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Cu multilayers with giant magnetoresistance
Thin Solid Films 424 (2003) 229–238
Structural evolution during growth of electrodeposited Co–CuyCu multilayers with giant magnetoresistance ´ Cziraki ´ a, L. Peter ´ b, B. Arnoldc, J. Thomasc, H.D. Bauerc, K. Wetzigc, I. Bakonyib,* A. a
¨ ¨ University, P.O. Box 32, H-1518 Budapest, Hungary Department of Solid State Physics, Eotvos Research Institute for Solid State Physics and Optics, Hungarian Academy of Sciences, P.O. Box 49, H-1525 Budapest, Hungary c ¨ Festkorperanalytik ¨ ¨ Festkorper-und ¨ Institut fur und Strukturforschung, Institut fur Werkstofforschung, Helmholtzstrasse 20, D-01069 Dresden, Germany b
Received 19 March 2002; received in revised form 11 October 2002; accepted 24 November 2002
Abstract The maximum room-temperature giant magnetoresistance (GMR) of electrodeposited Co–CuyCu multilayers produced during this work was approximately 9% at 8 kOe, and it was found to decrease with increasing bilayer repeat number. A transmission electron microscopy study has revealed the fine details of the microstructure formed during growth. At the beginning of the deposition very small, nano-sized crystallites formed with both hexagonal close-packed (hcp) and face-centred cubic (fcc) crystal structures containing a high level of internal stress. The Cu-content of these small crystallites was found to depend strongly on their crystal structure (fcc or hcp). After this initial polycrystalline region, the size of crystallites increases, forming an fcc superlattice with increasing average Cu concentration at the first hundreds of repeat periods. This increase is not monotonous across the whole sample thickness. As another effect, the bending of layer planes becomes more remarkable as the growth progresses. The above inhomogeneities formed during the deposition of hundreds of bilayers could be responsible for the decrease in GMR with increasing total thickness of the multilayered samples. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Multilayers; Structural properties; Transmission electron microscopy; Giant magnetoresistance
1. Introduction The discovery of giant magnetoresistance (GMR) observed in magnetic multilayers w1,2x has opened up a new material class for magnetic devices, such as magnetic read out heads, magnetic sensors and other information storage and retrieval devices. Exhibiting very large GMR (up to 50%) at room temperature w3–5x, the Co–Cu system is one of the most important among the taylored magneticynon-magnetic multilayers. A simple and cheap way to produce relatively thick self-supporting multilayered films in the Co–Cu system is electrodeposition. Nevertheless, as a result of the nature of the deposition method, the magnetic layers cannot consist of pure Co but rather a Co-rich Co–Cu alloy with *Corresponding author. Tel.: q36-1-392-2628; fax: q36-1-3922215. E-mail address: [email protected] (I. Bakonyi).
magnetic properties very similar to those of pure Co. The GMR values measured on these electrodeposited Co–CuyCu multilayered films have remained below 20% w6–19x, regardless of whether they were deposited with the control of potential, current or both alternating. In our recent studies of electrodeposited Co–CuyCu multilayers w17,19x, the aim was to clarify the deposition parameters that govern the GMR value. Direct-current plating, pