this study was to explore the temperature evolution of the two kinds of GMR contribution and, especially, to see how their ratio changes with temperature. The two multilayers containing 300 bilayers each were deposited by two-pulse plating from a single pure sulfate bath in G/P mode on rough Ti substrate at a non-optimized Cu deposition potential ECu(SHE) = 0.01 V [99,128]. Since at this potential the exchange reaction is active during the Cu deposition pulse, the actual layer thicknesses were determined from an analysis of the overall multilayer composition [99,128] to be Co(2.0 nm)/Cu(1.6 nm) for sample V4 and Co(2.1 nm)/Cu(3.8 nm) for sample V6. The Cu content of the magnetic layers was about 5 at.% in both multilayers. More details of the sample preparation and characterization including structural studies were described in Refs. [99,128]. The longitudinal MR(H) curves are shown in Fig. 48 for both multilayers at various temperatures. In agreement with the results of Kubota et al. [267] on sputtered Co/Cu multilayers, the GMR increases significantly towards lower temperatures also for the ED Co–Cu/Cu multilayers. A particular observation is that the shape of the MR(H) curve remained practically unchanged for the whole temperature range of investigation. The usual GMR decomposition analysis [103] was carried out for the measured MR(H) curves shown in Fig. 48. The saturation values of the total LMR (MRs) and of its FM and SPM components are shown as a function of temperature in Fig. 49a. Apart from the GMR magnitude, the observed temperature evolution of the GMRFM term shows very good agreement with the results of Kubota et al. [267], especially for sample V6 in which the GMRSPM term was fairly small.

Fig. 48. Temperature dependence of the longitudinal magnetoresistance for sample V4 (a) and sample V6 (b). Reprinted from Ref. [132] with permission, copyright (2006) of the American Physical Society.

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Fig. 49. Temperature dependence of various magnetoresistance and magnetic parameters determined for both samples V4 and V6: (a) total (saturation) longitudinal magnetoresistance MRs and its FM and SPM contributions; (b) fractional SPM contribution to the saturation magnetoresistance MRs and saturation magnetization Ms; (c) average magnetic moment l of SPM regions (the inset shows the same data as the main frame except that they are plotted on a logarithmic vertical scale in order to better demonstrate that the ratio of the average SPM moments for the two samples V4 and V6 is fairly independent of temperature); (d) magnetoresistance peak position Hp and coercive field Hc. Lines in (c) are only a guide for the eye to follow the trend of the SPM moment as derived from magnetoresistance data. Reprinted from Ref. [132] with permission, copyright (2006) of the American Physical Society.

The data in Fig. 49a reveal a strong difference in the temperature evolution of the GMR for the two selected multilayers. Whereas at room temperature the total magnetoresistance values are nearly the same for both samples, the increase of the magnetoresistance towards low temperatures is much larger for sample V6. The decomposition analysis furthermore reveals that in both samples the FM and SPM components of the GMR increase in line with each other. As a consequence of this behavior, the relative weight of the SPM contribution to MR hardly changes with temperature as shown in Fig. 49b. This trend leads one to the conclusion that the relative importance of the various spin-dependent scattering events is independent of temperature. Fig. 49c shows the temperature dependence of the average moment l of an SPM region. For a given temperature, the average SPM moment as obtained from the MR measurements by assuming noninteracting SPM regions is larger for sample V6 than for sample V4. The ratio of the moments established for samples V4 and V6 is fairly independent of temperature, as seen from the data in the inset of Fig. 49c. The two samples exhibit a similar behavior in that l decreases strongly with decreasing temperature as observed also in MBE-grown Co/Cu multilayers [198,199]. The temperature evolution of the MR peak position (Hp) is shown for both samples in Fig. 49d. An approximately linear increase of the peak position with decreasing temperature can be seen. Magnetic measurements on the same two multilayer samples gave very similar temperature dependence of the parameters as was deduced from the MR(H) curves (see Fig. 49c for the average SPM moment and Fig. 49d for a comparison of Hc and Hp values).

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It is mentioned that Sasaki et al. [113] also reported a systematic evolution of the MR(H) curves as a function of temperature on ED Co–Cu/Cu multilayers with a dominating GMRSPM contribution. The temperature variation of the total GMR measured was qualitatively very similar to that of Ref. [132] but no separation of the individual GMR terms was carried out. According to the results described above on the basis of the work by the present authors [132], ED Co–Cu/Cu multilayers V4 and V6 exhibit very different GMR behavior in spite of the fact that their room-temperature GMR values as measured in a magnetic field of 8 kOe are very similar. The MR data indicate a dominance of the SPM contribution to the GMR for sample V4 in contrast to sample V6 in which the FM contribution dominates. On the other hand, it was also found [132] that in the total magnetization the SPM contribution differed only slightly for the two samples, amounting to about 20% for both multilayers (cf. Fig. 49b). The size of the average magnetic moment l of SPM regions agreed well for sample V4 when the values derived from magnetic and magnetoresistance data are compared (Fig. 49c). On the other hand, for sample V6 the l values from MR data are higher and those from the magnetic data are smaller than the corresponding values of sample V4. An explanation of the observed differences should be sought in the details of the underlying electrochemical processes governing the morphology of the multilayer growth. All the preparation parameters of samples V4 and V6 were identical except for the nominal Cu layer thickness. For multilayer V4 (nominal magnetic/nonmagnetic layer thickness: 3.4 nm/0.2 nm), the Cu layer with a final thickness of 1.6 nm is built up almost exclusively at the expense of the dissolution of the previously deposited magnetic Co–Cu layer (the latter being reduced to an average thickness of about 2 nm). Such a consumption of the magnetic layer occurs randomly over the cathode area and this leads to a strong fluctuation of the magnetic layer thickness. At the extreme, such fluctuations may lead to a discontinuity of the magnetic layer at some places and, thus, SPM regions can form which are magnetically decoupled from the FM part of the magnetic layers. This tendency may be further enhanced as a result of some unavoidable degree of intermixing due to the random Co-dissolution and Cu-deposition processes which may take place during almost the total duration of the Cu deposition pulse. These features may well explain the dominance of the GMRSPM term in the magnetoresistance of multilayer V4. This picture is in accordance with the discussion in Section 5.2.7 in that in electrodeposited Co–Cu/Cu multilayers the Cu layers are not continuous below a certain thickness and this also strongly contributes to a fragmentation of the magnetic layers into SPM type regions. Thus, the preparation conditions of sample V4 having an average Cu layer thickness of 1.6 nm promote the formation of a granular type quasi-multilayer structure in which both FM and SPM regions can occur in the magnetic layer. The situation is quite different for multilayer V6 for which the nominal magnetic/nonmagnetic layer thicknesses are 3.4 nm/2.5 nm and the actual layer thicknesses are 2.1 nm/3.8 nm. Again with reference to Section 5.2.7, for large Cu layer thicknesses as is the case here the Cu layer is definitely continuous and this favors the formation of continuous magnetic layers even at thicknesses where their fragmentation occurs at thinner Cu layers. Therefore, the occurrence of dispersed SPM regions in the magnetic layers is then less favored what is reflected in the fact that the GMRSPM contribution is much smaller for sample V6 than for multilayer V4. Also, the Cu layer thickness is much larger for sample V6 which may well lead to a ‘‘leveling” effect and, thus, each subsequent Co layer may start to grow on a much smoother surface and may become more uniform. It will, thus, be less inhomogeneously consumed during the Cu deposition pulse than for multilayer V4. These considerations are further supported by the coercive field data. For a bulk Co95Cu5 alloy obtained by direct-current electrodeposition at the same current density as used for the magnetic layer in the multilayer structure, the room-temperature coercive field was found to be about 30 Oe [132]. In view of this result, the much higher coercive field values for sample V6 (Fig. 49d) indicate the presence of continuous and well-separated thin magnetic layers in the multilayer structure as well-known for thin ferromagnetic films. On the other hand, the typically 50% lower coercive field values for sample V4 with an effective (average) magnetic layer thickness identical to that of sample V6 hint at larger effective magnetic layer thicknesses, at least in some areas, for multilayer V4. This suggests the presence of larger magnetic regions percolated over several layers through the pinholes in the discontinuous Cu layers.

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For both Co–Cu/Cu multilayers described in this study, a significant increase of the MR values with decreasing temperature can be observed whereas the overall shape of the room-temperature MR curves is retained. This is in contrast to the case of ED Ni–Cu/Cu multilayers for which even the shape of the MR curves changed significantly with temperature [104]. However, a discussion of the difference between the observed behaviors of the two system will be deferred to Section 5.7.2 where similar data for the Ni–Cu/Cu multilayers are presented. 5.3. Co–Ag/Ag multilayer films There are relatively few reports on the GMR study of ED multilayer films [40,60,67,112,148]. This may partly be explained by the difficulties in preparing good-quality multilayers in this system by electrodeposition and partly by the discouraging results on PD Co/Ag multilayers [269–271] which were found to exhibit much smaller GMR in comparison with PD Co/Cu multilayers [177,178,188– 190]. The different equilibrium crystal structures of Co (hcp) and Ag (fcc) metals accompanied with a large atomic size difference (the latter results in a lattice mismatch Da/a = 14% between fcc-Co and fcc-Ag) are very unfavorable factors for growing defect-free multilayers in this system. The insufficient structural quality is certainly one of the major reasons for the low GMR even in PD Co/Ag multilayers. First, we shall review reports in which attempts were made to fabricate Co–Ag/Ag multilayers by electrodeposition [40,60,67,112,148,247]. Valizadeh et al. [247] used a bath based on cobalt chloride or cobalt sulfate and silver cyanide (AgCN) and prepared Co-Ag(5nm)/Ag(5nm) multilayers in G/G mode on Cu foil substrates, with the magnetic layers containing a few percent of Ag only. A cross-sectional TEM study revealed a multilayered structure although both the Co and Ag layers were nanocrystalline. The XRD pattern reported indicated the presence of fcc-Ag and fcc-Co or hcp-Co layers (the only XRD peak that can be assigned to Co was around the expected position of the hcp-Co(0 0 2) and fcc-Co(1 1 1) peaks which nearly completely coincide with each other, see Ref. [149]). No sign of multilayer satellite reflections were revealed. The broad XRD lines were also indicative of a finegrained structure. Neither magnetic nor magnetotransport data were reported in the work of Valizadeh et al. [247]. Ueda and coworkers [40,60] used a sulfate-based bath with sodium citrate addition to prepare multilayers in the Co–Ag system by electrodeposition in the G/G mode on glass plate with an evaporated Cu substrate layer. The composition of the FM/NM layers was reported to be Co70Ag30/Co8Ag92 but the indication of the composition was not accompanied by any structural characterization. The magnetization and the magnetoresistance were studied as a function of the thickness of the Co-rich and the Ag-rich layers. The GMR measured at room temperature in a magnetic field of 21 kOe increased from about 5% to 9% with the Co-rich layer thickness between 0.4 nm and 1.6 nm, respectively. When the GMR was studied as a function of the Ag-rich layer thickness in the 0.3 nm to 1.8 nm range, a GMR maximum of about 9% was found at about 1.2 nm. This GMR maximum increased to about 13% at 5 K [60]. Subsequently, Ueda and coworkers [112] also reported on the electrodeposition of multilayers in the Co–Ag system in the G/G mode from a bath containing CoSO4, AgI and KI. Under the preparation conditions used, the magnetic layer had the same composition as in their former works [40,60] but here the non-magnetic layer composition was not specified and no structural results were reported either. For these multilayers, the GMR measured under the same conditions as previously [40,60] exhibited a maximum as a function of both the magnetic and the non-magnetic layer thickness, with both thicknesses varying from a few tenths of a nanometer to about 3 nm. The maximum GMR achieved was about 6–7%, with the maximum being between 0.5 nm and 1.5 nm for both types of layer and the maximum position varied slightly with the thickness of the other layer. It should be noted that Ueda and coworkers [60] reported MR(H) curves also for Co–Ag alloy films obtained by d.c.-plating which were actually granular alloys. Since the shape of the MR(H) curves and the magnitude of the GMR were very similar for their granular alloys and multilayers, it cannot be excluded that their multilayers [40,60,112] were to a large extent actually also granular type alloys. The layer compositions for their multilayers (each type of layer containing a large amount of the other constituent metal) also suggest this since these elements are highly immiscible [202]. A further hint for

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the granular character of their multilayers is the fairly large GMR magnitude [40,60,112] in comparison with the small (about 1% or less) GMR in ED Co–Ag/Ag multilayers reported in the two works [67,148] to be described below. Fedosyuk et al. [67] reported the electrodeposition of multilayers in the Co–Ag system in the G/G mode from a bath containing CoSO4, AgNO3 and KI on amorphous Ni-P substrates. The magnetic layer contained some 10–15 at.% Ag and XRD revealed the presence of fcc-Ag and fcc-Co phases in the asdeposited state for magnetic layer thicknesses up to about 100 nm. A GMR of about 0.7% was measured at room temperature in a magnetic field of about 8 kOe for a multilayer with unspecified layer thicknesses. The MR(H) curve had a single peak at zero field and this behavior changed to a symmetrically split double-peaked curve with Hp  3 kOe after annealing at 400 °C for 30 min whereby the maximum GMR value reduced by about a factor of two. Although no XRD patterns were presented, these authors claimed that the heat-treated sample exhibited satellite diffraction peaks around the fcc-Co reflections from which the bilayer repeat period was deduced and found to be in good agreement with the nominal bilayer thickness. Actually, it is very hard to rationalize the results obtained after the heat-treatment (the appearance of XRD satellites and the unusually high value of the MR(H) peak position). In a most recent work by Garcia-Torres et al. [148], the GMR was studied for ED Co–Ag/Ag multilayers prepared in the G/P mode from quiescent electrolytes containing Co(ClO4)2, AgClO4 and NaClO4 and prepared with analytical grade chemicals and ultrapure water. The pH of the electrolytes was between 2.0 and 2.5, depending on the concentration of the metal salts. Electrodeposition was performed in a tubular cell with an upward looking cathode at the bottom of the cell as suggested in Refs. [99,138]. A Si(1 0 0)/Cr(5 nm)/Cu(20 nm) substrate was used, with the adhesive Cr layer and the Cu seed layer prepared by evaporation on the Si wafer. The number of bilayer repeats was varied in a manner as to maintain a nearly constant total multilayer thickness of about 800 nm. As a first step, it was established at which current densities and bath compositions a magnetic layer with an AMR of about 1% and with a low saturation field (<300 Oe) can be obtained by galvanostatic d.c.-plating. Under these conditions, the magnetic layer contained 1–2 at.% Ag only. Next, the appropriate Ag deposition potential for avoiding both Co dissolution and Co codeposition was optimized by using the current transient analysis (as suggested in Ref. [231]) and the optimum was found to be EAg(SHE) = 0.25 V. A structural study of Co–Ag/Ag multilayers obtained under these optimized conditions with various layer thicknesses was carried out by XRD. The XRD patterns did not reveal satellite reflections and indicated the presence of an fcc-Ag structure with strong (1 1 1) texture. Apart from an extremely weak hcp-Co(0 0 4) diffraction peak, no other Co diffraction lines could be indentified separately since the expected position of the hcp-Co(0 0 2) and fcc-Co(1 1 1) peaks was overwhelmed with the still fairly intense fcc-Ag(2 0 0) peak. Nevertheless, all the XRD peaks were much narrower than those reported by Valizadeh et al. [247] for comparable layer thicknesses, indicating significantly larger grain sizes. High-resolution TEM together with selected area electron diffraction (SAED) was used [148] to gain additional information about the crystal structure. The analysis of the SAED pattern indicated the presence of fcc-Ag and hcp-Co. The SAED pattern also revealed that no fcc-Co phase was present since the most important reflections (1 1 1) and (2 0 0) were not detected. The room-temperature magnetoresistance measurements of ED Co-Ag/Ag multilayers with layer thicknesses ranging from 2 to 10 nm yielded MR(H) curves with splitting and not saturating up to 8 kOe, the maximum field applied. The GMR magnitude measured for both the LMR and TMR components was around 1% at 8 kOe. The usual decomposition procedure [103] revealed that the FM and SPM components to the GMR were of comparable size. The GMR exhibited a shallow maximum around dAg = 6 nm as a function of the Ag layer thickness for dCo = 3 nm. The temperature dependence of the GMR for a Co–Ag(3 nm)/Ag(6 nm) multilayer exhibited very similar features as reported for ED Co–Cu/ Cu alloys (see Section 5.2.8). According to the analysis of the results on the ED Co–Ag/Ag multilayers discussed above, one can conclude that neither the Co-rich magnetic layers, nor the non-magnetic Ag layers can be considered as fully continuous, although there are definitely some regions of the magnetic layers which exhibit FM behavior. Although the GMR magnitude is usually higher in PD Co/Ag multilayers [269–271], there are strong similarities with the ED Co–Ag/Ag multilayers. It was found by van Alphen and de Jonge [270] on sputtered Co/Ag multilayers that for a Co-layer thickness of 2 nm which can be considered as continuous,

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the room-temperature GMR increased from about 1% to 7% when increasing the Ag layer thickness from 0.5 nm to 4 nm, after which a slight decrease of the GMR with Ag layer thickness was observed. On the other hand, for a Co layer thickness of 0.6 nm, the GMR values rose quickly to about 8% by around 1 nm Ag layer thickness and showed a slight decrease beyond about 3 nm. The higher GMR values for the Co/Ag multilayers with a 0.6 nm Co layer were attributed to the granular nature of the magnetic layer. For a Co layer of 2 nm thickness, the multilayers are of the normal FM/NM type although the authors concluded from the magnetization data that no antiferromagnetic coupling between the FM layers exists. Similarly high GMR values (typically between 5% and 10%) were obtained also by Araki [269] at room temperature on MBE-grown Co/Ag multilayers with 0.6 nm thick Co layers when the Ag layer thickness was varied between 1 nm and 4 nm; nevertheless, the GMR observed should at least partly arise from a granular type behavior due to the very thin Co layer, as it can be judged from the high saturation field of the reported MR(H) curves. This is also supported by the fact that, for an Ag layer thickness of 2.5 nm, the GMR reduced strongly for thicker Co layers. For Co layer thicknesses around and above 2 nm (i.e., for Co layers with clear FM behavior) the GMR was already as small as about 1% only. Furthermore, the lack of a discernible GMR oscillation with spacer layer thickness seems to be also a common feature for multilayers in the Co–Ag system prepared by any preparation method. The size difference and immiscibility of the constituent metals as mentioned at the beginning of this section can be considered as the major reasons for the low structural quality and inferior GMR behavior of these multilayers. 5.4. Co–Au/Au multilayer films Ueda and coworkers [112] used a bath containing CoSO4, KAu(CN)2, Na3-citrate and NaCl to prepare ED Co95Au5/Au multilayers in the G/G mode on glass plate with an evaporated Cu substrate layer. The current density values used varied in the range 0.1–25 mA/cm2 although it was not specified which values were used for depositing the individual layers. No MR(H) curves were reported just the room-temperature GMR values measured at H = 21 kOe were displayed. Similarly to the results on the ED Co–Ag/Ag multilayers reported in the same work (see Section 5.3), the GMR values exhibited a maximum as a function of both the Co–Au and the Au layer thickness in the range 0.5–2.5 nm. The maximum GMR values varied between 2% and 5%, being definitely smaller by about a factor of two in comparison with the ED Co–Ag/Ag multilayers. In lack of any structural characterization, the strong similarity of the magnetic and magnetoresistance data to the ED Co–Ag/Ag–Co and the ED Co–Au/Au multilayers [112] might suggest that the GMR originates also for the latter system from a granular magnetic behavior instead of a clear multilayer GMR effect. 5.5. Co–Ru/Ru multilayer films ED Co98.5Ru1.5/Ru multilayers were prepared [83] in the G/G mode on polished Ti substrates by using a bath at 50 °C which contained CoSO4, RuCl3, Na2SO4 and HCOOH. The nominal layer thicknesses were varied between 0.2 and 1.2 nm for Ru and 3.5 and 7 nm for Co–Ru. Due to the unavoidable exchange reaction in the G/G mode, the actual Ru layer thicknesses are definitely larger than the nominal values but no estimate of the layer thickness changes was made. TEM SAED patterns indicated a fully hcp structure and the presence of superlattice reflections in these multilayers. The layered structure could be well observed from cross-sectional, partly high-resolution TEM studies and a columnar growth with a large density of lattice defects (twinning and stacking faults) was also revealed. By measuring the room-temperature longitudinal and transverse MR(H) curves, it was found that the AMR behavior at low Ru thicknesses gradually transformed into a GMR type behavior with the increase of the Ru layer thickness. At sufficiently large Ru layer thickness, both the LMR and TMR components were negative (Fig. 50), indicating a clear GMR effect in spite of the small GMR magnitude. Measurement of the LMR component at 4.2 K up to 30 kOe confirmed the room-temperature result. It is not clear whether the non-saturating magnetoresistance at low temperature is due a SPM behavior or arises as a consequence of a strong AF coupling in the Co/Ru multilayers. The original motivation for this work on ED Co–Ru/Ru multilayers was that in previously studied Co/Ru multilayers [175,272–274] or sandwiches [273,275–277] prepared by sputtering or evaporation

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Fig. 50. Room-temperature and low-temperature MR(H) curves of an ED Co(3.5 nm)/Ru(1.2 nm) multilayer with GMR behavior. Layer thicknesses are nominal (see text). Reprinted from Ref. [83] with permission of The Electrochemical Society.

the GMR was always found to exhibit quite small values (typically 0.1% at room temperature), independently of the applied deposition technique and the applied deposition conditions. This finding was in contrast with expectation on the basis of the resistivity change caused by Ru impurities diluted in a Co matrix [169] from which a high GMR is expected for Co/Ru multilayers. The surprisingly small observed GMR value was interpreted as a consequence of a strong intermixing at the interfaces in the Co/Ru multilayers during the layer growth process [275]. The idea then arose that the relatively lowtemperature electrodeposition process might help to avoid such an intermixing but this attempt was not successful. This is because the magnetic layer dissolution, at least when the G/G method is applied, can apparently lead also to a significant interfacial intermixing: as was shown above, the GMR values were comparable for the multilayers prepared by either PD or ED methods. It was pointed out in Ref. [83] that the strong reduction of the GMR as a consequence of interfacial intermixing in Co(Ru)/Ru multilayers prepared by any method has its roots in electronic structure effects by Ru impurities in a Co matrix. First, it should be considered that according to the results of Ref. [278], a Co98.5Ru1.5 dilute alloy exhibits a much smaller AMR than pure Co metal although the saturation magnetization of Co reduces upon alloying by 1.5 at.% Ru in proportion with the alloying element concentration only [83,278]. Apparently, the presence of a few percent of Ru atoms in bulk hcp Co induces electronic band structure changes and modifies the spin asymmetry of the up and down dbands at the Fermi level, which then leads to a decrease of spin-dependent scattering effects. This change in the asymmetry between majority- and minority–spin electronic states at the Fermi level has been well demonstrated for Co92Ru8 and Co84Ru16 bulk alloys by electronic band structure calculations [277]. The change in the calculated electronic band structure may be consistent with the observed small saturation magnetization reduction mentioned above. Taking into account that there is a significant chemical intermixing at the interfaces of Co/Ru multilayers prepared by any method indicates that a Co–Ru alloy with strongly reduced AMR exists between the magnetic Co [or Co(Ru)] layers and the non-magnetic Ru layers in these multilayers. Therefore, the formation of such an intermixed interface with strongly reduced spin-dependent scattering properties can be a possible reason for the very small GMR observed in Co/Ru multilayers for any preparation technique. A microscopic mechanism of this effect could be that in such an interfacial region the population of the spin-up and spin-down conduction subbands equalizes due to the reduced spin asymmetry at the Fermi level, and therefore, the spin polarization cannot be transferred between the neighboring magnetic layers. 5.6. Co–Zn/Cu multilayer films El Bahraoui and coworkers [106,117] used a dual-bath technique to produce ED Co–Zn/Cu multilayers with 10–20 bilayer repeats on glass or Si substrates which were covered by a 240 nm thick

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Cu buffer layer sputter-deposited at room temperature. The 12 nm thick magnetic Co-Zn layer was deposited from a bath containing CoCl2, CoSO4, ZnSO4 (NH4)2SO4, and H3BO3; the deposition current density or deposition potential was not specified. The 3 nm thick Cu layer was deposited at a current density of about 20 mA/cm2 from a bath containing CuSO4 and H2SO4. When the magnetic layer composition was close to Co5Zn95, XRD revealed, in addition to the Bragg peaks of the monoclinic CoZn13 phase, a multilayer peak with weak first-order satellites. Since the multilayer was found to be magnetic and the majority-phase CoZn13 structure is not expected to be magnetic, there should be some small magnetic precipitates (possibly tiny Co-rich particles) embedded in the majority phase which, however, escape the conventional XRD study due their small volume fraction and small individual size. In this manner, one should conclude that the magnetic layer in the Co5Zn95/Cu multilayer is itself a granular magnetic alloy containing small SPM regions. The small observed coercive field (<50 Oe) and the small relative remanence (ca. 0.1) is consistent with this picture. The shape of the reported room-temperature MR(H) curve is indicative of a classical granular magnetic alloy but the observed GMR is extremely small (partly due to the large shunting effect of the buffer layer) and does not saturate at the maximum field (15 kOe) applied. It may well happen that the observed GMR is not at all due to the multilayer structure revealed by XRD but rather due to the granular nature of the magnetic layer itself. This means that the majority of the spin-dependent scattering events giving rise to the GMR effect observed derive from electron paths not passing through the Cu spacer layer at all. 5.7. Ni-Cu/Cu multilayer films It was shown by Kubota et al. [189] that the GMR of sputtered Ni/Cu multilayers is by nearly by an order of magnitude smaller than for corresponding Co/Cu multilayers. This may explain the fact that much less efforts were devoted to the study of GMR in ED Ni–Cu/Cu multilayers [13,15,18,29, 37,59,70–72,82,89,97,102,104,108,110] in comparison with the wealth of literature on ED Co–Cu/Cu multilayers (see list of references in Section 5.2.1). It should be noted at the outset that the optimization of the Cu deposition potential has not yet been carried out for ED Ni–Cu/Cu multilayers and, therefore, for this system the actual layer thicknesses may deviate from the nominal values. For this reason, in this section we shall use mainly the nominal layer thicknesses except when otherwise stated. As a general remark, it is noted that the dissolution of the magnetic layer during the Cu deposition pulse is a less serious problem for Ni than for Co. On the other hand, the non-optimized Cu deposition potential can lead to a significant codeposition of the magnetic element in the Cu spacer also in ED Ni– Cu/Cu multilayers. 5.7.1. Deposition from a single sulfamate/sulfate bath in P/P mode As it was already noted in Section 4.2, the first ED multilayers with sufficiently high structural quality (as indicated by the appearance of multilayer satellite reflections in the XRD pattern) were Ni–Cu/ Cu multilayers produced by Lashmore and Dariel [256]. These authors used a bath containing Ni-sul-

Table 5 Deposition parameters of ED Ni–Cu/Cu multilayer films with GMR which were obtained under P/P control from a single sulfamate/ sulfate type bath with buffering agents X (H3BO3, together with NaOH or HCl for pH regulation) and additive A. The cathode potentials applied for deposition are referenced to the standard hydrogen electrode (SHE), see Table 1. Authors and Refs. in []

X+A

Ni-sulfamate + CuSO4 Lashmore et al. [13] Bird and Schlesinger [15] Shima et al. [37] Myung and cow. [70,71] Tokarz and cow. [89,97] Alper et al. [102] Kazeminezhad and Schwarzacher [108]

H3BO3 Antipitting agent + H3BO3 H3BO3 H3BO3 + NaOH or HCl H3BO3 H3BO3

pH

Ni2+/Cu2+ ionic ratio

ENi(SHE)/ECu(SHE)

4.0 3.5 2–3 2.0

120 8000 100 94 150 50 46

1.36 V/0.06 V 0.95 V/+0.07 V 0.95 V/+0.45 V 0.96 V/0.46 V 1.06 V/0.66 V 1.46 V/+0.04 V 1.66 V/0.16 V

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famate, CuSO4 and H3BO3 and many authors have applied such a sulfamate/sulfate type bath to produce ED Ni–Cu/Cu multilayers for studying their GMR behavior [13,15,37,70,71,89,97,102,108]. Some details of the preparation conditions used in these studies are listed in Table 5. Lashmore et al. [13] investigated the GMR of ED Ni–Cu/Cu multilayers prepared in P/P mode with a total thickness of about 2 lm on annealed pc-Cu(1 0 0) substrates. The deposition conditions were not described in detail in this report; therefore, the parameters specified by Lashmore and Thompson in Ref. [236] were included in Table 5. The XRD study revealed an fcc multilayer structure with pronounced (1 0 0) texture, i.e., it was identical with that of the substrate. Sharp and intense satellite reflections around the main (2 0 0) Bragg peak were reported for a Ni–Cu(2.0 nm)/Cu(2.5 nm) multilayer for which even the (1 1 1) peak exhibited satellites. Superlattice satellites, though much less intense, could be seen around the (2 0 0) peak also for multilayers with 1.7 nm and 0.6 nm nominal Cu layer thicknesses. On the other hand, satellite peaks were not observed for multilayers with nominal magnetic layer thicknesses below 2 nm. The bilayer periods deduced from satellite peak positions agreed within 5% with the nominal values determined from the deposition parameters. After dissolving the Cu substrates, both the LMR and TMR components of the magnetoresistance were measured on strip-shaped samples. Multilayers with 2 nm nominal thickness of the magnetic layers and with the Cu layer thickness varying from 0.8 nm to 4.3 nm were found to exhibit MR(H) curves characteristic for well-defined FM/NM multilayers (cf. Fig. 3): the magnetoresistance decreased sharply for both the L and T components up to about 1 kOe beyond which a slight, nearly linear MR decrease was only observed for higher magnetic fields. The MR values at 1 kOe were around 2% and this can be ascribed to a GMRFM term whereas the linear term (a few tenths of a percent change from 1 to 8 kOe) is due to the paraprocess characteristic for any ferromagnetic metal beyond technical saturation [160,164–166]. The isotropic MR data obtained at the maximum applied field by Lashmore et al. [13] on a series of ED Ni–Cu(2 nm)/Cu(dCu) multilayers as a function of the Cu layer thickness are shown in Fig. 51. The data indicate an oscillatory type behavior of the GMR with maxima roughly at the positions corresponding to those reported for sputtered Ni/Cu multilayers [189]. For assessing the validity of this observed oscillatory behavior, it is important to consider the shape of the MR(H) curves as well. Of the series shown in Fig. 51, Lashmore et al. [13] reported the MR(H) curve for the Ni–Cu(2 nm)/Cu1.9 nm) multilayer only but the Cu layer thickness dependence of the full width at half maximum (FWHM) of the MR(H) curves was also presented. It can be inferred from this latter that around the second GMR maximum (dCu  2 nm), the FWHM value was about 50–100 Oe. For dCu = 0.8 nm (first GMR maximum) and dCu = 3.5 nm (third GMR maximum), the FWHM values were about 150 Oe and 50–100 Oe, respectively, whereas the FWHM values are much higher for Cu layer thicknesses corresponding to the GMR minima (550 Oe and 250 Oe). The narrow MR(H) curves at the GMR maximum positions indicate the dominance of a GMRFM term which can be saturated in fairly

Fig. 51. Isotropic magnetoresistance measured at H = 8 kOe for an ED Ni–Cu/Cu multilayer series with 2 nm magnetic layer thickness as a function of the Cu spacer thickness. The solid line is a guide to the eye only. Reprinted from Ref. [13] with permission of The Electrochemical Society.

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small magnetic fields (1 kOe). On the other hand, the maxima of the saturation field of the AF coupling in sputtered Ni/Cu multilayers [189] coincide with the positions of the GMR maxima as a function of dCu, a feature just the opposite to the finding of Lashmore et al. [13]. Therefore, we have to conclude that although a GMRFM term provides the dominant contribution to the observed GMR in these ED Ni–Cu/Cu multilayers, the reported oscillatory behavior cannot be considered as originating from an oscillatory exchange coupling. The possible origin of the observed GMR vs. dCu variation can have its roots in the microstructural changes occurring when varying the Cu layer thickness and this should still be explored. Evidently, in the ED Ni–Cu/Cu multilayers the Cu layers can grow in a continuous form down to lower thicknesses than was found for the ED Co–Cu/Cu multilayers (compare Figs. 47 and 51) since in the latter system a GMRFM term arises for dCu values well above 1 nm only. Another issue is the actual value of the spacer layer thickness (dCu). As noted above, there is no direct information to what extent the Ni–Cu and Cu layer thicknesses might have changed with respect to the nominal values for the multilayers studied by Lashmore et al. [13] due to the non-optimized Cu deposition potential. An indication for this may be taken, however, from the fact that a Ni–Cu(1 nm)/ Cu(1 nm) multilayer exhibited a dominating SPM type GMR contribution since the LMR and TMR values which were about 0.2% and 0.4%, respectively, at H = 5 kOe did not show any sign of saturation at this maximum applied field. This result indicates that the magnetic layer deposited during the highpotential pulse may have dissolved to the extent that it is no longer continuous but rather breaks up and contains also small SPM regions (on the other hand, the MR(H) curves did not indicate the appearance of SPM regions for a 2 nm thick magnetic layer). This hints at an increase of the Cu layer thickness over the nominal value set during the deposition process. Therefore, the actual thicknesses in Fig. 51 may be larger than the values displayed and, thus, the true GMR peak positions should certainly be situated at higher dCu values. These considerations are further supported by the XRD data discussed beforehand for the same samples. Namely, it was noted by Lashmore et al. [13] that no multilayer satellite peaks could be observed for nominal magnetic layer thicknesses below 2 nm. This was attributed to the less-perfect structure the cause of which, of course, can well be the magnetic layer dissolution, leading to a non-uniformity of the layer thickness as a result of which the structural coherence along the thickness should get lost. We should make a final remark concerning the GMR magnitude in ED Ni–Cu/Cu multilayers. Namely, the GMR increases strongly when Co or Fe is alloyed into the magnetic layer of Ni/Cu multilayers [189]. On the other hand, due to the anomalous codeposition mode of the iron-group metals (Fe, Co and Ni) with each other (see Section 3.2.1), if there is a small amount of Co or Fe in the chemicals serving as the source of Ni for Ni–Cu/Cu multilayer electrodeposition, a significant amount of Co or Fe can be codeposited into the magnetic layer (the contamination of Co and Fe in a Ni deposit is usually by an order of magnitude higher than in the bath [72]). This possible contamination should be taken into account when comparing the GMR values of ED Ni–Cu/Cu multilayers from various reports. Since the GMR magnitude of ED multilayers is usually smaller than that of the corresponding PD multilayers, the maximum GMR value of 3.5% (Fig. 51) reported by Lashmore et al. [13] for ED Ni–Cu/Cu multilayers can be considered as a good typical value in view of a GMR of 7% observed by Kubota et al. [189] for sputtered Ni/Cu multilayers. Bird and Schlesinger [15] reported on the GMR of ED Ni–Cu/Cu multilayers prepared in P/P mode from a sulfamate/sulfate bath under identical conditions as used for their ED Co–Cu/Cu multilayers (see Section 5.2.3.4) except for the deposition potentials which are given in Table 5. Similarly to the Co–Cu/Cu multilayers, an oscillatory GMR behavior was reported with a maximum GMR of 7%, i.e., identical to that obtained on sputtered Ni/Cu multilayers [189]. In lack of a detailed description of preparation conditions, structural characterization and MR(H) curves, one cannot really assess the validity of these data, similarly to the case of ED Co–Cu/Cu multilayers studied by Bird and Schlesinger [15]. Myung and coworkers [70,71] reported on the GMR of ED Ni–Cu/Cu multilayers prepared in P/P mode from a sulfamate/sulfate bath on a Si-wafer covered with a sputtered Au layer. The layered structure was demonstrated for a Ni–Cu(60 nm)/Cu(5 nm) multilayer by taking a SEM picture of the deposit cross-section after selectively etching the Cu layers. The XRD pattern reported for a multilayer with unspecified layer thicknesses showed separate Cu and Ni peaks for both the (1 1 1) and (2 0 0) reflections, the latter ones being of somewhat higher intensity for Ni. The split Cu and Ni peaks

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indicate fairly large individual layer thicknesses. No further structural characterization was reported. No MR(H) curves were reported although it was noted that the maximum field applied (1.5 kOe) was not sufficient to saturate the magnetoresistance. The TMR component was measured in the maximum magnetic field for two multilayer series: (i) Ni–Cu(dNi–Cu)/Cu(2.7 nm) with 5 nm 6 dNi–Cu 6 27 nm and (ii) Ni–Cu(5 nm)/Cu(dCu) with 0.5 nm 6 dCu 6 2.7 nm. The MR values of these ED Ni–Cu/Cu multilayers varied in the range from 1% to 2.2% whereas the authors also reported a magnetoresistance value of 1.4% for a d.c.-plated Ni foil (bulk Ni). Since the TMR component was only measured, on the basis of the close proximity of the bulk Ni and the multilayer values, it is highly probable that there was no GMR contribution to the observed magnetoresistance in the multilayers; a magnetoresistance of GMR origin can be concluded only on the basis of the LMR component (cf. Fig. 3). Tokarz and coworkers [89,97] investigated the GMR of ED Ni95Cu5/Cu multilayers prepared in P/P mode from a sulfamate/sulfate bath on a Si(1 0 0) wafer covered with an evaporated Cu(50 nm) layer exhibiting a (1 1 1) texture. The Ni95Cu5/Cu multilayers with typically 1 0 0 bilayer repeats exhibited comparable fractions of (1 1 1) and (1 0 0) oriented crystallites. Clear first-order satellites were observed around the (2 0 0) reflections and some very weak second-order satellites were also present. From the LMR(H) and TMR(H) curves reported [97], we can establish an AMR behavior for a [Ni95Cu5(4 nm)/Cu(0.65 nm)]100 multilayer. This is probably due to the discontinuous Cu layers at a such a small thickness which allows a FM coupling of adjacent magnetic layers via pinholes in the spacer layer. A GMR effect was observed for a [Ni95Cu5(4 nm)/Cu(2 nm)]50 multilayer for which the LMR and TMR components reached a value of 0.3% and 1.6%, respectively, in a magnetic field of about 2.5 kOe whereas further the magnetoresistance did not show a significant change up to 15 kOe. Tokarz and coworkers [89] reported also the temperature dependence of the magnetoresistance for several Ni95Cu5/Cu multilayer series from 300 K to 3.5 K (without specifying if LMR or TMR was measured). A minimum in the magnetoresistance was found for each sample at a different temperature but no explanation was offered for the existence of such a minimum which is hard to understand. In the same work, the magnetoresistance was reported at room temperature for two Ni95Cu5(4 nm)/Cu(dCu) multilayer series with either 50 or 100 bilayer repeats. The MR values scattered from 0.5% to 1.7% but in lack of knowledge whether the LMR or TMR component was displayed, one cannot even be sure if the observed magnetoresistance was due to a GMR effect since bulk Ni itself can have a magnetoresistance of this size (see the result quoted from the work of Myung and coworkers [70] above). Alper et al. [102] investigated the GMR of ED Ni–Cu/Cu multilayers prepared in P/P mode from a sulfamate/sulfate bath with pH values ranging from 2.0 to 3.0 on pc-Cu(1 0 0) sheets. The total multilayer thickness was 2 lm for the XRD studies and about 1 lm for the MR measurements, the latter carried out after dissolving the substrates from the multilayer deposits. XRD revealed a strong (1 0 0) multilayer texture with clear first-order satellites and with a weak (1 1 1) reflection. The bilayer periods deduced from the superlattice peak positions were in fairly good agreement with the nominal values. It should be noted, however, that no satellites were observed for Cu layer thicknesses below 2 nm which was attributed to the relatively large total multilayer thickness (2 lm). For so large thicknesses, a significant surface and interface roughening may develop. The necessity of a relatively large Cu layer thickness for the satellite peak observation is an indirect evidence that roughening takes place mainly during the Cu layer deposition step. For the MR studies, ED Ni–Cu/Cu multilayers were prepared with fixed nominal thickness (1.5 nm) of the magnetic layer and the Cu layer thickness was varied from 0.4 nm to 2.0 nm. A dominating AMR effect was observed for dCu < 0.6 nm. This can be ascribed to a percolation of the Ni–Cu layers via pinholes in the discontinuous Cu layers. For dCu P 0.7 nm, a dominating GMR effect was observed (LMR < 0, TMR < 0). However, the MR(H) curves did not saturate for magnetic fields up to 8 kOe. This was ascribed to the presence of SPM regions in the magnetic layer. Since the nominal magnetic layer thickness is fairly small (1.5 nm), a partial dissolution due to the non-optimized Cu deposition potential may lead to the formation of SPM regions. Another finding of this work was that the GMR magnitude decreased from 2.8% to 1.4% when increasing the bath pH from 2.2 to 2.9 for identical layer thicknesses. The electrolyte pH influences the cathode surface during deposition as well as the layer growth mode, thus the interface quality and grain size. Another possible explanation was also given for the pH dependence of the GMR on the basis of chemical analysis results on both the multilayers and d.c.-plated Ni-rich alloys. It was found, namely, that from the electrolyte used the Cu content of the magnetic layer increases with increasing pH. The Cu content in

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the magnetic layer was about 10, 20 and 30 at.% for pH values of 2.0, 2.5 and 3.0, respectively. For pure Ni metal and Ni–Cu alloys with low Cu content, there is a significant asymmetry of the spin-up and spin-down d-bands at the Fermi level and this leads to strong spin-dependent scattering events which then give rise to the GMR effect. In Ni–Cu alloys with high Cu contents, this Fermi level asymmetry is destroyed and this results in a diminished GMR effect. Kazeminezhad and Schwarzacher [108] investigated the GMR of ED Ni–Cu/Cu multilayers prepared in P/P mode with 200 bilayer repeats from a sulfamate/sulfate bath with a low pH of 2.0 on Au-coated glass substrates. Neither structural data nor MR(H) curves were reported, only the LMR and TMR values measured at H = 3.5 kOe were displayed as a function of the magnetic layer thickness for dCu = 1.0, 1.5 and 2.0 nm. For each dCu, the multilayers exhibited GMR when the magnetic layer thickness did not exceed about 2.5 nm beyond which the LMR component took positive values. This fact indicated that, here, the AMR effect of the thick magnetic layers dominates over the GMR effect arising from the multilayer structure. In the GMR thickness range, the LMR values were typically around 0.5% and the TMR values varied between 0.5% and 1.0%. Shima et al. [37] investigated the GMR of about 0.6 lm thick ED Ni–Cu/Cu multilayers prepared in P/P mode from a sulfamate/sulfate electrolyte on annealed pc-Cu(1 0 0) substrates. The high ionic ratio Ni2+/Cu2+ = 100 ensured that the magnetic layer contained less than 5 at.% Cu. The multilayers had a quasiperiodic structure which was created by forming a so-called Fibonacci sequence by appropriate combinations of two building blocks. Block A was a bilayer Ni(11 ML)/Cu(9 ML) and block B was a bilayer Ni(11 ML)/Cu(5 ML) where 1 monolayer (ML) corresponds roughly to 0.2 nm layer thickness. For comparison, two regular, periodic multilayers were also prepared, one by using a periodic repetitions of block A and another one with block B. According to an XRD study, the quasiperiodic multilayer exhibited a strong (1 0 0) texture with four satellites around the main peak, corresponding to the quasiperiodic structure. The two periodic multilayers displayed the same texture with clear first-order satellites around the (2 0 0) peak. The presence of a small amount of (1 1 1)-oriented grains was also indicated by the XRD patterns. The superlattice peak positions corresponded well to the nominal layer thicknesses except for the Ni layers for which the difference was ascribed to an eventual slight dissolution of the magnetic layers during the Cu deposition pulse. For the MR measurements, the Cu substrate was dissolved from the multilayers. The MR(H) curves presented without specifying if the LMR or TMR component was measured indicated a non-saturating behavior up to the maximum field applied (2 kOe) with MR values between 0.2% and 0.3% for all three multilayers. In spite of the relatively good structural quality of the multilayers investigated, as a consequence of the very low MR values, no definite conclusions can be drawn concerning the effect of the quasiperiodic structure on GMR. The non-saturating character of the MR(H) curves can certainly be ascribed to the SPM regions formed during the partial dissolution of the fairly thin (1.5 nm) magnetic layers. 5.7.2. Deposition from a single sulfate/citrate bath in G/G mode The present authors [18,29,59,72,82,104] investigated the GMR of ED Ni–Cu/Cu multilayers prepared in G/G mode from a bath containing NiSO4, CuSO4, Na3-citrate and NaCl. Although due to the low ionic ratio Ni2+/Cu2+ = 7.5 of the electrolyte used the magnetic layer composition in these studies was Ni81Cu19, this still enabled the observation of GMR since the Curie point of the magnetic layer was close to 200 °C [72,279] (since Ni and Cu are completely soluble in each other in the fcc structure [202], phase separation problems discussed in Section 5.2.6 for Co–Cu/Cu multilayers do not arise here even at such a high Cu content of the magnetic layer). Electrodeposition was carried out on mechanically polished Ti plate substrates from which the multilayer deposits could be peeled off and the structural studies as well as the magnetic and transport property measurements were performed on the self-supporting multilayers having a typical total thickness of several micrometers. The results of XRD and TEM studies were already summarized in Section 4, by giving also some details on the correlation between deposition conditions, microstructure and GMR. A fairly good multilayered structure was established with large grains (columns) in cases where GMR could be observed. The magnetoresistance was studied, by measuring the field dependence of both the LMR and TMR components, for several ED Ni–Cu/Cu multilayer series with a range of layer thicknesses (1 nm 6 dNi–Cu 6 12 nm; 0.5 nm 6 dCu 6 3 nm). Due to the G/G method applied, the actual Cu layer thicknesses were definitely larger than the nominal values and the reverse is true for the Ni–Cu layers. This is caused by the

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exchange reaction discussed in Section 3.2.2. In Ref. [72], an important piece of evidence for the exchange reaction was provided by delicate EQCM data. In the same report, it was also shown for the first time that by applying a two-pulse plating sequence with zero Cu deposition current, the exchange reaction results in such a thick Cu layer between the magnetic layer deposited during the applied pulses that the deposit obtained clearly exhibits GMR, i.e., a multilayer structure is formed (a similar finding was described in Section 5.2.4.1 for ED Co–Cu/Cu multilayers [77]). The zero-field resistivity was also found to be larger in such deposits than the value for the corresponding d.c.-plated bulk Ni–Cu layer and was similar to the data reported for related ED Ni–Cu/Cu multilayers [279]. The MR(H) curves of those ED Ni–Cu/Cu multilayers [59,72,82,104] which exhibited GMR behavior (i.e., LMR < 0; TMR < 0) were characterized, similarly to those investigated by Lashmore et al. [13] by a rapidly decreasing MR component for magnetic fields up to about 1 kOe beyond which only a small magnetoresistance change occurred up to 8 kOe (Fig. 52a [59]). The MR contribution saturating in low magnetic field can be unambiguously identified as a GMRFM term. As a function of the Cu layer thickness, the GMR exhibited a fairly broad single maximum-like behavior as shown in Fig. 52b [59]. A similar evolution of GMR was reported in another work [82] although the maximum was less well defined and the GMR magnitude was smaller by about a factor of two. In the latter work, some possible origins of the irreproducibility were also discussed such as the importance of the cathode position and orientation as well as the temporal instability of the sulfate/citrate bath applied. Another source of different magnitude of GMR in ED Ni–Cu/Cu multilayers can be the various level of Fe and Co impurities of starting chemicals used for the Ni–Cu bath preparation. In any case, these GMR vs. dCu data show a quite different behavior than that reported by Lashmore et al. [13] (compare Figs. 51 and 52b). Evidently, further studies by using an optimized Cu deposition potential are needed to reveal the true variation of GMR with dCu in ED Ni–Cu/Cu multilayers. Nevertheless, the general evolution of GMR in Fig. 52b is very similar to the data reported on ED Co–Cu/Cu multilayers (see Figs. 35 and 47). As to the dependence of GMR on magnetic layer thickness, it was found [18,29,82] that the largest GMR can be obtained around a magnetic layer thickness of about 2–3 nm (in these studies, the Cu layer thickness was fixed in the range from 0.7 nm to 2 nm). For low magnetic layer thicknesses, the saturation field of the magnetoresistance strongly increased and, quite often, no saturation occurred up to 8 kOe, a feature observed also by Lashmore et al. [13] in some cases. This was due to the appearance of SPM regions as a consequence of the exchange reaction which can ‘‘consume” the previously deposited magnetic layer to the extent that the residue of it contains already magnetically isolated regions exhibiting SPM behavior as discussed also for ED Co–Cu/Cu multilayers in Section 5.2. For magnetic

Fig. 52. (a) Longitudinal (L) and transverse (T) magnetoresistance of Ni–Cu(3 nm)/Cu(dCu) multilayers with dCu = 1.5 nm (full symbols) and 2 nm (open symbols) at 300 K. (b) Longitudinal (L) and transverse (T) magnetoresistance of Ni–Cu/Cu multilayers measured at H = 8 kOe as a function of the spacer layer thickness (dCu) at 300 K. The solid line is a guide to the eye only. Reprinted from Ref. [59] with permission of Elsevier.

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Fig. 53. Temperature dependence of the longitudinal magnetoresistance curves of a Ni–Cu(3 nm)/Cu(2 nm) multilayer. Reprinted from Ref. [104] with permission of Elsevier.

layer thicknesses beyond the GMR maximum, the AMR effect within the magnetic layer started to dominate and the relative importance of the multilayer GMR effect gradually diminished. The overall behavior of GMR as a function of the magnetic layer thickness corresponds well to that found in ED Co–Cu/Cu multilayers (see Figs. 36a and b). The non-saturating MR component was measured [104] for some ED Ni–Cu/Cu multilayers up to a magnetic field of 18 kOe and saturation still did not occur at room temperature, giving a suggestion for the SPM origin. In order to prove this, detailed low-temperature MR(H) measurements were carried out for a multilayer with fairly strong non-saturating behavior at room temperature. It turned out that at low temperatures (Fig. 53) the MR(H) curves saturated in magnetic fields as low as 5 kOe or even smaller. This could be considered as a strong evidence for the SPM origin of the non-saturating character of MR(H) curves in ED Ni–Cu/Cu multilayers. A similar lowering of the MR(H) saturation fields of ED Ni–Cu/Cu multilayers with decreasing temperature was found also by Tokarz et al. [89]. We may recall now the related results discussed for ED Co–Cu/Cu multilayers in Section 5.2.8 according to which for this system the shape of the MR(H) curves hardly showed any changes when reducing the temperature. The most probable reason for this difference is that for the Ni–Cu/Cu system all the SPM regions become FM regions at the lowest measuring temperature whereas for the Co–Cu/Cu system some regions always remain in the SPM state. In the case of the Ni–Cu/Cu multilayers, the occurrence of a SPM–FM transition (blocking) at low temperature has at least two reasons. On the one hand, by cooling the sample, the measuring temperature becomes lower than the blocking temperature of some SPM regions (this effect can equally work for both Ni–Cu/Cu and Co–Cu/Cu multilayers). On the other hand, however, one has to take into account also a possible PM–FM transition of the boundary separating the SPM and FM regions. By comparing the composition dependence of the Curie temperature of the Ni–Cu and Co–Cu alloys, one obtains that dTC/dc = 11.8 K/at.% for the Ni–Cu system [202,279] while for metastable Co–Cu alloys the same data is 30.8 K/at.% [280]. This means that the composition interval that can undergo a PM–FM transition during the same temperature change is 2.6 times higher for the Ni–Cu system than for the Co–Cu system. It can also be assumed that for Ni–Cu samples the local composition may vary continuously because of the miscibility of the components. In contrast, the very limited miscibility in the case of the Co–Cu system may result in precipitates, and the boundary regions are probably so Cu rich that they cannot undergo a PM–FM transition at any temperature. Hence, a PM–FM transition of the boundary regions occurring with temperature in the Co–Cu system is much less likely than for the Ni–Cu system.

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It was already noticed above that the sulfate/citrate bath used by the present authors [18,29,59,72,82,104] for the preparation of ED Ni–Cu/Cu multilayers with GMR is temporally unstable. This is probably due to the relatively low pH of the bath since it was shown [281] that a sulfate/citrate bath for Ni–Cu alloy deposition can be stabilized for pH approaching values as high as 6. This was achieved by adding NaOH to the bath containing NiSO4, CuSO4 and Na-citrate. The preparation of ED Ni–Cu/Cu multilayers in the G/P and G/0/G mode (the zero in the second pulse sequence stands for a zero-current pulse) from such a stabilized sulfate/citrate electrolyte in an electrochemical flow reactor was described in detail [263], by estimating even the layer thickness changes on the basis of overall multilayer analysis data (layer thickness changes ranging from 0.5 nm to 1.5 nm were found, depending on the specific deposition conditions). The magnetoresistance was also studied for the same ED Ni–Cu/Cu multilayers by Meuleman et al. [110] but it was obtained that the bath stabilization had a very detrimental effect on the GMR magnitude. The multilayers with 100 bilayers were deposited on glass/Ti(18 nm)/Au(30 nm)/Cu(28 nm) substrates where the Ti and Au layers were obtained by sputtering and the Cu layer by electrodeposition for the sulfate/citrate bath immediately prior to the multilayer deposition. XRD patterns revealed strong (1 1 1) and (2 2 2) reflections of comparable intensity, just as the Au sublayer, indicating an epitaxial growth on gold. From the width of the diffraction lines, the grain size was estimated to be fairly small (10–15 nm). In spite of the small grain size, first-order satellites could be revealed around the (2 2 2) reflection for several multilayers in both series (deposited in either G/P or G/0/G modes), at least for Cu layer thicknesses above about 2.5 nm. The bilayer repeats derived from the XRD satellite peak positions agreed well with the nominal bilayer repeat periods. For multilayers with GMR behavior (i.e., LMR < 0 and TMR < 0) which was the case for sufficiently thick (>2 nm) Cu layers, the MR(H) curves did not saturate up to 8 kOe and their shapes indicated a dominant SPM type contribution. Also, the total GMR measured at the maximum field applied did not exceed 0.35% for any of the multilayers investigated. This was certainly to a large extent due to the extremely small crystal size which, on the other hand, is a consequence of unfavorable multilayer formation conditions from the stabilized sulfate/citrate bath of high pH. The less suitable character of the stabilized bath for good deposit microstructure formation showed up already in the MR behavior of the d.c.-plated Ni98Cu2 alloy which material serves as the magnetic layer in the multilayers. The LMR(H) and TMR(H) curves of this alloy revealed a strong magnetic anisotropy and an AMR magnitude of 1.2% only whereas the AMR of pure Ni is 2.4% at room temperature [164]. For comparison, it is noted that a d.c.-plated Ni81Cu19 deposit from a non-stabilized sulfate/citrate bath of low pH exhibited an AMR of 1.2% whereas no magnetic anisotropy was indicated by the MR(H) curves [72]. 5.7.3. Deposition by dual-bath techniques In the work of Myung et al. [71] already discussed in Section 5.7.1, results were given also for ED Ni/Cu multilayers prepared by two different dual-bath approaches. In the first case, multilayers were prepared by transferring the substrate between two different electrolytes. The Ni plating solution contained NiCl2, NaCl, H3BO3 and saccharin, whereas the Cu plating solution contained CuSO4, H2SO4 and brightener; for both types of layers, a constant current density of 5 mA/cm2 was applied. The authors chose these plating baths on the basis of experience that they were found to yield smooth and contiguous deposits even at nanoscale thicknesses. In the second approach, a recirculating electrochemical flow reactor was used with three reservoirs (magnetic metal electrolyte, non-magnetic metal electrolyte, deionized rinse water). Ultra-high purity Ar gas (99.999%) was used to sparge each reservoirs and to facilitate flow of the plating solutions and rinse water through the reactor. The plating sequence was as follows: one kind of layer was electrodeposited by recirculating its plating solution through the reactor. When the deposition was completed, a valve was switched to enable deionized water pumped through the reactor in order to displace the remaining plating solution. During the rinsing step, a low cathodic current was applied to the substrate to preclude oxidation of the deposited film. After the rinsing step, the other plating solution was recirculated through the reactor and the second film was electrodeposited. This was followed by another rinsing step and, in this manner, multilayer films were deposited by repetitive cycling. The ingredients of the plating solutions and the deposition current/potential were not specified though they may be assumed to have been identical with those used in the classical static dual-bath

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method described above. As for the ED Ni–Cu/Cu multilayers deposited from a single bath (see Section 5.7.1), the MR measurements were carried out also for the dual-bath ED Ni/Cu multilayers in a magnetic field of 1.5 kOe which was not sufficient to saturate the magnetoresistance; similarly, it was not specified either here if the LMR or TMR component data are presented. By both dual-bath methods, ED Ni/Cu multilayers were first prepared with dNi = 57 nm and the MR(1.5 kOe) values were displayed as a function of the Cu layer thickness. For the static dual-bath multilayers, the magnetoresistance varied in the range between 1.5% and 2.5% when dCu increased from 0 up to 27 nm. For the multilayers from the electrochemical flow reactor, the MR(1.5 kOe) values were presented in the range 0 6 dCu 6 7 nm and were around 1.5% for dCu = 0, 2 nm and 7 nm, whereas they were much larger at intermediate thicknesses (3% at 3 nm and 5% at 5 nm). These data indicate a maximum-like behavior for Cu layer thicknesses around 3–6 nm. The magnetic layer thickness dependence of the magnetoresistance was also studied for both preparation methods by using a constant Cu layer thickness of 5 nm with Ni layer thicknesses up to 60 nm. The reported MR(1.5 kOe) values ranged typically from 1% to 4% in both cases and it is hard to really assign any particular systematic layer thickness dependence to the data. On the basis of experience with other ED Ni/Cu–Cu multilayers [72,82] having magnetic layer thicknesses as high as about 10 nm, it is hard to imagine to obtain so high GMR values as reported by Myung et al. [71] here. Namely, for magnetic layer thicknesses above 10 nm, the AMR effect should already dominate and for pure Ni at room temperature, the value of either of the LMR and TMR components is typically around 1% (however, a few percent of Co or Fe contamination can raise the magnetoresistance to the values reported here). 5.8. Co–Ni–Cu/Cu multilayer films An intensive study of GMR has been carried out on ED Co–Ni–Cu/Cu multilayers [912,16,23,28,38,39,43,44,48,50,54,56,66,69,81,86,87,105,122–124,138,142]. This wide interest was motivated mainly by the fact that whereas the addition of Ni to the magnetic layers of Co/Cu multilayers prepared by PD methods was found to reduce the GMR [189], the deleterious exchange reaction and magnetic metal dissolution is less effective in the case of using Ni as magnetic metal when preparing GMR multilayers by electrodeposition as compared to the case of Co. The first observation of GMR in ED multilayers which was made in the Co–Ni–Cu/Cu system [9] was achieved by finding a successful compromise between reduced GMR and better controlled electrochemistry during multilayer preparation. Similarly to the ED Co–Cu multilayers, various bath combinations were elaborated to prepare ED Co–Ni–Cu/Cu multilayers for GMR studies and the review of results will be organized mainly along this line. Following the recipe of the first successful preparation of an ED multilayer film with GMR behavior from a sulfamate/sulfate bath [9], it became popular to work with some variants of such an electrolyte for the production of ED Co–Ni–Cu/Cu multilayers although there have also been reports without using a sulfamate component. 5.8.1. Deposition from variants of sulfamate/sulfate type baths in P/P mode 5.8.1.1. Deposition from a bath containing Ni-sulfamate, Co-sulfate and Cu-sulfate. Most GMR studies on ED Co–Ni–Cu/Cu multilayers were carried out on samples prepared in P/P mode from an electrolyte containing Ni-sulfamate, CoSO4, CuSO4 and H3BO3 [9–12,16,23,28,38,39,50,54,81,87]; it is noted that in Refs. [39,50], multilayers prepared in G/G mode were also studied for purposes of comparison. In the pioneering works of Alper and coworkers [9,10,28], GMR was reported on ED Co–Ni-Cu/Cu multilayers prepared on pc-Cu(1 0 0) plates from such an electrolyte at deposition potentials ECo–Ni–Cu(SHE) = 1.56 V and ECu(SHE) = +0.09 V [in Ref. [28], ECu(SHE) was +0.04 V]. The substrates were dissolved after deposition and the multilayers containing typically 100 bilayers were mounted on a glass plate for further studies. Compositional analysis showed [9,10,28] that the amount of Co was about 3–4 times higher than Ni in the magnetic layers which also contained Cu up to 25 at.% or even more. XRD studies revealed that the multilayers inherited the texture of the pc-Cu(1 0 0) substrate and first-order satellites could clearly be observed [28,81] which yielded bilayer lengths in fairly good agreement with the nominal values. The magnetoresistance was measured in the van

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Fig. 54. (a) Longitudinal magnetoresistance curve of an ED Co–Ni-Cu(2 nm)/Cu(0.7 nm) multilayer; (b) the evolution of the LMR () and TMR (s) components measured at 8 kOe for an ED Co-Ni-Cu/Cu multilayer series with constant magnetic layer thickness and a constant total multilayer thickness of about 300 nm. Reprinted from Ref. [10] with permission, copyright (1994) of American Institute of Physics.

der Pauw geometry (cf. Fig. 1a). A typical MR(H) curve recorded for the longitudinal component is shown in Fig. 54a and a very similar curve was measured for the transverse component with somewhat larger magnitude. These curves are characterized by a saturation field well above 8 kOe and with a small splitting only. An evolution of the magnitude of the LMR and TMR components measured at 8 kOe is shown in Fig. 54b as a function of the Cu layer thickness. The magnitude of the observed magnetoresistance and the common (negative sign) of the LMR and TMR components indicates that we have to deal here with a GMR effect. The shape of the reported MR(H) curves, at least at small Cu layer thicknesses (1 nm or less) [9,10,28,81] strongly resembles the MR(H) curve shapes shown in Fig. 48a which could be decomposed into a GMRFM and a GMRSPM term. Therefore, the MR(H) curves in Fig. 54a certainly indicate the presence of a large amount of SPM regions. Although a direct Langevin type fitting was not carried out but from the strongly nonsaturating character (Hs well above 8 kOe) and the small splitting, it can be concluded that the measured GMR is definitely dominated by the GMRSPM term. This is probably due to the fact that the Cu deposition potential chosen was certainly optimized for a large GMR and not for completely avoiding Co and Ni dissolution which processes may have taken place and could give rise to the formation of SPM regions as was discussed for ED Co–Cu/Cu multilayers. The dominance of the SPM term in the GMR implies also that the observed large GMR is not due to an AF exchange coupling of adjacent magnetic layers. This conclusion is further supported by the fact that the M(H) curves for these multilayers [9,10,28] were found to exhibit a fairly high relative remanence (0.7), a feature similar to the case of ED Co–Cu/Cu multilayers and, therefore, the same discussion given previously for the latter system applies also for ED Co–Ni–Cu/Cu multilayers. For the ED Co–Ni–Cu/Cu multilayers studied by Alper and coworkers [9,10,28], the MR(H) curves for larger Cu layer thickness (e.g., 2 nm) exhibited a shape very similar to those shown in Figs. 37c and d for ED Co–Cu/Cu multilayers obtained at a non-optimized Cu deposition potential: a characteristic splitting due to a larger GMRFM component but still with a non-saturating component (GMRSPM) of comparable size. This change of the MR(H) curve shape was accompanied with a strong reduction of the total GMR for ED Co–Ni–Cu/Cu multilayers as seen in Fig. 54b. This may imply that the thicker Cu layers effectively prevent of the formation of SPM regions; however, concomitantly, in this manner the major source of GMR also dies out and the total GMR strongly reduces. The above described GMR results [9,10,28,81] were obtained on ED Co–Ni–Cu/Cu multilayers prepared at low pH (around a value of 2). Alper and coworkers [28] have also studied the dependence of GMR magnitude on bath pH (in this work, a Cu layer of at least 5 nm thickness was first deposited on the substrate in order to achieve a diffusion-limited condition for Cu before the start of the magnetic layer deposition, i.e., to better ensure a constancy of the Cu-content in the subsequent magnetic layers). It was observed that there is a drastic reduction of GMR with increasing bath pH (Fig. 55).

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Fig. 55. Variation of the GMR magnitude measured at 8 kOe with the electrolyte pH for a series of ED Co–Ni–Cu(1.5 nm)/ Cu(0.8 nm) multilayers. Reprinted from Ref. [28] with permission of The Electrochemical Society.

Compositional analysis revealed that, at the same time, the Cu content in the magnetic layer increased roughly by a factor of two, a finding similar to that found on ED Ni–Cu/Cu multilayers [102]. In Section 5.7.1, we have already discussed to some extent the possible reasons of the influence of pH on GMR in the case of ED Ni–Cu/Cu multilayers that may partly apply also to ED Co–Ni–Cu/Cu multilayers. Here, we should add some more remarks on the basis of the results of Ref. [28]. Namely, the saturation magnetization was found to decrease, for both low (0.8 nm) and high (2 nm) Cu layer thicknesses, by nearly a factor of two from the lowest to the highest pH value, in accordance with the increased Cu content of the magnetic layers. For thin Cu layers (the same series as shown in Fig. 55), the MR(H) curves indicated with increasing pH an evolution of the magnetoresistance from SPM-dominated GMR to AMR at the highest pH values [the MR(H) splitting and the coercive field remained constant]. On the other hand, for the multilayer series with the thicker Cu layers, the MR(H) curves evolved from the case of comparable GMR terms due to SPM and FM contributions to fully SPM-dominated GMR, accompanied by a reduction of the coercive field and the MR(H) splitting by about a factor of 2. These findings indicate that the bath pH value has a tremendous effect on the growth mode of the individual layers and this effect can even be drastically different for various individual layer thicknesses. It was also revealed by XRD measurements that the texture does not change with pH in the range from 1.8 to 3.3 and the visibility of first-order satellite reflections remained unchanged. However, a comparison of the nominal bilayer repeat lengths with values derived from the satellite peak positions indicated a reduction of the current efficiency of the magnetic layer deposition from 75% (pH = 1.8) to 50% (pH = 3.3), i.e., the bath pH can drastically influence the growth rate as well. In case of fairly low nominal magnetic layer thicknesses (in most cases, also in Fig. 55, the nominal value was 1.5 nm), the magnetic layers can become so thin indeed where the formation of SPM regions is strongly promoted as was actually observed for most multilayers. Alper and coworkers [11,39] investigated also the GMR of ED Co–Ni–Cu/Cu multilayers prepared in P/P mode from the same sulfamate/sulfate electrolyte on n-type (1 0 0)-oriented GaAs wafers. Prior to multilayer deposition, the wafers were covered for ohmic contact with a Cu layer electrodeposited

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Fig. 56. Field-dependence of the transverse magnetoresistance of an ED [Co–Ni–Cu(5 nm)/Cu(5 nm)]100 multilayer prepared on n-type (1 0 0)-oriented GaAs wafer covered with a Cu layer. Reprinted from Ref. [11] with permission of The Electrochemical Society.

galvanostatically from a pyrophosphate bath containing Cu2+ ions. The deposition potentials chosen for the multilayer preparation were ECo–Ni–Cu(SHE) = 2.06 V and ECu(SHE) = 0.26 V. It turned out from energy-dispersive X-ray analysis by TEM that the Co:Ni ratio in the multilayers (i.e., also in the magnetic layers) was 2:3. The XRD pattern of a [Co–Ni–Cu(4 nm)/Cu(2 nm)]200 multilayer revealed a strong (1 1 1) peak and a smaller (2 0 0) peak. A small additional peak on the low-angle side of the (1 1 1) peak was attributed to a first-order superlattice reflection the position of which was used to deduce a bilayer period of 3.5 nm. Cross-sectional TEM images recorded at underfocused conditions revealed the layered structure in an ED Co–Ni–Cu(5 nm)/Cu(5 nm) multilayer. A Fourier transform of the TEM image of this multilayer yielded an average bilayer period length of 6.0 nm with a spread of 0.7 nm. The fact that both XRD and TEM provides bilayer lengths by about 40% smaller than the nominal values is in agreement with the reduced current efficiency mentioned above on the basis of other works of the same research group. As to the microstructure of these multilayers, an analysis of TEM images showed the films to be comprised of small and randomly oriented grains with an average size of 20 nm. A large transverse magnetoresistance (about 10%) was reported [11] for a [Co–Ni–Cu(5 nm)/ Cu(5 nm)]100 multilayer (Fig. 56) and nearly 7% was measured for the longitudinal component. This is clearly due to a GMR effect with high field-sensitivity (0.07%/Oe over a range of 100 Oe). The observed GMR is due to a GMRFM term as shown by the low saturation field of about 500 Oe and by the split MR(H) curves. The large difference between the LMR and TMR components is due to the AMR effect inside the relatively thick (5 nm) magnetic layers. For comparison, it is noted that in the same work [11] an AMR (LMR  TMR) of about 7% was reported for a 30-nm-thick Co–Ni–Cu film deposited under similar conditions on the same type of substrate. It may be appropriate at this point to comment on a striking difference between the MR(H) curves in this latter work [11] and the formerly discussed works [9,10,28] of the same research group. Namely, whereas the shapes of the reported MR(H) curves indicate practically the complete absence of SPM regions in the multilayers of Ref. [11], the majority of the multilayers in the other works [9,10,28] are characterized with the dominance of a GMRSPM term. This difference is due to at least three reasons. (i) In Ref. [11], the Cu deposition potential was chosen much more negative [ECu(SHE) = 0.26 V] than in Refs. [9,10,28] [ECu(SHE) = +0.09 V or +0.04 V]. Although a true optimization by any of the methods suggested in Section 3.6 has not yet been carried out for the specific sulfamate/

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sulfate bath used by Alper and coworkers in all these works (and, unfortunately, the pH value was not reported at all in Ref. [11]), the more positive ECu values of Refs. [9,10,28] are certainly more prone to promote a dissolution of the previously deposited magnetic (less noble) metals during the Cu deposition pulse. As discussed for ED Co–Cu/Cu multilayers in Section 5.2.3.2 on the basis of the works of Weihnacht and coworkers [99,115,128,132,139], such a dissolution can lead to a thinning of the magnetic layers to the extent that magnetically decoupled SPM regions can form. (ii) In the case of such a dissolution process at not sufficiently negative Cu deposition potentials, the formation of SPM regions is more probable for smaller nominal magnetic layer thicknesses as also pointed out in Section 5.2.3.2 (see Fig. 42). In Refs [9,10,28], the magnetic layer thicknesses were in the range from 0.7 nm to 2 nm, whereas in Ref. [11], this thickness was 5 nm. In the latter case, even if the dissolution is strong, it can be expected that a continuous, fully FM magnetic layer forms in the ED Co–Ni–Cu/Cu multilayers. (iii) Finally, the comparison of the results of Ref. [11] with those of Refs. [9,10,28] fully supports the conclusions obtained on ED Co–Cu/Cu multilayers [128] (also discussed in Section 5.2.3.2) in that the formation of SPM regions in ED multilayers with GMR behavior can be effectively suppressed by using sufficiently smooth substrates (such as semiconductor wafers). Hua and coworkers [12,16] studied the GMR of ED Co–Ni–Cu/Cu multilayers prepared in P/P mode from the sulfamate/sulfate electrolyte on highly textured pc-Cu(1 0 0) foils. The substrates were dissolved from the deposits and the multilayers were mounted on glass plates for further studies. The deposition potentials chosen were ECo–Ni–Cu(SHE) = 1.56 V and ECu(SHE) = 0.02 V. The magnetic layer composition was established to be about Co64Ni31Cu5. The multilayers comprised typically 400 bilayers and the nominal magnetic layer thickness was kept constant at 2.2 nm whereas the Cu layer thickness varied from 0.7 nm to 3.3 nm. XRD revealed [12] a dominant (1 0 0) texture epitaxial to the substrate, with some (1 1 1) contribution as well. First-order satellites around the (1 0 0) reflection for both dCu = 1 nm and 2.3 nm indicated the presence of well-defined periodicity. A cross-sectional TEM study by Lorenz microscopy was also performed for an ED [Co–Ni–Cu(5 nm)/ Cu(1 nm)]200 multilayer deposited on mechanically polished (1 1 1)-oriented Cu single crystal and a clear multilayered structure was revealed. Both the LMR and TMR components of the magnetoresistance were found to be negative, indicating the presence of a GMR effect. The evolution of the magnitude of the measured GMR with Cu layer thickness is shown in Fig. 57. In order to be able to discuss these results, we should first consider the MR(H) and M(H) curves reported for dCu = 1.0 nm and 2.3 nm (Fig. 58). As Fig. 58a1 shows, the magnetoresistance does not saturate up to 10 kOe for dCu = 1 nm and the shape of the MR(H) curve, which is very similar to that

Fig. 57. Dependence of magnetoresistance DR/Ro on Cu layer thickness t(Cu) for ED [Co–Ni–Cu(2.2 nm)/Cu(tCu)]400 multilayers. The triangles indicate the measured maximum GMR data for applied field up to 8 kOe, and the squares give the extrapolated saturation GMR. Reprinted from Ref. [12] with permission, copyright (1994) of American Institute of Physics.

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Fig. 58. MR(H) (upper graphs) and M(H) (lower graphs) curves for ED [Co–Ni–Cu(2.2 nm)/Cu(dCu)]400 multilayers with dCu = 1 nm (left column) and 2.2 nm (right column). Note: 1 T = 10 kOe. Reprinted from Ref. [12] with permission, copyright (1994) of American Institute of Physics.

shown in Fig. 31b or in the lower left panel of Fig. 34, indicates the dominance of a GMRSPM term. The corresponding M(H) curve (Fig. 58a2) indicates low coercivity, in agreement with the small MR hysteresis, and high remanence, the latter being an indication for the absence of an AF coupling and the presence of a FM coupling. According to Fig. 58b1, for dCu = 2.3 nm, the magnetoresistance has a component saturating in a few kilooersted magnetic field and there is an increased hysteresis as well (split MR(H) curves). All this means that the observed magnetoresistance, being much larger than was for dCu = 1 nm is dominated by a GMRFM term due to uncoupled magnetic layers. The M(H) curve (Fig. 58b2) exhibits an increased hysteresis similarly to the MR(H) curve (Fig. 58b1) and the remanence is somewhat reduced with respect to the case of dCu = 1 nm (Fig. 58a2). This latter feature may be due to the more random alignment of adjacent layer magnetizations for dCu = 2.3 nm since the layers become much less coupled here than was for dCu = 1 nm where a residual FM coupling can contribute to a larger remanence. With reference to Section 5.2.7, we can therefore see an analogy between the Cu-layer-thickness dependencies of the GMR in ED Co–Cu (Fig. 47) and Co–Ni–Cu/Cu (Fig. 57) multilayers. The overall evolution of GMR is similar in both cases and the larger GMR magnitude with increasing dCu is due to the formation of more and more perfect Cu layers, as a consequence of which the FM coupling being strong at low dCu progressively dies out and at large enough dCu, the magnetic layers become mostly uncoupled. The small GMR peak at dCu = 1 nm for the ED Co–Ni–Cu/Cu multilayers (Fig. 57) may arise from a dominance of a GMRSPM term due to the specific microstructure formed here, as was seen also in some studies of ED Co–Cu/Cu multilayers (cf. Fig. 44). The possible differences in the microstructures of ED Co–Cu/Cu and Co–Ni–Cu/Cu multilayers may well explain also the different Cu layer thicknesses where an AMR ? GMR transition and a maximum of the GMR magnitude occur. Hua and coworkers [16] carried out the measurement of the MR(H) curves also at T = 15 K for those two multilayers for which the room-temperature data were already shown in Fig. 58. For both Cu layer thicknesses, the GMR at the maximum field strength doubled at low temperature whereas the shape of the MR(H) curves remained unchanged in both cases. The only significant variation at low temperature was the increased coercivity for dCu = 2.3 nm as indicated by the larger splitting of the corresponding MR(H) curve. This is an expected result for (uncoupled) individual layers as was seen also in Fig. 49d for ED Co–Cu/Cu multilayers, supporting in this manner the picture depicted above for ED Co–Ni–Cu/ Cu multilayers. It may also be noted that the temperature dependence of the GMR of ED Co–Ni–Cu/Cu

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multilayers follows the behavior [unchanged MR(H) curves] of ED Co–Cu/Cu multilayers (Section 5.8) and not that of ED Ni–Cu/Cu multilayers (Section 5.7.2). This is in agreement with the composition ratio Co:Ni = 2:1 of the magnetic elements in the ED Co–Ni–Cu/Cu multilayers. By using the same sulfamate/sulfate bath as Alper and coworkers [9–11,28,81], Nabiyouni and coworkers produced ED Co–Ni–Cu/Cu multilayers in order to study the influence of substrate orientation [23,39,50], deposition pulse mode [39,50], heat treatment [38,54] and Co:Ni ratio [87] on GMR. In all these studies deposition was on single crystal or polycrystalline Cu substrates which were then dissolved from the deposits and the multilayers were mounted on glass plates. In the first of these works, Nabiyouni and Schwarzacher [23] investigated the GMR of ED Co–Ni–Cu/ Cu multilayers grown with ECo–Ni–Cu(SHE) = 1.36 V and ECu(SHE) = +0.04 V on pc-Cu(1 0 0) as well as (1 0 0)- and (1 1 1)-oriented Cu single crystal substrates. The ratio of the magnetic elements was Co:Ni = 3:1 and XRD revealed an epitaxial growth in that the multilayer texture/orientation was the same as that of the substrate whereas the presence of satellites was not reported. The GMR of multilayers grown on single crystal substrates was generally about two third of the largest values achieved on polycrystalline substrates. For a constant magnetic layer thickness of 3 nm, the GMR magnitude measured at 8 kOe on multilayers grown on Cu(1 0 0) single crystals had a maximum of 15% at dCu = 0.7 nm and another one with 4.5% at dCu = 1.9 nm. For multilayers on Cu(1 1 1) single crystals, a GMR maximum of 8.5% was obtained at dCu = 0.7 nm and then the GMR rapidly decreased to a constant value of 3.5% in the Cu-layer thickness range from 1.3 nm to 2.3 nm. The MR(H) curve was reported only for the Co–Ni–Cu(3 nm)/Cu(0.7 nm) multilayer grown on Cu(1 0 0) single crystal and it exhibited the same shape as that shown in Fig. 54a, indicating a dominant GMRSPM contribution. Therefore, the GMR maximum at dCu = 0.7 nm cannot be considered as arising from an AF coupling. In lack of further details on the structure as well as the MR(H) and M(H) curves for multilayers with dCu  1.9 nm, one cannot properly assess the origin of the second GMR peak. Nabiyouni and Schwarzacher [38,54] also studied the influence of annealing on the GMR of ED [Co–Ni–Cu(3 nm)/Cu(1 nm)]1000 multilayers for which the MR(H) curve in the as-deposited and annealed states was the same as shown in Fig. 54a, i.e., with dominating GMRSPM contribution. An anneal at 100 °C for 1 h in vacuum caused a small but reproducible increase of the GMR magnitude. Similarly to the work of Uhlemann et al. [98] on ED Co–Cu/Cu multilayers (see Section 5.2.3.3), this slight increase might have been caused by interface relaxation and/or decomposition. Upon further annealing up to 600 °C, the GMR in ED Co–Ni–Cu/Cu multilayers [38,54] was found to decrease continuously whereas the shape of the MR(H) curves remained unchanged. At 600 °C annealing temperature, the multilayer exhibited AMR only (LMR > 0, TMR < 0), probably due to the break-up of the multilayer structure. A further long-term (15 h) annealing at 600 °C recovered a small GMR (1%) by converting probably the sample into a granular magnetic alloy. Nabiyouni and coworkers [39,50] compared the GMR characteristics of ED Co–Ni–Cu/Cu multilayers grown on pc-Cu(1 0 0) and pc-Cu(1 1 0) substrates in both G/G and P/P mode with identical nominal layer thicknesses. An XRD study showed that the multilayers grown in P/P mode exhibit a preferred orientation identical with that of both the pc-Cu(1 0 0) and pc-Cu(1 1 0) substrates whereas multilayers prepared in G/G mode tend to grow preferentially with (1 1 1) texture on both substrates. A GMR behavior was observed for multilayers grown on pc-Cu(100) when deposited either in P/P or G/ G mode and also on pc-Cu(1 1 0) in P/P mode (with the GMR magnitude measured at 8 kOe decreasing from about 8% to 4% in this sequence) whereas the multilayer grown on pc-Cu(1 1 0) in G/G mode exhibited AMR only. It is hard to assess directly the influence of substrate texture and deposition pulse combination on GMR but a structural study by cross-sectional TEM [50] revealed that the GMR magnitude and the absence of GMR is unequivocally correlated with the microstructures developed in the multilayers under the specific deposition conditions. Namely, whereas a multilayered structure could be observed on a small scale in each multilayer investigated and a columnar growth structure was also typical, the column widths were quite different in the individual multilayer samples. Specifically, the column widths ranged from 200 nm to 2 lm for the multilayer with the best GMR [pc-Cu(1 0 0) substrate, P/P mode)] and from 20 nm to 1 lm for the two multilayers with intermediate GMR magnitudes [pc-Cu(1 1 0) substrate and P/P mode as well as pc-Cu(1 0 0) substrate and G/G mode] whereas the column widths were 10–60 nm for the AMR multilayer [pc-Cu(1 1 0) substrate and G/G mode]. With decreasing column size, the volume fraction of the disordered intergrain regions in the

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multilayer evidently increases and, thus, causes an increase of the background resistivity, leading naturally to a reduction of the GMR magnitude. For very small column (grain) sizes, the multilayer structure may become disrupted to an extent that no more GMR contribution to the observed magnetoresistance can occur and the sample will exhibit AMR only. Nabiyouni and coworkers [87] investigated the GMR of ED Co–Ni–Cu/Cu multilayers with varying Co:Ni ratio in the magnetic layers for two series with constant magnetic (3 nm) and non-magnetic (1 nm or 2 nm) layer thicknesses. The Cu-content of the magnetic layers was approximately constant at around 20 at.% whereas the Co-content was increased continuously from 0 to about a Co:Ni ratio of 4:1 (this was achieved by adding more and more Co-sulfate to the bath). A maximum GMR of about 7% was achieved for Co:Ni  3:1. The shape of the MR(H) curves indicated the presence of SPM regions for all compositions and a fitting with the Langevin function could be carried out, yielding average SPM region sizes of 3000–5000 lB. 5.8.1.2. Deposition from a bath containing Ni-sulfamate, Co-sulfate and Cu-sulfate with sulfamic acid addition. The second variant of the sulfamate/sulfate baths was obtained by adding sulfamic acid to the electrolyte by Schwarzacher and coworkers [44,56,66,69,86,122] with the purpose of shifting the pH to a value down to 2 or even below, by recognizing the beneficial effect of low pH for the GMR magnitude from earlier studies [28]. Schwarzacher and coworkers [44,56,66,69] investigated the GMR of ED Co–Ni–Cu/Cu multilayers grown on single crystal semiconductor substrates. First, multilayers with 20 bilayer repeats [44,56] were obtained on n-type (1 0 0)-oriented Si wafers having a resistivity in the range 3–7 lX cm and electrodeposition was carried out directly on the Si wafer by applying a GaIn back contact. The deposition potentials were ECo–Ni–Cu(SHE) = 2.06 V and ECu(SHE) = +0.06 V. To ensure a full coverage of and good adhesion to the Si substrate, a Co–Ni–Cu layer of 10 nm thickness was first deposited. A representative SEM image of an ED [Co–Ni–Cu(25 nm)/Cu(2 nm)]20 multilayer showed metallic appearance but at the same time also a granular structure. The GMR was studied for multilayers with magnetic layer thicknesses from 2 nm to 5 nm and Cu layer thicknesses from 0.7 nm to 5 nm. Irrespective of the applied magnetic layer thicknesses in the above range, the shape of the MR(H) curves was decisively determined by the Cu layer thicknesses. For Cu layer thicknesses from 0.7 nm to 1 nm, the MR(H) curve had the same shape as that shown in Fig. 54a, i.e., no saturation occurred up to 8 kOe (where the GMR value was about 3%) and this can be ascribed to a dominating GMRSPM contribution. For dCu = 4 nm and 5 nm, the reported MR(H) curves were split and showed saturation in magnetic fields around 1 kOe, just as in Fig. 56, with GMR magnitudes as high as 10%. This indicates a dominant GMRFM term for large Cu layer thicknesses. For a multilayer series with constant magnetic layer thickness (2 nm), the GMR showed a sharp increase as a function of the Cu layer thickness from 0.7 nm to 2 nm and there was a small further increase up to 4 nm. This behavior can be considered as being quantitatively similar to that found for ED Co–Cu/Cu multilayers (cf. Fig. 47) and the increasing GMR with increasing Cu layer thickness can be equally explained by the continuous structural improvement of the Cu layers separating the magnetic layers. The GMR was found to show a nearly monotonous decrease with increasing magnetic layer thickness in the range from 1.5 nm to 4 nm [44] and this behavior is again similar to the case of ED Co–Cu/Cu multilayers (see Figs. 36a and b). However, it should be mentioned that according to the notice of the authors of Ref. [44], there was poor reproducibility of the GMR in some cases. This could be caused eventually by the very low (20) bilayer numbers of their multilayers since at this early stage of the multilayer deposition, the growth may not have yet reached quasistatic conditions and strong fluctuations could occur. Schwarzacher and coworkers [66,69] investigated also the GMR of ED [Co–Ni–Cu(3 nm)/Cu(dCu)]50 multilayers as a function of the Cu layer thickness, which were grown directly on Si-doped GaAs wafers with either (0 0 1) or (1 1 1) crystal orientation. Electrodeposition was carried out by applying a GaIn back contact on the semiconducting wafer and the multilayers were prepared at the deposition potentials ECo–Ni–Cu(SHE) = 1.96 V and ECu(SHE) = 0.06 V. XRD revealed strong multilayer texture corresponding to the GaAs substrate orientation although with different degree in the two cases [I002:I111 = 3:1 for GaAs(0 0 1) and I111:I002 = 10:1 for GaAs(1 1 1) substrate]. Whereas clear satellites could be observed for most multilayers grown on GaAs(0 0 1), which gave bilayer repeat periods in good agreement (within about 10%) with the nominal values, satellites were not discernible on the

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Fig. 59. Dependence of the transverse magnetoresistance on Cu layer thickness for ED [Co–Ni–Cu(3 nm)/Cu(dCu)]50 multilayers grown on GaAs(0 0 1) substrates (solid square) and on GaAs(1 1 1) substrates (open circles). Reprinted from Ref. [69] with permission of The Electrochemical Society.

multilayers deposited on GaAs(1 1 1). Plain-view TEM imaging confirmed the texture degree and orientation for both kinds of multilayers. From the reported SAED patterns, one can also infer that there may be much smaller column widths in multilayers grown on GaAs(1 1 1) substrates, supporting the presence of stronger disorder already indicated by the absence of XRD satellite reflections in these samples. The MR(H) curves were very similar to that shown in Fig. 56 (splitting and saturation for magnetic fields around 1 kOe) for both kinds of substrates. There was, however, a difference in the magnitude of the GMR as shown in Fig. 59. In agreement with less perfect structure of multilayers electrodeposited on GaAs(1 1 1) substrates, the GMR magnitude also was lower, at least for not too large Cu layer thicknesses. Apparently, the quality of Cu layers (smooth growth with constant layer thicknesses and absence of pinholes leading to FM coupling) is attained already at 2 nm Cu layer thickness when growing the Co–Ni–Cu/Cu multilayers on GaAs(0 0 1) wafer whereas this degree of structural perfectness is reached at around 3 nm Cu layer thickness only for multilayers on GaAs(1 1 1) substrates. This implies that the layer growth modes and, thus, microstructural features, are strongly dependent on the deposition conditions, influencing indirectly the development of GMR characteristics. A similar behavior was seen in Fig. 47 for ED Co–Cu/Cu multilayers in that the onset of transition from AMR to GMR behavior varied from study to study, evidently due to the non-identical deposition conditions the details of which are not known to the extent as here in the case of the two sets of ED Co–Ni–Cu/Cu multilayers, i.e., the different crystal orientations of the two substrates. Schwarzacher and coworkers [86] investigated the GMR of ED [Co–Ni–Cu(4 nm)/Cu(dCu)]30 multilayers grown in P/P mode in a channel flow cell on glass plates covered with sputtered Cr(3 nm) and Au(25 nm) layers. The same electrolyte was used as in the previously discussed work and the deposition potentials applied (ECo–Ni–Cu = 1.8 V and ECu = 0.5 V) were specified with respect to a Pt wire reference electrode immersed in the bath. Unfortunately, the uncertainty of the actual cathode potentials when referenced to such an electrode can be as high as 0.5 V, one cannot reliably convert the specified potential to values referencing to the SHE scale. No structural characterization of the multilayers was carried out. Chemical analysis of [Co–Ni–Cu(4 nm)/Cu(6 nm)]20 multilayers by energy-dispersive X-ray spectroscopy in SEM revealed that at the maximum flow rate applied (60 cm/s) the Co:Ni ratio was typically 2:1 whereas the overall Cu content was relatively constant around 80 at.% for a wide range of Cu2+ ionic content of the electrolyte (0.004–0.016 M). A stagnant bath with much higher Cu2+ concentration (0.055 M) yielded multilayers with the same thicknesses which had a Co:Ni

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ratio of 3:1 and the same Cu content as the multilayers prepared in flow electrolytes. As to the magnetoresistance, for a moderate flow rate (13 cm/s), the transverse GMR exhibited a maximum-like behavior as a function of dCu, just as observed for ED Co–Cu/multilayers (see Fig. 47). The GMR reached a maximum of about 7% for 5–6 nm Cu layer thicknesses where the MR(H) curves were similar to that shown in Fig. 56, with a saturation field below 1 kOe and with splitting. At the same flow rate, another identical multilayer series was prepared in G/G mode and the TMR magnitude exhibited the same evolution with Cu layer thickness except that the maximum GMR was reduced to about a half value at the peak and the peak itself was also much narrower than for multilayers deposited in P/P mode. The GMR magnitude was also studied for a series prepared with varying flow rates: it was found that the GMR decreases from 6% at zero flow (stagnant bath) to about 1% at around 60 cm/s flow velocity. However, this strong reduction of GMR was observed for a relatively high Cu content in the bath since for an electrolyte with much lower Cu content, the MR(H) curve shape and the GMR magnitude was the same at 60 cm/s flow rate as obtained under stagnant conditions from the bath with higher Cu content. The authors discussed all these effects on the basis of the influence of flow rate on limiting Cu deposition current at various Cu-contents in the bath. It was also assumed that the electrolyte flow might create rougher interfaces; however, in lack of any structural characterization of the multilayers investigated, one cannot assess the validity of this argumentation. Another issue is that the flow conditions may affect the potential at the Pt wire reference electrode and, thus, the actual electrode potential of the cathode may have been different for the various flow rates. Nabiyouni and Schwarzacher [122] also reported on the GMR of ED Co–Ni–Cu/Cu multilayers prepared in P/P mode from the same stagnant bath on Cu(1 0 0) and Cu(1 1 1) single crystals as well as pcCu(1 0 0) substrates. High-angle XRD showed that multilayers on single crystals grow with a texture corresponding to the substrate orientation but no satellites were observed. Low-angle XRD investigation at a synchrotron source, on the other hand, revealed superlattice peaks up to the second order on 100 repeats of a Co–Ni–Cu/Cu multilayer grown on Cu(1 0 0) single crystal. The MR results on these multilayers resemble very much the behavior reported by the same authors previously [23] for similar multilayers prepared under identical conditions from the sulfamate/sulfate bath without sulfamic acid addition which results were already discussed in Section 5.8.1.1. 5.8.1.3. Deposition from baths containing Ni-sulfamate, Co-sulfamate and Cu-sulfate with various additives. The third variant of the sulfamate/sulfate baths was used by Kainuma et al. [43] who investigated the GMR of ED Co–Ni–Cu/Cu multilayers grown in P/P mode (without specifying the electrode potentials applied) onto vacuum-deposited Cu films from electrolytes containing Ni-sulfamate, Cosulfamate and Cu-sulfate with various additives. Two kinds of additives were applied: (a) H3BO3 and Triton X-100 and (b) Na-citrate and NaCl. Solutions (a) and (b) differed also in the sense that the ionic ratio of the magnetic elements to the non-magnetic one was nearly three times higher in solution (a) than in solution (b). No MR(H) curves were reported, only the dependence of transverse magnetoresistance measured at H = 15 kOe was displayed as a function of the Cu layer thickness for multilayers deposited from both electrolytes with 50 bilayers but without specifying the magnetic layer thickness. For bath (a) where the Cu content in the magnetic layers was lower, the TMR values were around 6% between 5 and 10 nm Cu layer thicknesses and dropped to about 2% outside this range. For bath (b) where the Cu content in the magnetic layers was presumably significantly larger, even larger TMR values (13%) were obtained around dCu = 5 nm but they decreased rapidly for both thicker and thinner Cu layers down to the level of about 2–3%. 5.8.2. Deposition from a sulfamate bath in G/0/G mode Cavallotti et al. [48] investigated the GMR of ED Co–Ni–Cu/Cu multilayers grown from an electrolyte containing Ni-sulfamate, Co-sulfamate, Cu-sulfamate and NaKC4H4O6. Electrodeposition was carried out from baths with pH of 6.4 under gentle stirring at 48 °C onto Si wafers coated with various sputtered seed layers such as (A) Ni80Fe20(100 nm), (B) SiO2(100 nm) + Ni80Fe20(100 nm), (C) Cr(20 nm) + ED Au flash, (D) ITO(160 nm) + ED Au flash and (E) SiO2 + Cr(70 nm) + ED Au flash. A G/ 0/G deposition pulse sequence was applied where a zero-current pulse (0.1 s long) was inserted after the magnetic layer deposition pulse of 5–6 mA/cm2 amplitude and 1–3 s length, whereas Cu was deposited during a long pulse (4–60 s) with an amplitude of 0.5 mA/cm2. The magnetic layer was

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determined to have a Co:Ni ratio of 20:1 and a Cu content of about 30 at.%. XRD revealed an fcc-(1 1 1) preferred orientation in all multilayers although the degree of texture as measured by the intensity ratio I111:I200 varied depending on bath composition, substrate and Cu layer thickness. Satellite reflections were also obtained in some cases. Special care was taken to examine the substrate roughness by SEM and, in general, it was found that multilayers deposited on smooth substrates with shiny appearance yielded higher GMR magnitude, higher field sensitivity and lower coercivity than those on rough substrates. The longitudinal magnetoresistance values were measured and since no corrections were made for the relatively thick substrate layers in some cases, the maximum GMR values did not exceed about 7%. The MR(H) curves reported for [Co–Ni–Cu(3 nm)/Cu(5 nm)]40 multilayers deposited on various substrates and from various bath compositions all exhibited splitting and saturated in magnetic fields at or below 1 kOe. This means that in such multilayers the observed magnetoresistance is dominated by a GMRFM term. Below about 3 nm Cu layer thickness, the MR(H) curves were found to become gradually broadening and this is indicative of an increasing saturation field as can directly be seen in the MR(H) curve of a [Co–Ni–Cu(3 nm)/Cu(1.2 nm)]255 multilayer for which the magnetoresistance did not saturate up to about H = 3 kOe. This is an indication that for low Cu layer thicknesses an SPM type GMR contribution dominates the observed magnetoresistance. As to the dependence of GMR on Cu layer thickness, it exhibits a similar behavior as already shown for ED Co–Cu/Cu (Fig. 47) and Co–Ni–Cu/Cu (Fig. 59) multilayers: the GMR starts to increase monotonously from dCu  2 nm, it reaches a maximum in the range from 5 nm to 8 nm and decreases only slightly up to about dCu = 11 nm. 5.8.3. Deposition from a sulfate/citrate bath in P/P mode Dulal and coworkers [105,142] investigated the GMR of ED Co–Ni–Cu/Cu multilayers prepared in P/ P mode from an electrolyte containing Co-sulfate, Ni-sulfate, Cu-sulfate and Na-citrate. The same type of a flow channel cell with flow rates ranging from 10 to 30 cm3/s was used as discussed in Section 5.8.1.2 [86]. The deposition potentials were ECo–Ni–Cu = 2 V and ECu = 0.6 V, both potentials measured with respect to the anode which was a Cu plate facing the cathode and having an area 20 times larger than that of the cathode. Deposition was carried out on a quartz disc coated with Ti and Au layers. The multilayer growth adopted the (1 1 1) texture of the Au seed layer. Most of the multilayers exhibited visible first-order satellites around both the (1 1 1) and the (2 2 2) reflections. The MR(H) curves of a [Co–Ni–Cu(2 nm)/Cu(3 nm)]200 multilayer prepared at a pH of 4.5 [105] were similar for both the LMR and TMR components to that shown in Fig. 54a and did not seem to show saturation up to 8 kOe where the values were 1.8% and 2.1%, respectively. This indicates a dominance of an GMRSPM type contribution to the measured magnetoresistance. The LMR(H) and TMR(H) curves reported in Ref. [142] for a [Co–Ni–Cu(2 nm)/Cu(2 nm)]100 multilayer deposited at a pH of 6 were very similar with smaller GMR values at 8 kOe (1.2%). In this latter work, the influence of Co2+ and Ni2+ concentration in the bath, the Ni/(Ni + Co) molar ratio in the multilayer, the nominal layer thicknesses, the number of bilayers, the flow rate and the electrolyte pH on the magnetoresistance were investigated in great details. Most listed parameters were found to affect the GMR magnitude to some extent. Although the negative sign of both LMR and TMR components indicated that GMR was exhibited by all multilayers studied, the largest GMR obtained at 8 kOe was 2.5% only and this was always dominated by the GMRSPM contribution. The low GMR might be thought caused by the relatively high pH as was also for the case of ED Ni–Cu/Cu multilayers prepared also from sulfate/citrate bath in the same flow channel cell at similarly high pH (see end of Section 5.7.2). Surprisingly, it was found, however, in Ref. [142] that at low pH values (1.5–3.0) this specific Co–Ni–Cu sulfate/citrate electrolyte yielded multilayers with practically zero magnetoresistance (not even AMR was observed). A chemical analysis revealed that the amount of Ni and Co in these latter multilayers reduced to the level of 1% which naturally explains the absence of any magnetoresistance present. This finding points out the fact that the choices of supporting electrolyte, bath pH and hydrodynamic conditions are strongly interrelated in forming the microstructure of multilayers and, thus, their magnetotransport properties. 5.8.4. Deposition from a sulfate bath with additives in G/G mode Zhang et al. [123,124] investigated the GMR of ED Co–Ni–Cu/Cu multilayers prepared in G/G mode from an electrolyte containing Co-sulfate, Ni-sulfate, Cu-sulfate, H3BO3, saccharin and triton X-100 at

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a pH of 5.2. The multilayer preparation and investigation was identical with conditions described in Section 5.2.4.4 for ED Co–Cu/Cu multilayers and here we intend to provide only some results obtained by these authors upon the addition of Ni to the deposition bath. A small Ni-content in the multilayer (the Co:Ni ratio was 96.5:1 in wt.%) did not influence the magnetoresistance whereas for a Co:Ni ratio of 90.5:3.5, the TMR at 90 kOe reached a value of about 10%. However, saturation was not achieved even in this high field, so the observed GMR was dominated by a contribution of SPM regions. 5.8.5. Deposition from a sulfate bath in G/P mode The present authors [138] investigated the GMR of ED Co–Ni–Cu/Cu multilayers prepared with a typical total thickness of 1.5 lm in G/P mode from an electrolyte containing Co-sulfate, Ni-sulfate, Cu-sulfate, Na2SO4, H3BO3, HSO3NH2 and NaOH (the latter for adjusting the pH to 3.25). This bath was elaborated in order to make it economic by avoiding the use of the expensive sulfamate salts and still retaining the benefits of the advantages of the sulfamate-containing electrolytes by the addition of sulfamic acid. A tubular type cell with upward facing cathode at the bottom (see Fig. 17c) was used as in Ref. [99], the advantages of which, in avoiding edge effects and ensuring lateral homogeneity, were discussed in Section 3.4. A mechanically polished Ti sheet was used as cathode. The magnetic layer was deposited at current densities of 15 mA/cm2 or 45 mA/cm2. The Cu deposition potential was optimized according to the method described in Ref. [231] (see Section 3.6) and the optimum value was found to be ECu(SHE) = 0.345 V. Under these conditions, in the absence of Co and Ni dissolution, the actual layer thicknesses remain equal to the nominal values determined by assuming 100% current efficiency for Cu deposition and using 96% current efficiency for the magnetic layer deposition determined from measured thicknesses of d.c. plated deposits. After preparation, the multilayers were peeled off from the Ti substrates for further studies. An fcc-(1 1 1) preferred orientation was established by XRD and satellite reflections were observed in most cases, often even around the small intensity (2 0 0) reflections. From structural point of view, the multilayers deposited with higher current density for the magnetic layers were definitely of higher quality (it should be noted, however, that higher current density yielded also lower Cu content in magnetic layers). The bilayer lengths derived from the satellite peak positions were systematically higher by 10–30% than the nominal values. Both the LMR and TMR component of the magnetoresistance were measured up to 8 kOe for multilayers having increased Cu layer thickness from 1 nm to 3 nm and a dominating GMR was observed for all multilayers. The GMR magnitude was larger for higher current density and thicker Cu layers, with the largest GMR values being around 10%. A decomposition of the MR(H) curves revealed that there was a large GMRSPM contribution (non-saturating behavior and no splitting) for low current densities (higher Cu-contents) and low Cu layer thicknesses. On the other hand, for lower Cu contents and thicker Cu layers, the observed magnetoresistance was dominated by a GMRFM term (splitting and saturation fields around 1 kOe). The GMRFM contribution exhibited a nearly linear increase with increasing Cu layer thickness, a result very similar to the observation for ED Co–Cu/Cu multilayers (see Fig. 47). 5.9. Co–Ni–Ag/Ag multilayer films Fedosyuk et al. [67] reported the electrodeposition of Co–Ni–Ag/Ag multilayers in the G/G mode from a bath containing CoSO4, NiSO4, KI and AgNO3 on amorphous Ni-P substrates. Similarly to the ED Co–Ag/Ag multilayer films studied in the same paper (see Section 5.3), the magnetic layer contained some 10–15 at.% Ag and XRD revealed the presence of an fcc multilayer structure. The authors provided, however, very few results on these ED Co–Ni–Ag/Ag multilayer films. It was only mentioned that the addition of Ni to the bath had a beneficial effect on the two-dimensional character of layer growth and this improved the GMR effect over the ED Co–Ag/Ag multilayers but no MR data were presented for the Ni-containing multilayer films. 5.10. Fe–Co–Cu/Cu multilayer films Kakuno et al. [53] investigated the GMR of ED Fe–Co–Cu/Cu multilayers prepared with typically 20 bilayers in P/P mode from a solution containing CoSO4, FeSO4(NH2)SO4, CuSO4 and H3BO3 on Si(1 1 1) wafers covered with an evaporated Cu(1 1 1) layer. The deposition potentials used were

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Fig. 60. Dependence of the transverse magnetoresistance of ED Co–Fe–Cu/Cu multilayers on Cu layer thickness tCu. Reprinted from Ref. [53] with permission of the Institute of Physics.

EFe–Co–Cu(SHE) = 1.00 V and ECu(SHE) = 0.15 V. No justification for the choice of these potentials was given, it was only noticed that at these electrode potentials the Fe–Co–Cu electrodeposits exhibited a very shiny surface whereas the Cu electrodeposits a slightly opaque surface. It was established from energy dispersive X-ray spectroscopy analysis of the multilayers that the Fe:Co ratio was 17:83 in the magnetic layers which also contained Cu up to 3 at.%. On the basis of an XRD study, it was reported that the multilayers exhibited an fcc(1 1 1) texture without evidence of satellite reflections from multilayer periodicity. The (1 1 1) diffraction peaks were very broad, indicating small grain sizes. TEM SAED patterns revealed a polycrystalline fine-grained character with very low texture; in addition, the observed diffraction rings indicated the presence of fcc-Cu, fcc-Co and hcp-Co phases. As to the magnetoresistance, the TMR component was only measured. Due to the magnetic softness of the Co–Fe alloys, the MR(H) curves of the ED Co–Fe–Cu/Cu multilayers saturated in magnetic fields around 0.5 kOe for sufficiently thick (2–4 nm) magnetic layers. However, for magnetic layer thicknesses at and below 1 nm, the MR(H) curve shape indicated an SPM-type behavior as was also the case for ED Co–Cu/Cu and Ni–Cu/Cu multilayers. The dependence of the magnetoresistance of ED Co–Fe–Cu/Cu multilayers on the Cu layer thickness is shown in Fig. 60. The high relative magnetic remanence (see inset) indicates a predominantly FM coupling between adjacent magnetic layers, even for Cu layer thicknesses as high as 4 nm where the magnetoresistance is the highest. This raises the question whether the observed peaks in the Cu layer thickness dependence of the GMR for these multilayers can have any significance. At low Cu layer thicknesses (below about 3 nm), the measured magnetoresistance is anyway so small (0.5–1.5%) that it may arise even from an AMR effect [165]. Due to the low structural quality of these multilayers, the most that can be established from Fig. 60 is that the magnetoresistance, on the average, increases with Cu layer thickness, just as was found for ED Co–Cu/Cu multilayers (see Fig. 47).

5.11. Fe–Ni–Cu/Cu multilayer films Attenborough et al. [14] investigated the GMR of ED Fe–Ni–Cu/Cu multilayers prepared onto textured pc-Cu(1 0 0) and single-crystalline Cu(1 0 0) substrates in P/P mode from a sulfate-based electrolyte containing NiSO4, FeSO4, CuSO4, sodium lauryl sulfate and sodium saccharin (the bath composition was taken from Romankiw and coworkers, see, e.g., Ref. [282]). The deposition potentials used were EFe–Ni–Cu(SHE) = 2.26 V and ECu(SHE) = 0.16 V. No justification for the choice of these potentials was given. Electron probe microanalysis measurements on a Ni–Fe–Cu(2 nm)/Cu(2.5 nm) multilayer yielded a Fe:Ni ratio of 4:6 for the magnetic layer. From the deposition current densities applied, the Cu content in the magnetic layer was estimated to be about 9 at.%. By using low-angle

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Fig. 61. Transverse (a) and longitudinal (b) magnetoresistance curves of an ED [Fe-Ni-Cu(2 nm)/Cu(2.5 nm)]200 multilayer. Reprinted from Ref. [14] with permission of Elsevier.

X-ray reflectivity under conditions of anomalous scattering for enhanced contrast between the two kinds of layers, a bilayer repeat of 3.0 nm was obtained for a [Ni–Fe–Cu(4 nm)/Cu(2 nm)]200 multilayer which was taken as an indication for 50% current efficiency during deposition. This low current efficiency was ascribed to hydrogen ions being reduced at the substrate in addition to the metal ions. The authors noted that XRD studies of other multilayers showed that the bilayer period could still be controlled reproducibly. The longitudinal and transverse MR(H) curves for a Fe–Ni–Cu multilayer are shown in Fig. 61. The identical sign of the LMR and TMR components indicates a GMR effect of the order of 1%. Due to the magnetic softness of the magnetic layer consisting predominantly of Fe and Ni, there is an extremely sharp MR component saturating in small magnetic fields (well below 1 kOe) which can arise from a GMRFM term. In addition, one can clearly observe also a component not saturating up to 8 kOe which can certainly be ascribed to a GMRSPM term. Chassaing et al. [19,20] investigated the GMR of ED Fe–Ni–Cu/Cu multilayers prepared in P/P mode from a solution containing FeSO4, NiSO4, CuSO4, H3BO3 and C6H8O6 on glass plates covered with a sputtered Au(18 nm) layer. The ionic concentration ratios were Fe2+:Ni2+:Cu2+ = 60:810:7. The electrode potentials for multilayer deposition were chosen on the basis of cyclic voltammetry data and the values used were EFe–Ni–Cu(SHE) = 1.46 V and ECu(SHE) = 0.36 V. Under these conditions, the Fe:Ni ratio in the magnetic layer was about 20:80 and this layer contained also about 5 wt.% Cu. The grazing incidence XRD patterns of the multilayers showed the characteristic lines of an fcc structure with a (1 1 1) texture although there was significant intensity also at the (2 0 0) reflection. The experimental lattice parameter of an ED [Fe–Ni–Cu(1 nm)/Cu(1 nm)]20 multilayer was 0.360(5) nm which was intermediate between the lattice parameters of pure Cu and a Fe20Ni80 alloy. The MR(H) curves measured for a [Fe–Ni–Cu(3 nm)/Cu(1.5 nm)]30 multilayer at 77 K [19] were nearly linear for the LMR and the TMR components, both being negative, and did not show a sign of saturation up to the maximum magnetic field applied (2 kOe) where the GMR was about 1%. Chassaing et al. [20] reported also the LMR(H) and TMR(H) curves for a [Fe–Ni–Cu(1.3 nm)/Cu(1.2 nm]20 multilayer at T = 4.2 K. Both components were again negative and around 3% for magnetic fields around 10 kOe where, however, still no saturation of the magnetoresistance was observed. Due to the non-saturating character of the MR(H) curves, it should be concluded that the [Fe–Ni–Cu(3 nm)/Cu(1.5 nm]30 multilayers studied by Chassaing et al. [19,20] exhibit a GMRSPM type behavior.

5.12. Co–Zn–Cu/Cu multilayer films The present authors [111] prepared ED Co–Zn–Cu/Cu multilayers with several hundred bilayers on mechanically polished Ti sheets both in the G/G and the G/P modes from a bath containing CoSO4, CuSO4 and ZnSO4. Besides studying the electrochemical effect of the presence of an element with anomalous codeposition, the motivation of the incorporation of large Zn atoms into the magnetic layer was to increase its lattice parameter, thus reducing lattice mismatch with the Cu layer. It was expected

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that this may result in reduced interfacial stresses and less structural defects in comparison with the Co–Cu/Cu multilayers. The structure of the Zn-containing multilayers was fcc with increased lattice parameters and the presence of Zn increased the degree of (1 1 1) texture. It could be deduced from a careful compositional analysis of the multilayers, by taking into account electrochemical considerations as well, that a Zn-rich layer forms at the Cu/magnetic layer interface. By measuring the LMR and TMR components of the magnetoresistance, a GMR of about 3% was obtained at 8 kOe for multilayers in which the magnetic layer contained about 6 at.% Zn. However, the GMR of the Co–Zn–Cu/Cu multilayers was only about half of the GMR value obtained for similar samples free of Zn. The decrease in GMR can be attributed to an increase of the resistivity of the magnetic layer due to the codeposition of a third element (i.e., Zn), a change in the electronic structure of the deposit and also to the presence of a Zn-rich layer at the interfaces. This Zn-rich zone produced by the anomalous codeposition results in an interface with high resistivity, thus lowering the mean free path of the electrons and leading to the loss of spin memory while the electrons are transferred from one magnetic layer to the adjacent one. 5.13. Fe–Co–Ni–Cu/Cu multilayer films Since the best soft magnetic properties (low coercivity and high induction) among the alloys of the iron-group metals with each other are exhibited by some ternary Fe–Co–Ni alloys, there were attempts to produce also ED multilayers containing all three FM metals and to investigate their GMR behavior. Huang and coworkers [85,96] studied the GMR of ED Fe–Co–Ni–Cu/Cu multilayers prepared with a bilayer number of 730 in G/G mode from an electrolyte (pH of 2.5) containing FeSO4, CoSO4, NiSO4, CuSO4, NaKC4H4O6, sulfamic acid, Na saccharin and Triton X-100 onto a rotating disk electrode (Aucovered stainless steel from which the deposits were mechanically peeled off). A two-compartment cell was used to avoid a mixing of the anodic Fe3+ and Co3+ products in the cathodic region and, in order to suppress the oxidation of the Fe2+ and Co2+ ions, nitrogen was continuously sparged into the solution. Current densities of 35.4 mA/cm2 and 0.354 mA/cm2 were used to deposit the magnetic layers (composition: Fe16Co73Ni5Cu6 obtained from analysis of d.c. plated deposits) and the Cu layers, respectively. Although a TEM study of multilayers with individual layer thicknesses of several tens of nanometers indicated a layered structure, it also revealed a very fine-grained microstructure, the grain size being of the order of some 10 nm. It turned out from XRD studies of multilayers with nanometric layer thicknesses that the majority of the grains have an fcc-(1 1 1) or a bcc-(1 1 0) orientation, with the presence of additional fcc-(2 0 0) grains. Both the relative intensity of the two major reflections and the grain size (20–80 nm) deduced from XRD line broadening varied with Cu layer thickness for a multilayer series with fixed 2 nm magnetic layer thickness. No satellite reflections due to multilayer periodicity were observed. The room-temperature MR(H) curves in most cases had a rapidly varying component apparently saturating around 10 kOe whereas in several cases usually a slowly saturating component also survived up to magnetic fields around 90 kOe. The GMR magnitude was typically as high as 5% in the maximum magnetic field. At 4 K, the central component of the magnetoresistance strongly increased, reaching about 15% at around 10 kOe whereas there remained also a non-saturating component up 90 kOe. The overall behavior of the MR(H) curves indicates a strong SPM type contribution to the GMR and this speaks for a very disordered or even granular multilayer structure as suggested already by the results of structural studies (small grain size, lack of superlattice reflections, presence of both fcc and bcc phases). Based on the experience with ED Co–Cu/Cu multilayers (see Section 5.2), this may probably be ascribed to a large extent to the presence of deleterious additives in the bath used. Gong et al. [118,119] studied the GMR of ED [Fe–Co–Ni–Cu/Cu]50 multilayers prepared in P/P mode from an electrolyte (pH of 2.8) containing FeSO4, CoSO4, NiSO4, CuSO4 and H3BO3. The electrolyte composition was previously optimized, without CuSO4, to yield low-coercivity, high-moment Fe–Co–Ni alloys [283]. The solutions were freshly prepared and purged with N2 for 1 h before use and deposition was carried out under quiescent conditions. As substrates, n-type Si(0 0 1) wafers with an active area of 1 cm2 were used which were contacted from the back by using a liquid GaIn eutectic alloy. The cathode potential for the FM layer deposition was chosen as EFe–Co–Ni–Cu(SHE) = 1.66 V and the

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Fig. 62. Magnetoresistance curves (a) and normalized hysteresis loops (b) of ED [Fe-Co-Ni-Cu(4.5 nm)/Cu(5.5 nm)]50 multilayers with various Fe contents in the magnetic layer: (1) Co86Ni14; (2) Fe9Co80Ni11; (3) Fe18Co70Ni12; (4) Fe51Co34Ni15. Reprinted from Ref. [119] with permission, copyright (2005) of the IEEE Magnetics Society.

optimum Cu deposition potential was established by the method of analyzing the current transients during the Cu deposition cycle [231]. The optimum Cu deposition potential was found to be ECu(SHE) = 0.34 V to 0.38 V, a range surprisingly close to the optimum established for ED Co–Cu/ Cu multilayers prepared from an additive-free sulfate bath (see Section 5.2.3.1). Gong et al. [118,119] established that, indeed, when approaching the optimum Cu deposition potential, the MR(H) curves indicated that the GMRSPM term was reduced (the GMRFM term became dominating and the saturation field reduced). It was also shown by an AFM study that around the optimum ECu value, the surface roughness decreased as a result of the diminished selective dissolution. The reduction of the GMRSPM term for low surface/interface roughness is in line with the model suggested by Ishiji and Hashizume [200] as one possible origin for the appearance of SPM regions in magnetic/ non-magnetic multilayers. Gong et al. [118,119] achieved a room-temperature GMR of 8% in magnetic fields as low as a few hundred oersteds, by establishing, through the control of the Fe2+ concentration in the bath, the Fe-content in the magnetic layer which yielded the largest GMR (Fig. 62a) and smallest coercivity (Fig. 62b). Nevertheless, the lack of a significant AF exchange coupling between adjacent layers was also concluded from the splitting of the MR(H) curves (Fig. 62a) and the large relative magnetic remanence (typically 0.7, see Fig. 62b). This finding agrees well with the analysis given for ED Co–Cu/Cu multilayers in Section 5.2.7, together with the similar overall evolution of the GMR with layer thicknesses in both multilayer systems. Gong et al. [118,119] also tried to estimate, from an analysis of the MR(H) behavior and GMR magnitude, the minimum thickness of the magnetic and the nonmagnetic layers in the ED Fe–Co–Ni–Cu/Cu multilayers which provide a continuous coverage of the previously deposited layer. These minimum layer thicknesses were found to be 1.5–2.2 nm for FM layer growth on Cu and >3.5 nm for the Cu layer growth on the FM layer. The difference of these values was interpreted as implying a different nucleation behavior for the two kinds of layers, a feature also observed for evaporated Co/Cu [284] and ED Co–Cu/Cu [109] multilayers. 6. GMR of ED spin-valve type structures and multilayers with view on possible applications Since the late 1970s, thanks to the devoted and pioneering activity of Lubomyr Romankiw (IBM), electrodeposition has become an important part of modern electronic technology. This is well witnessed by the historical overview of the development of electrochemical processes for magnetic thin film heads as given by Romankiw [285] and by the papers presented at the Session ‘‘Electrochemical Thin Film Technology in Honor of Dr. L.T. Romankiw” of the 8th International Symposium on Magnetic Materials, Processes and Devices in 2004 [286]. Although current applications of the GMR effect exclusively use layered structures produced by PD methods, there have already been several reports on GMR studies of ED spin-valve type structures [46,47,58,64,73,76,79,91,140,143] and multilayer films prepared with view on possible applications [63,92,101,135]. In case of further progress in the ED technology, there are prospects for the application of ED GMR structures as well.

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For introducing a GMR device as a sensor into high-tech applications, the feasibility of depositing the multilayer structure on semiconductor wafers is usually an important issue, simply due to the need of easy integration into the well-established wafer technology lines. In addition, the low level of roughness of semiconductor wafers is itself a factor being beneficial for the formation of high-quality layers and multilayers, as could be seen also above when comparing the structure (Section 4) and GMR behavior (Section 5) of ED multilayer films prepared on rough and smooth substrates. The preparation of sandwich structures such as an exchange-biased GMR spin-valve [287] which came finally in real application in 1997 as the read-head of magnetic hard disk drives [6] could be successfully synthesized on substrates only which have a roughness comparable to the semiconductor wafers. Whereas layer deposition by PD methods can be carried out on any kind of substrates with appropriate buffer layers, the situation is different for electrodeposition since for this process a substrate with sufficient conductivity is required. Therefore, in this section we first discuss the possibilities of preparing high-quality multilayer films by ED methods on semiconductor wafer substrates. Then, the principles on the basis of which GMR multilayers and spin-valves can be used for magnetic field detection will be summarized. This will be followed by a summary of works with application-oriented layered GMR structures produced by electrodeposition. 6.1. Electrodeposition of multilayer films on semiconductor wafers Semiconductor wafers to be used as substrates for electrodeposition should be either doped to achieve sufficient conductivity or metalized (covered with a thin metal film of good adherence). In the case of a doped semiconductor wafer, a metallic contact is attached to the back (unpolished side) of the wafer (either by metalizing it with a thin film via physical or electroless deposition or by placing the wafer on a metal plate and providing electrical contact between them with the help of a low-melting-point Ga alloy). By applying now an electrical current through the sufficiently conducting wafer itself, electrodeposition onto the polished wafer surface can then be carried out (this could be termed as through-plating on wafer). Before electrodeposition, the cathode surface needs to be thoroughly cleaned from oxides by etching with a solution of HF. A proper metallic buffer layer can also be deposited onto the etched wafer surface by a PD process or by electrodeposition before carrying out the electrodeposition of the multilayer film itself via through-plating. Recognizing the importance of the feasibility of electrodeposition on semiconducting wafers, it was demonstrated as early as 1994 [11,282] that high-quality ED multilayer films can be prepared by through-plating on doped semiconductor wafers. Schwarzacher and coworkers [11] obtained 10% GMR (see Fig. 56) on ED Co–Ni–Cu/Cu multilayer films prepared on n-type GaAs(1 0 0) wafers with an ED Cu buffer layer. Chang and Romankiw [282] prepared ED Fe–Co–Cu/Cu and Fe–Ni–Cu/Cu multilayers with an ED Fe-Co or Fe–Ni buffer layer on an n-type Si wafer. It was pointed out by the latter authors [282] that a Schottky barrier forms between the first metal layer and the n-type wafer material. This implies that whereas the relatively high electrical conductivity of the wafer permits electrodeposition in the through-plating mode, the polarization behavior of an uncoated semiconductor wafer can be very different from that of a metal substrate, and significantly larger overpotentials may be necessary to achieve metal deposition. From the view-point of the MR measurement, the Schottky barrier formed suppresses the shunting effect of the wafer. This was an important point since it demonstrated that GMR multilayer films for applications as MR sensors can be electrodeposited by through-plating on n-type semiconductors (for a p-type semiconductor material, this does not work in lack of the formation of a Schottky barrier). In addition, it is also a common practice to metalize a Si wafer with the help of deposited metal layer(s) thereon, providing in this manner a conducting cathode surface towards the electrolyte. For this purpose, usually a Si wafer of (1 0 0), (1 1 0) or (1 1 1) orientation is metalized by physical deposition methods (evaporation or sputtering). This implies the deposition of a first adhesive layer on the wafer (typically Cr or Ta which ensure good adherence to the semiconductor surface) and a thickness of 5–10 nm is usually sufficient. The next step is the deposition of a seed layer, mostly deposited by the same method as the adhesive layer. For a seed layer, usually Cu or Au is chosen which are sufficiently noble not to be attacked seriously when bringing the cathode in contact with the electrolyte.

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A thickness of about 20 nm is usually sufficient for the seed layer in order to avoid a strong shunting effect for the magnetoresistance measurements. When the seed layer is too thick, layers may become very rough, hence deteriorating GMR properties, as it was observed for sputtered Co/Cu multilayers [200]. A PD adhesive/seed layer pair can provide, under proper preparation conditions, a surface roughness as low as 1 nm, i.e., very close to the original roughness of the Si wafer. Several reports of the GMR on ED multilayer films and sandwiches prepared onto semiconductor substrates have been published and many of them were already explicitly mentioned when discussing the GMR of specific ED multilayer systems in Section 5. Works describing the GMR studies on multilayers or sandwich structures electrodeposited on semiconductor wafers can be categorized into three main groups. (i) Through-plating directly on semiconductor wafer (no buffer layer). Various semiconductor wafer materials were used such as n-type Si(1 0 0) with a resistivity q of 1–10 X cm [44,56,76,79] or 0.006–0.020 X cm [118,119], n-Si(1 1 1) with q = 0.001–0.009 X cm [127], n-GaAs(1 0 0) [39], n-GaAs without orientation specification [66], n-GaAs(1 0 0) which was Si-doped to a carrier concentration of 108/cm3 [46,47,64,73] and n-GaAs(1 1 1) which was Ga-terminated and Sidoped to a carrier concentration of 1–5 108/cm3 [69]. (ii) Through-plating on semiconductor wafer with buffer layer. The semiconductor wafer/buffer layer combinations used were n-GaAs(1 0 0)/Cu with the Cu layer obtained also by electrodeposition from a pyrophosphate bath [11], n-Si(1 0 0)/Cu with an evaporated Cu layer of 20–40 nm thickness [80,88] and n-Si(1 0 0)/Co(26 nm)/Cu(6 nm) where the Si wafer had a resistivity of 2– 4 X cm and the Co and Cu seed layers were obtained by electrodeposition from the same bath used in the next step to deposit the multilayer [144]. (iii) Plating on metalized semiconductor wafer. Studies of GMR have been reported on ED multilayer films prepared on Si/adhesive-layer/seed-layer substrates in numerous works [17,45,48,53, 55,57,58,63,89,93,95,97,98,100,101,107,113,114,127–129,131,135,136,139–141,143,147–149].

6.2. Application of the GMR effect in magnetoresistive sensors The phenomenon of magnetoresistance can be utilized for the measurement of magnetic field strength or, in a simpler case, for the establishment of the presence or absence of a magnetic field larger than a threshold value. These magnetoresistive sensors made of a soft magnetic alloy such as, e.g., Ni80Fe20 (Permalloy), have already been used in various applications for a long time. Permalloy MR sensors served as read heads in bubble memories of the early 1970s and they were also introduced, in order to replace the inductive elements, in the read heads of magnetic hard disk drives in 1991. This contributed to the substantial increase in the growth rate of the storage density [6]. By around 1997, in order to maintain the annual increase rate of the storage density, the AMR read heads were replaced by GMR heads which were based on the exchange-biased spin-valve structure invented by Dieny et al. [287] in 1990. This was of vital importance for creating a GMR-based device matching application-related requirements. First, we shall describe how GMR multilayers can be used for detecting magnetic field and, second, the GMR spin-valve structures will be introduced. For more detailed information about the possible applications of GMR devices, the interested readers should consult, e.g., the papers of Dieny [287] and Daughton [288]. 6.2.1. Field-sensitivity of GMR multilayers A basic requirement towards a magnetoresistive field sensor [288] is a nearly linear MR(H) curve which should provide, at the same time, a significant field-sensitivity (large change of resistivity upon a unit change of magnetic field to be detected). Further requirements are that the MR(H) curve should be a singe-valued function (absence of hysteresis), at least in the range of fields of interest, and the steepest section (high-sensitivity range) should fall possibly close to H = 0. In order to better visualize how the GMR effect can be used for magnetic field detection, we should first consider the MR(H) curves of two PD Co/Cu multilayers from the work of Mosca et al. [177]. Fig. 63 shows the MR(H) curves measured at T = 4.2 K for spacer layer thicknesses at the first and

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Fig. 63. Magnetoresistance curves of PD [Co(1.5 nm)/Cu(dCu)]30 multilayers at T = 4.2 K for (a) dCu = 0.9 nm (first AF maximum) and (b) 2.0 nm (second AF maximum). Reprinted from Ref. [177] with permission of Elsevier.

second AF maximum. At room temperature the GMR magnitude is reduced to about 50% and 20%, respectively (see Fig. 8), whereas the MR saturation fields remain practically unchanged. At the first AF maximum, due to the strong AF coupling, the magnetization reversal from the demagnetized zero-field state with high resistivity to the low-resistivity saturated state in high magnetic fields proceeds by the gradual alignment of all layer magnetizations (each layer can be considered as being in a single-domain state) along the external magnetic field. This results in an MR(H) curve (Fig. 63a) which shows a small splitting (hysteresis) only and a nearly linear monotonic decrease toward saturation. In addition, we can see that MR saturation against the strong AF coupling can be achieved at around 5 kOe only. Taking a room-temperature GMR of 50%, this yields a sensitivity of 0.01%/Oe. This is rather low by considering that the sensitivity of AMR sensors made of magnetically soft alloys can be as high as 0.4%/Oe [288] but at least the shape of the MR(H) curve at the first AF maximum of Co/Cu multilayers approximately fulfills the requirements for the field-dependence of a magnetoresistive sensor [288]. On the other hand, we can see in Fig. 63b that at the second AF maximum the saturation field is around 500 Oe only, i.e., reduced by a factor of 10 with respect to the first AF maximum whereas the room-temperature GMR magnitude is still as high as 20%. All this yields a room-temperature sensitivity of 0.04%/Oe at the second AF maximum. By choosing a magnetically softer alloy for the magnetic layer (Fe–Co, Fe–Ni or Fe–Co–Ni) in the multilayer structure while retaining the high GMR, the sensitivity can be tuned to about 0.1%/Oe for a GMR multilayer [288]. Even so, the sensitivity does not reach that of the AMR sensors; furthermore, the MR(H) curve is strongly nonlinear and it is not single-valued but exhibits a large hysteresis. The situation is very similar if we consider the room-temperature GMR characteristics of ED FM/NM multilayer films. With reference to Fig. 54, a GMR of about 20% for an ED Co–Ni–Cu/Cu multilayer with the thinnest Cu layers (below 1 nm) in a magnetic field of 8 kOe yields a sensitivity of 0.0025%/Oe only (although sensitivity is somewhat better in the low-field region); this value is even lower than the sensitivity that can be achieved with granular alloys [288]. On the other hand, for Cu layer thicknesses in the range of 3–5 nm where the multilayers exhibit MR(H) curves with significant hysteresis (see Figs. 41 and 56 for ED Co–Cu/Cu and Co–Ni–Cu/Cu multilayers, respectively), the GMR is about 10% in magnetic fields around 1 kOe. This yields a sensitivity of 0.01%/Oe, although sensitivity is much larger at the steepest part of the MR(H) curves. Gong et al [118,119] produced ED Fe–Co–Ni– Cu/Cu multilayer films and a magnetoresistance change of 8% was achieved over a magnetic field range of 200 Oe, yielding a sensitivity of 0.04%/Oe. 6.2.2. GMR spin-valve type structures and applications of GMR sensors Fig. 64 shows the structure of a GMR spin-valve device as well as the schematic indication of the variation of the magnetization and the magnetoresistance between the two extreme values of the

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Fig. 64. Cross-sectional view of a GMR spin-valve structure (left panel, left side) with the indication of the layer magnetization orientations in zero magnetic field (left panel, right side) and with the schematic change of the variation of the magnetization and magnetoresistance with magnetic field (right panel). Beside the M(H) function, the upper and lower arrows refer to the pinned and the free magnetic layers, respectively.

external magnetic field. The spin-valve structure [287] consists of an FM layer pair separated by a NM spacer (typically Cu) whereby a further AF layer (e.g., a Ni–Mn or Fe–Mn alloy) is deposited on top of the upper FM layer. Due to the strong direct exchange coupling at the interface between the AF layer and the upper FM layer, the magnetization of the latter always points either left or right (this depends on whether AF or FM coupling dominates at the interface) since a change in the magnetization orientation of the AF layer can occur in very high magnetic fields only. The spacer layer thickness is chosen in a way that the upper ‘‘fixed” FM layer is either only weakly coupled to the lower FM layer or the two FM layers are completely uncoupled. The lower FM layer is called ‘‘free” layer since an external magnetic field can easily change the orientation of its magnetization being uncoupled or only weakly coupled to the fixed layer. We can observe in Fig. 64 that in the spin-valve structure we get an MR(H) curve which exhibits a fairly linear, hysteresis-free steep section around H = 0 with high sensitivity. This advantageous feature is created via the exchange biasing of the fixed layer with the help of the AF layer and this bias is reflected in the shift of the position of the magnetization reversals along the field axis (exchange or unidirectional anisotropy) and resulting in the hysteresis-free reversal process at H = 0. The essence of this spin-valve structure is that there is a range of low magnetic fields where the two FM layers have an antiparallel alignment irrespective of whether there is an AF coupling between them. For a weak coupling or for the absence of a coupling between the FM layers, the spacer layer thickness is chosen to be at least 2 nm or higher. This choice also has the technological advantage that a better control of the spacer layer thickness can be achieved at higher spacer thicknesses. As to the FM layers, they are usually composed of soft magnetic alloys such as a Ni–Fe pair or a Co fixed layer and a Ni–Fe free layer. Daughton [288] provided a comparison of capabilities of various MR sensors which shows that the exchange-biased GMR spin-valve structure can exhibit a sensitivity of 1%/Oe, the largest among the MR sensors. It is also pointed out, however, that the various MR sensors are partly complementary to each other. Whereas the AMR sensors have a fairly large sensitivity, their saturation fields are very low (below 10 Oe). Therefore, in applications where larger fields are to be detected, GMR devices may be more favorable since for a relatively large magnetoresistance change the saturation field of a multilayer structure can be independently tuned by an appropriate choice of layer materials or layer thicknesses. Further particular layered GMR structures not yet mentioned are the so-called pseudo spin-valves without exchange bias. Such a structure can be either a sandwich or a multilayer containing FM

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and NM layers only with the peculiarity that two kinds of FM layers with different coercivities are present and there is no coupling between adjacent FM layers. Shinjo and Yamamoto [289] prepared a [Co(3 nm)/Cu5 nm)/Ni–Fe(3 nm)/Cu(5 nm)]15 multilayer by evaporation with magnetically hard (h) Co and soft (s) Ni–Fe layers. Such an FMh/NM/FMs/NM structure can also be termed as a ‘‘quartet” multilayer consisting of a periodic repetition of a layer quartet (in this sense, a conventional FM/NM multilayer with one type of the magnetic layer only is a duet multilayer). The M(H) and MR(H) curves of this structure are shown in Fig. 65. A significant hysteresis is present in both the magnetization and the magnetoresistance, very similar to that of Fig. 36b or Fig. 56 but with the magnetoresistance at H = 0 being almost as low as the saturation value. This also means that at the MR peaks there is a wide range of magnetic fields with antiparallel alignment of the adjacent layer magnetizations. On the magnetization curve (Fig. 65A), one can recognize a break point. For magnetic fields smaller than this value, the magnetization change is very sharp and this corresponds to the magnetization reversal of the soft layers at small magnetic fields whereas the magnetization change at a slower rate above the break point corresponds to the reversal of the hard layers. After complete reversal of the soft layers, the hard layers are still in their original saturated state so there is a range of magnetic fields where the two kinds of layers have an antiparallel alignment with respect to each other. Therefore, in the vicinity of the break point of the M(H) curve, the MR(H) curve exhibits a maximum. Since the break point is around 50 Oe, most of the total MR change takes place over a fairly small magnetic field range. Another approach to a pseudo spin-valve structure is an FMh/NM/FMs sandwich and the magnetization and magnetoresistance results of Dieny et al. [290] on a sputtered pseudo spin-valve sandwich structure Ni–Fe(4 m)/Cu(6 nm)/Ni–Co(4 nm) are shown in Fig. 66. Again, the break in the M(H) curve (around 3 Oe) indicates the completion of the magnetization reversal of the soft Ni–Fe layer, whereas we can observe that most of the magnetization of the hard Ni–Co layer reverses in a magnetic field around 13 Oe only. Therefore, the magnetization of the soft and the hard layer is antiparallel aligned in the field range in between, which results in a corresponding plateau in the MR(H) curve. The field-sensitivity of PD pseudo spin-valve sandwiches or multilayers can be as high as 0.2%/Oe, and they also can find their specific applications [288] in spite of the hysteretic behavior. 6.3. GMR results on ED pseudo spin-valve type structures Most of the attempts to prepare GMR spin-valve type sandwich structures by an ED process were carried out in the Co–Cu system [46,47,58,64,73,76,79], in each case from a pure sulfate bath in P/P mode and some of deposition parameters used in these works were summarized in Table 2. Attenborough and coworkers [46,47,64,73] performed electrodeposition at ECu(SHE) = 0.35 V which value corresponds well to the optimum deposition potential for the sulfate bath of Co and Cu (see Section 5.2.3.1). This also means that the nominal layer thicknesses specified correspond well

Fig. 65. Magnetic hysteresis curve (A) and magnetoresistance curve (B) of an evaporated ‘‘quartet” multilayer [Co(3 nm)/ Cu(5 nm)/Ni–Fe(3 nm)/Cu(5 nm)]15. Reprinted from Ref. [289] with permission of the Physical Society of Japan.

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Fig. 66. Magnetic hysteresis curve (dashed line) and magnetoresistance curve (solid line) of a sputtered pseudo spin-valve sandwich structure Ni-Fe(4 m)/Cu(6 nm)/Ni–Co(4 nm). Reprinted from Ref. [290] with permission of Elsevier.

to the actual thicknesses. The high Co2+/Cu2+ ionic ratio of the bath and the fairly negative Co deposition potential applied ensured a very low Cu content in the magnetic layer. A doped n-type GaAs(1 0 0) wafer was used as substrate and the layers were deposited directly on the wafer by through-plating. A symmetrical sandwich structure was designed the cross-section of which is shown schematically in Fig. 67a. This structure consists of a Co/Cu/Co/Cu/Co core with thin Co and Cu layers and this core is surrounded on each side by a thick Cu layer followed by a thick Co layer. The design idea was that the core forms an artificial antiferromagnet (AAF) in that the two thin Cu layers ensure a strong AF coupling between the three Co layers whereas the thick Cu layers decouple the outer thick Co layers from the core. In case of a strong AF-coupling in the core and with the outer Co layers decoupled from the core, then both the upper and the lower part of this structure resembles the exchange-biased GMR spin-valve depicted in Fig. 64. The M(H) and MR(H) curves measured on this ED sandwich structure are shown in Fig. 67b. The magnetization curve exhibits a very high remanence (Mr/Ms is close to 1), indicating that when the field is reduced from saturation to H = 0, practically all the layer magnetizations remain in their saturation states. A strong angular dependence of the MR(H) and M(H) curves in the multilayer plane indicated a significant in-plane anisotropy due to the epitaxial growth, resulting in a single-crystal behavior of the magnetic layered structure (this magnetic anisotropy of ED Co layers on the same substrate was also studied separately [291]). The large remanence can be ascribed to the well-defined in-plane anisotropy. The three steps which can be observed in the M(H) curves can be interpreted as arising from the subsequent magnetization reversals of the (i) top thick Co layer, (ii) the bottom thick Co layer and (iii) the three thin Co layers of the core. The slightly different growth conditions, i.e., the different microstructure of the top and bottom thick Co layers can properly explain their different but still closely spaced coercive fields, even if their thicknesses are identical. The somewhat larger coercive field of the thin Co layers in the core structure is accounted for by the layer thickness difference with respect to the outer Co layers as discussed in Section 5 (see Fig. 32 and Ref. [264]). At the same time, the fairly sharp outermost magnetization reversal step and its (low) coercive field (around 100 Oe) suggests that there is no AF coupling between the Co layers of the core. Actually, this cannot even be expected since the Cu spacer thickness in the core is 3.2 nm at which value the AF coupling is fairly weak even in PD Co/Cu multilayers. Furthermore, in Section 5.2.7, a lack of AF coupling was concluded for Cu spacer thicknesses at this value in ED Co/Cu multilayers. The absence of AF coupling in the core was demonstrated also by Attenborough et al. [46] in that the remanence was found to be fairly high (Mr/Ms = 0.81) for an ED [Co(2.7 nm)/Cu(3.2 nm)]20 multilayer grown on the same substrate under identical conditions. All this means that the designed core structure does not represent an AAF system, just a set of three thin and uncoupled Co layers. The total thickness (8.1 nm) of the Co layers in the core is close to the thickness of the outer Co layers (10 nm), well explaining the fact that each of the three subsequent magnetization reversal steps has practically the same height.

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Fig. 67. (a) Cross-sectional view of the ED sandwich structure investigated by Attenborough et al. [46,47,64,73]. In the Co/Cu/ Co/Cu/Co core structure the Co and Cu layer thicknesses are 2.7 nm and 3.2 nm, respectively. The thicknesses of the outer Co and Cu layers are 10 nm and 4.8 nm, respectively. The top Cu layer is deposited to prevent oxidation; (b) Magnetization (dashed line) and magnetoresistance (solid lines) curves of the ED sandwich structure. The arrows along the MR(H) curve indicate the magnetic field cycling sequence. Reprinted from Ref. [64] with permission of Elsevier.

Having in mind the individual magnetization reversal steps, the observed field dependence of the magnetoresistance (Fig. 67b, solid line with arrows) can also be understood now. The high remanence after saturation implies that at H = 0 all the layer magnetizations are still parallel aligned and, thus, the MR value should also be very small at H = 0 as actually observed. The first drop in the magnetization is accompanied by a rise of the MR(H) curve due to an antiparallel alignment between the top thick Co layer and the Co layers in the core. A similar MR(H) increase follows at a slightly higher magnetic field upon the reversal of the bottom thick Co layer whereas the core magnetic layers are still in their originally saturated orientation. A high MR value (a plateau) persists until the magnetizations of the core Co layers are also reversed and at this magnetic field, the MR value suddenly drops. A comparison with Fig. 66 clearly reveals a similarity with FMh/NM/FMs pseudo spin-valve structure of Dieny et al. [290]. It can be established that whereas the original idea of an exchange-biased GMR spin-valve structure could not be realized by Attenborough and coworkers [46,47,64], also demonstrated by the symmetrical nature of the M(H) and MR(H) curves (see Fig. 67b), they definitely were able to produce a highquality magnetic/non-magnetic layered sandwich structure by an ED process. This sandwich exhibited a GMR magnitude of about 5%, a low saturation field (about 100 Oe) and very steep MR(H) changes, the latter yielding a field-sensitivity of about 0.55%/Oe) which is the best value among ED layered structures (by patterning in a top-down approach the ED sandwich to micrometer-scale wide strips, an even larger sensitivity of 1.5%/Oe was achieved [47]). As noted above, due to the symmetrical MR(H) curve, the ED sandwich structure studied by Attenborough and coworkers [46,47,64,73] cannot be considered as an exchange-biased GMR spin-valve but rather a pseudo spin-valve with magnetic layers of different coercivity. Actually, the same is true for all the ED sandwich structures to be described later in this section. Thus, the preparation of an exchange-biased GMR spin-valve structure by an ED process is still awaited. Shima et al. [58] investigated the magnetic and magnetoresistance behavior of the same sandwich structure with approximately the same layer thicknesses as studied by Attenborough et al. [46,47,64]. The deposition conditions used by Shima et al. [58] were summarized in Table 2. The Cu deposition potential was chosen just a bit more negative than the optimum value and, in assessing the layer thicknesses, the actual current efficiencies were also taken into account. A GaAs(0 0 1) substrate was used and electrodeposition was carried out by through-plating directly on the wafer. The M(H) and MR(H) curves were measured along the h1 1 0i and h1 0 0i directions of the wafer substrate. High remanence (Mr/Ms  1) and sharp magnetization reversals around 50 Oe were observed for both orientations. For H > 50 Oe, the magnetization change slowed down and reached saturation at about 200 Oe. The strongly different rate of the magnetization change below and above 50 Oe indicates that

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the thick and thin Co layers undergo a magnetization reversal in different field ranges. The MR(H) curves exhibited a behavior similar to that shown in Fig. 65b. Two sharp MR peaks with a plateau-like behavior were observed in the field range of the sharp magnetization reversal, indicating an AP alignment of the thin and thick Co layers. The MR peak values were slightly above 1%, with a negligibly small magnetoresistance in the field range between the peaks. Pasa and coworkers [76,79] investigated the GMR of ED sandwich structures prepared in P/P mode directly on n-type Si(1 0 0) wafers by through-plating. As it can be seen from the preparation conditions in Table 2, the Cu deposition potential applied [ECu(SHE) = 0.16 V] was much more positive than the optimum value. This means that a significant Co dissolution during the Cu deposition pulse cannot be excluded which may have an impact on the actual layer thicknesses with respect to the specified values. The following layer structures were investigated: Si/Co(27.5 nm)/Cu(6.5 nm) [Co(3.5 nm)/Cu(5.5 nm)]n/Co(25.5 nm)/Cu(0.7 nm). They correspond again roughly to that of Attenborough and coworkers [46,47,64] except that Pasa and coworkers used a stack of bilayers with various numbers. The final thin Cu layer served also here as a capping layer to prevent oxidation. The bilayer number in the sandwich core was n = 3, 5 and 8. The observed MR(H) curves indicated a GMR of about 8%, the value depending somewhat on the number of bilayer stacks in the sandwich core, and saturation was around 500 Oe. The MR(H) curves exhibited splitting with a shape characteristic of a pseudo spin-valve structure. Yamada et al. [91] investigated the magnetic and magnetoresistive properties of an ED quartet multilayer [FMh/Cu/FMs/Cu]5. Electrodeposition was carried out under galvanostatic control from a bath containing sulfates of Ni, Fe and Cu with an ionic ratio Ni2+:Fe2+:Cu2+ of 500:10:5. After establishing the dependence of deposit composition obtained by d.c. plating from this bath, appropriate current pulse amplitudes were chosen to deposit the individual layers. At the lowest current density (0.15 mA/ cm2), an almost pure Cu layer (Cu98Ni2) was deposited. By applying current densities of 5 and 25 mA/ cm2, a Permalloy type soft magnetic layer (Ni78Fe14Cu8) with a low coercivity and a magnetically harder Ni-rich alloy (Ni93Fe4Cu3) with three times higher coercivity were deposited, respectively. For such a quartet multilayer, a separate switching of the two kinds of magnetic layers could be observed by measuring the hysteresis loop. The MR(H) curves exhibited a splitting and showed saturation at the same magnetic field as the magnetization. The dependence of the magnetoresistance was studied as a function of the layer thicknesses. A maximum GMR of 3% was achieved for the nominal thicknesses 5 nm (Cu layer), 3 nm (hard layer) and 3 nm (soft layer). This room-temperature GMR value of 3% increased to 9% at T = 5 K. It was also attempted to increase the coercive field of the hard magnetic layer by depositing it with a pulse train, i.e., replacing the single hard magnetic layer by a Ni-rich alloy/Cu multilayer. Although the coercive field could be increased in this manner by a factor of three, the GMR remained the same as before. A very similar quartet multilayer was developed by Yao and coworkers [140,143] who fabricated a [Cu/Co/Cu/Ni80Fe20]30 sequence by dual-bath electrodeposition on an n-Si(1 1 1) wafer by throughplating. The wafer was first covered with an ED buffer layer of Ni80Fe20 under a constant potential of ENiFe(SHE) = 0.86 V from an acidic bath (pH of 3.0) composed of NiSO4, FeSO4, NiCl2, Na-citrate and H3BO3. Thereafter, this substrate was transferred to another electrolyte (pH of 6.0) containing a mixed solution of CoSO4, CuSO4, Na-citrate and NaCl and a Cu/Co/Cu trilayer stack was grown on it by two-pulse plating in P/P mode. By returning the substrate to the Ni–Fe bath, another Ni80Fe20 layer was deposited and the whole process was repeated 30 times. On the basis of cathodic polarization curves, the deposition potentials were chosen as follows: ECu(SHE) = 0.16 V for Cu deposition on Ni–Fe, ECu(SHE) = 0.31 V for Cu deposition on Co and ECo(SHE) = 0.81 V for the deposition of Co. The XRD pattern presented revealed a strong fcc(1 1 1) texture for an ED stacking of Si/Ni80Fe20(25 nm)[Cu(3.6 nm)/Co(1.2 nm)/Cu(3.6 nm)/Ni80Fe20(2.8 nm)]30. Two first-order satellites could be identified which gave a bilayer repeat of 11.8 nm, comparing well to the nominal value of 11.2 nm. The MR(H) curves of these quartet multilayers were similar to those shown in Fig. 65b for pseudo spin-valves (the two separate coercive fields could be seen in the M(H) curves for relatively thick Cu layers only) and saturation could be reached in magnetic fields around 350 Oe. The GMR magnitude was measured as a function of the constituent layer thicknesses, including the buffer layer. For a given stacking, the GMR was about 2% at 15 nm buffer layer thickness whereas the GMR increased to 5% at 25 nm thickness. An AFM study revealed that the roughness values on the surface

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of the multilayer stack were 7.9 nm and 4.3 nm, respectively, for the above two buffer layer thicknesses. At the lower buffer thickness, actually incomplete substrate coverage was observed. For buffer layer thicknesses above 25 nm, the roughness increased again and was accompanied by a decrease of GMR. These results indicate that the growth of high-quality ED multilayer films directly on a semiconductor substrate still involves a lot of challenges for improving their GMR performance. The overall evolution of GMR with magnetic layer thicknesses was similar to that presented in Figs. 36a and b, the maximum GMR (about 5%) achieved at about 1.2 nm for the Co layer and 3 nm for the Ni80Fe20 layer. The dependence of the GMR on Cu layer thickness was also very similar to that shown in Fig. 47 for ED Co–Cu/Cu multilayer films, with the largest GMR (5%) obtained at about 3.5 nm Cu layer thickness. 6.4. GMR results on application-oriented ED multilayer films In this section, reports on GMR studies of ED multilayer films will be described in which multilayers were prepared with the purpose of prospective application in magnetic field sensing devices. Following an earlier work by Parkin [292] who produced a flexible GMR sensor by sputtering an exchange-biased spin-valve structure on various organic films, Yan et al. [92] attempted the preparation of a flexible GMR sensor by electrodepositing a Co–Cu/Cu multilayer film on a conducting polymer substrate. For this purpose, a polypyrrole (PPy) conducting polymer was chosen which, especially when heavily doped with m-sulphobenzoic acid, exhibited a sufficiently high conductivity and high stability even in a reductive agent (such as NH3OH) and at a negative potential necessary for the electrodeposition of a multilayer on its surface. The PPy polymer itself was also deposited and polymerized by electrochemical methods on a stainless steel substrate under galvanostatic conditions. At the applied current density of 1 mA/cm2, the PPy growth rate was about 10 lm/h. The electrodeposition of the Co–Cu/Cu multilayer films was then carried out on the properly cleaned stainless steel/PPy substrate in P/P mode with ECo(SHE) = 0.66 V and ECu(SHE) = 0.16 V from a nitrogen-purified bath with an ionic ratio Co2+:Cu2+ of 200:1 (the anions were not specified and a buffering agent was used). In specifying the layer thicknesses, the actual current efficiencies were also taken into account on the basis of data obtained on d.c.-plated deposits. However, due to the too positive ECu value with respect to the optimum (see Section 5.2.3.1), a strong Co dissolution and its replacement by Cu certainly resulted in smaller Co and larger Cu layer thicknesses than those specified in the paper. It was established from an X-ray photoelectron spectroscopy (XPS) study that the magnetic layer composition was about Co95Cu5. After the electrodeposition of the multilayer films, the PPy/[Co–Cu/Cu]30 composite cold be easily peeled off from the stainless steel substrate for further investigations. The MR(H) curve did not show a splitting and a GMR of 4.5% was reached at 10 kOe whereas most of the resistance change occurred up to about 5 kOe. As a possible reason for the relatively low GMR, the authors mentioned that the PPy film, at a typical thickness of about 5 lm, exhibited a fairly rough surface as seen by SEM. Another reason mentioned by the authors is the conductivity of the PPy substrate itself which can partly shunt the Co–Cu/Cu multilayer films and, partly, due to its not sufficiently high conductivity, may cause a highly inhomogeneous current density distribution which then results in a non-uniformity of the layer thicknesses over the cathode surface area. In spite of the relatively low GMR achieved in the study of Yan et al. [92], such flexible sensors composed of multilayers with improved GMR characteristics can be useful for applications. The multilayer growth rate by the ED process can be fairly high and Yan et al. [92] described that the flexible sensor can be formed into various shapes and is easy to cut by a simple tool as a scissor. The authors of this paper have also carried out several mechanical tests on the PPy/multilayer composite. The mechanical properties of the composite were found practically to be those of the PPy substrate itself and were found to be almost constant in the temperature range from 20 °C to +50 °C. There were, furthermore, some reports [63,101,135] in which patterned ED Co–Cu/Cu multilayer films with GMR were prepared by using two different approaches. Zhang and coworkers [63,101] applied a top-down approach in which an ED Co–Cu/Cu multilayer film was first prepared in P/P mode on an oxidized Si(1 0 0) wafer covered with a sputtered Ti(10 nm) and a Cu(20 nm) metallic layer beforehand. The bath was deaerated before plating and contained CoSO4, CuSO4 and H3BO3 as buffering agent; the pH was adjusted to 5.6 with the addition of NH3

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solution. The Co2+:Cu2+ ionic ratio was 1500:8 and the applied deposition potentials were ECo(SHE) = 0.86 V and ECu(SHE) = 0.32 V. The latter value almost corresponds to the optimum Cu deposition potential, although the authors observed an anodic current transient indicating some degree of Co dissolution when switching from the Co deposition potential to the Cu deposition potential. Deposition was started with a Cu layer and typically 100 bilayers were deposited with a total thickness of about 0.5 lm. A lithographic technique was then applied to pattern the GMR multilayer film of original lateral size 10 mm by 4 mm into the shape of a continuous zigzag wire shown schematically in Fig. 68. Microwires with a width of 10 lm were formed along the longer edge of the deposit and were connected in series through the short perpendicular section, with a total wire length of about 150 mm which effectively increased the total resistance of the microdevice to be measured. During the lithographic process, also the Ti and Cu layers were etched down to the Si wafer in the regions between the wires. An XRD study revealed a strong fcc-(1 0 0) texture and first-order satellites around the main peak indicated a fairly good multilayer structure. The MR(H) curve measured on a zigzag pattern with the multilayer [Co(3 nm)/Cu(2 nm)]100 showed an almost linear decrease of resistivity up to the maximum field applied (3 kOe) where a GMR of 4% was obtained. Although no MR(H) curves were reported for the ED multilayer films before the patterning process, a GMR of 2% measured at 100 Oe was specified for an unpatterned deposit containing the same multilayer film as the zigzag wire pattern. The same reduction of sensitivity was obtained also for 3–5 times larger microwire widths. The authors considered that the deleterious effect of patterning may come from the elevated temperatures of the baking process during patterning the GMR multilayer film. Li et al. [135] used a bottom-up approach in that first a patterned substrate was created by lithography and the GMR multilayer film was then electrodeposited on this pattern. A similar zigzag structure with equal wire width and gap distance from 50 lm to 200 lm was defined by pattering as used by Yan et al. [101]. Li et al. [135] prepared ED Co–Cu/Cu multilayer films in G/G mode from an electrolyte containing CoSO4, CuSO4 and H3BO3 and with an ionic ratio Co2+:Cu2+ = 100:1 under quiescent conditions at a pH of 3.0 on an n-type Si(1 0 0) wafer with sputtered Ti(20 nm) and Cu(100 nm) metalizing layers. The magnetic and non-magnetic layers were deposited at current densities of 20 mA/ cm2 and 0.2 mA/cm2, respectively. The ED process was carried out in recessed areas defined by a remaining resist of the patterning process which covered the metal-free regions of the wafer; after electrodepositing the multilayer, this resist was also removed. Electrodeposited [Co–Cu(2.5 nm)/ Cu(3 nm)]1000 multilayers with a total thickness of about 5 lm were prepared on the zigzag pattern. The reported MR(H) curves were similar to those shown in Fig. 39: the magnetoresistance did not saturate up to 10 kOe where a GMR of about 2% was achieved. This is about half of the GMR value measured on identically prepared multilayers on the same substrate without patterning, as the authors noted. A SEM study revealed a significant edge effect on the multilayers electrodeposited on the patterned substrate that was more pronounced for narrower wires. This may be one of the factors contributing to the loss of the GMR in the patterned microdevice. A pulse train was also applied to deposit the magnetic layer but the MR(H) curves remained the same as with a single pulse and the GMR increased by about one third only.

Fig. 68. Schematic diagram of an ED GMR multilayer sensor with a zigzag wire pattern. Reprinted from Ref. [101] with permission of Springer Science + Business Media.

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7. Summary of progress achieved and problems ahead with GMR in ED multilayer films Although a vast amount of literature is available on the GMR study of ED multilayered films and sandwich structures [9–149] which were critically reviewed in Sections 5 and 6.3, respectively, the deposition conditions have not been optimized for each FM/NM multilayer systems which can be produced at all by the ED process. This is especially the case for those multilayer systems which turned out, when prepared by PD methods, to exhibit GMR characteristics attractive for applications. At the moment, one cannot see any obvious reasons why it could not be achieved to produce ED multilayers with GMR properties being competitive with those of PD counterparts. It is also evident, however, that in order to exploit the advantage of ED multilayers concerning their cheap and simple production, significant efforts should still be made until sandwich structures or multilayers with a low bilayer number, with a total thickness of a few tens of nanometers, can be manufactured in a reliable and reproducible manner. In current applications, GMR structures fabricated by PD methods and with a total thickness of at most a few tens of a nanometer are used exclusively. Alternatively, in case of advantageous GMR parameters of ED multilayer films, new applications should be sought for in which the cheap mass production of thick multilayers can represent an attraction. First of all, achievements are listed below by ordering the information in a sequence going from macroscopic to microscopic properties: (i) Macroscopic homogeneity of the multilayer film. The electrolyte solution used for the deposition of multilayer films is a reactive environment that itself also changes upon the deposition process. This gives rise to problems with temporal and spatial steadiness of the deposition conditions. The bath in the vicinity of the cathode becomes depleted with respect to the reacting components. Therefore, the composition changes in time, which has the largest influence on the deposition of the most dilute component. This is usually the NM metal. This is why it is a prerequisite for the homogeneity of the multilayer deposit in the growth direction that the quantity of the NM metal deposited within a single pulse is controlled. Time-based control is generally insufficient, and a charge-based control under optimized potentiostatic conditions is required. In contrast, the lateral homogeneity of the multilayer deposit is related to the cell geometry rather than the steadiness of the deposition conditions. A significant edge effect may occur if the deposition is carried out in a non-confined solution close to the cathode. Besides the reduction of the edge effect by means of the cell construction, a good care has to be taken to provide an even current distribution by positioning the cathode and the anode parallel to each other. Another factor influencing the lateral homogeneity of the deposit is the position and orientation of the cathode. If it is horizontal, the flow of the solution induced by either the gravity effect (due to the decrease in concentration and hence the decrease in density) or natural convection (caused by the hydrogen evolved in the side reaction) has the same impact on the deposition conditions at each point of the cathode. It turned out that this arrangement provides good lateral deposit homogeneity over the cathode surface area in contrast to a vertical cathode orientation. (ii) It is required that not only the homogeneity in both lateral and growth directions is good, but that the layer undulation and interface roughness are also as small as possible. For the sake of a fairly large planarity of the subsequent layers, the impact of the reactive environment on the layers already deposited has to be minimized. This is particularly important for the dissolution/corrosion of the less noble FM metal(s) during the deposition of the NM spacer layer. This can be ensured only by a potentiostatic deposition of the NM layer and by a careful optimization of the deposition potential of the NM metal. The planarity of the layers at a much larger lateral scale than the bilayer thickness is greatly influenced by the initial quality of the substrate. It was shown that substrates with very low initial roughness yield multilayers with higher GMR than those with a large surface roughness. This is also a step toward the integration of multilayer electrodeposition into the silicon wafer technology.

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(iii) It was established that complexing agents and other organic compounds classified as surfactants, levelers and brightener agents are all deleterious for the GMR of multilayer deposits. This can be well explained with the grain-refining effect of these additives. Upon grain refinement, the resistivity increases due to a larger contribution from grain-boundary electron scattering and hence the electron mean-free path is reduced. (iv) Besides the application of general trends previously established for the multilayer electrodeposition process, several detailed structural studies underpinned the relationship between the microstructure of the deposit and their GMR properties. Impact of the misfit of the constituents, the formation of twins and dislocations as well as the interrelation of in-plane and out-of-plane atomic distances are now understood. (v) Some asymmetry is observed in the behavior of the Co and Cu layer growth in Co–Cu/Cu multilayers. The desired FM behavior of the multilayer can be achieved with d(Co)  1 nm, and this lower limit does not depend too much on the electrolyte composition if FM behavior can be achieved at all with a particular bath. However, the threshold thickness of the Cu layer where AMR turns to GMR behavior is larger and depends on the bath used. On the one hand, this asymmetry can be explained by the difference of the lattice plane distances of the constituent metals and the multilayer. Since the lattice plane distances in the multilayer are intermediate between those of pure Co and Cu, the latter metal exhibiting the larger lattice plane distances the nucleation on the surface of Co with smaller lattice plane distances is more difficult than vice versa. Therefore, the growth of the Cu layer follows the Volmer–Weber type growth (formation of islands which then coalesce and, thus, a complete coverage is achieved at high average deposit thickness only), whereas the growth of the Co layer is much closer to the Frank–van der Merwe type growth mode where a complete coverage of the substrate surface is achieved already at a low deposit thickness. On the other hand, the conclusions drawn for the deposition mode are a bit indirect. Although the above explanation is in full agreement with the experimental data, it has to be noted that the presence of pinholes in the FM and in the NM layers has different consequences on the macroscopic properties observed. The pinholes in the NM layer immediately manifest themselves by giving rise to an FM coupling and hence to an AMR type magnetoresistance behavior. However, the pinholes in the FM layers have no significant impact on the magnetic or magnetoresistance behavior as long as their size and density is low enough to exclude the magnetic fragmentation of the FM layer into SPM regions. (vi) Concerning the quantitative description of the magnetic and magnetoresistance behavior of ED multilayer films, it was shown that the large saturation fields often observed at small layer thicknesses can be attributed to the SPM behavior of the multilayers. The method applied for the description of granular magnetic materials was adapted to multilayers, too. By this method, both the magnetization and the magnetoresistance curves (M(H) and MR(H), respectively) can be decomposed into FM and SPM contributions. The origin of these contributions can be ascribed to the presence of various kinds of spin-dependent scattering events. The temperature dependence of the magnetization and magnetoresistance behavior of Co-Cu/ Cu multilayer films could be elucidated only by the adaptation of the model of interacting SPM entities. This explained why the relative SPM contribution remains large down to the liquid helium temperature range, although the blocking of the SPM particles should also lead to a vanishing SPM contribution. The latter behavior was indeed observed for Ni–Cu/Cu multilayers. The difference in behavior for the two multilayer systems was also explained by taking into account the miscibility properties and the dependence of the Curie temperature on the composition. (vii) It was shown that no significant AF coupling exists between the magnetic layers of electrodeposited Co–Cu/Cu multilayers. A similar situation should probably hold for all ED multilayer systems, although the amount of experimental data available is still insufficient to make a general statement.

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A list of problems to be dealt with in order to achieve significant progress is provided below. While some of them should just be studied and have hitherto remained hidden rather by the negligence of the researchers, a few problems may prove to be very hard and the outcome of finding a proper solution to them can be quite uncertain. (i) It was shown that a horizontal cathode yields multilayer deposits with a good lateral homogeneity. However, this study was performed with upward-facing cathodes only. In this case, the cell construction is perfect for a lab scale research task, but it may prove to be inadequate for a large scale (mass production) where an immersion system could be more cost-effective. It would be worthwhile to study if the same multilayer quality can be achieved with downward facing cathodes, especially at small overall deposit thicknesses. (ii) The purity of the chemicals used for electrodeposition was seldom controlled or was completely neglected. While the electrodeposition of Cu interconnects in microelectronics was possible only with a rigorous control of the amount of additives and purity, such studies have not yet been performed for ED multilayers. The neglect of this issue can be likely a reason for the fact that for the same multilayer systems, research groups using similar electrolytes often publish not only different results but even contradicting trends. (iii) Other macroscopic conditions of the multilayer deposition have not been addressed either. For instance, it is unclear if the electrolyte flow can be optimized for obtaining multilayers with high GMR and conserving the lateral homogeneity. Another issue is the impact of the deposition temperature. Should it be surprising but no study can be found where the deposition from the same electrolyte was studied as a function of temperature. The expectation is that, by increasing the temperature, the surface diffusion distance of the intermediates/adatoms increases, hence leading to the growth of larger crystals, as established for electrodeposition in general. The grain-coarsening might lead to an increase in GMR by reducing the grain-boundary scattering contribution to the background resistivity. (iv) The deposition conditions of the NM component, especially its deposition potential, should be optimized for various electrolytes. This is the only manner by which the performance of different electrolytes can be compared. (v) The deposition of the NM layer can be optimized with the principles of electrochemistry when it is carried out in the P mode, while it cannot be done in the G mode. The mass production might require the regulation of the current rather than that of the electrode potential. It is unclear yet if there is a suitable optimization method to do so. (vi) For designing a bath for multilayer deposition, it would be important to relate the threshold Cu layer thickness (above which the GMR behavior is predominant) to elementary properties of the bath. For the time being, it is impossible even to predict what kind of properties should be sought for establishing such a relationship. A similar problem occurs with the growth mode. For a large GMR, the layer-by-layer growth of both types of layers would be favorable. Apart from the general principle of electroplating that the molecules strongly interacting with the metal surface hinder the surface diffusion and, therefore, lead to the increment in nucleation rate and the suppression of grain growth, there is no concept for bath development in which the layer-by-layer growth is preferred. The experimental methods for studying atomistic steps of electrodeposition have yet to improve in order to be able to make predictions on the growth mode on the basis of other elemental parameters. (vii) Although the impact of organic surfactants is generally deleterious for GMR, the PD methods take the advantage of elements that have surfacting effect different from the organic ones. These metal atoms form a kind of ‘‘floating” monatomic layer over the sample surface. The result is that the growth pattern with nucleation and grain coalescence turns into another growth mode where the nucleation step is almost eliminated and the adatoms are immediately built in the existing grains, conserving the top floating layer at the same time.

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This process surely involves the penetration of adatoms through the adlayer to the parent metal. It is a great challenge whether a similar process can be developed for the electrodeposition, too. (viii) The structural features of some electrodeposited multilayer systems have been widely investigated in the past decade. However, it is not yet understood why no antiferromagnetic exchange coupling exists between adjacent magnetic layers in ED multilayers which is probably the main reason for the absence of oscillatory GMR in this kind of multilayers. The clarification of this issue requires an extensive study in the future. (ix) The mathematical analysis of the M(H) and MR(H) curves revealed that they can be well elucidated by means of a Langevin-type function added to the ferromagnetic ‘‘background”. Nevertheless, it still remained a problem what kinds of objects (regions, particles, layers, segregations, etc.) can be responsible for the SPM behavior. The application of sophisticated diffraction and electronmicroscopic techniques as well as three-dimensional atomic probe methods will certainly be needed to solve this problem.

8. Concluding remarks In this paper, an attempt was made to give an overview on the current status of the preparation and structure as well as the magnetic and magnetotransport properties of ED multilayer films exhibiting GMR. Nearly one and a half hundred reports have been published on ED GMR multilayer films since the first paper in the field. These reported results were critically evaluated for each multilayer system accessible at all for preparation by electrodeposition. Whereas the basic features of GMR are the same for both PD and ED multilayer films, there are some aspects of the GMR characteristics which are naturally specific to the ED method. These specific issues related to the ED multilayer films include mainly the frequently observed high saturation field of the magnetoresistance and the absence of an oscillatory GMR as a function of the spacer layer thickness, the latter intimately connected to the absence of an AF coupling between adjacent layers in the ED multilayer films. Evidently, these specific features derive from the deposition conditions governed by electrochemical processes not yet completely understood and controlled. It appears from the wealth of literature of GMR in ED multilayer films that the experimental approach has hitherto been mostly biased in a sense that either the magnetic and magnetotransport properties were extensively investigated without a thorough consideration of the underlying electrochemical processes or a very thorough electrochemical characterization of the electrolyte used and the deposit formed was carried out which was, then, only ‘‘decorated” with some MR measurements. It was our intention, therefore, to make a unified approach by focusing on both aspects, namely on the physical processes underlying the GMR effect in various magnetic nanostructures on the one hand and on the influence of electrochemical conditions on the formation of various magnetic nanostructures on the other hand. This was considered the most viable approach since, evidently, the transport and magnetic properties, the microstructure and the electrochemical deposition conditions are strongly interrelated. It is hoped that the dual type of approach of this review will prove useful for those who intend to work in this field in the future. The text was intended to be made sufficiently self-containing and independent in the sense that a newcomer researcher to the field can profit out of it fairly quickly with a basic background of magnetism and electrochemistry. An experienced researcher of this specific topic can enjoy the systematic treatment of magnetotransport properties and electrochemical processes, both subfields on their own footing, and can notice commonalities and specialties which are although often contained in individual research papers but also often remain hidden among the wealth of details but can be brought more explicitly to light when presented in such a summary form. Finally, it is also the authors’ hope that individual sections alone can bring enlightening to the interested readers, even if not intending to work in this particular field. It would be the outermost pleasure if readers could enjoy the combined approach from both the magnetism and electrochemistry side with enthusiasm as we did when working in this field in the last decade and in writing this review. Anyway, it is

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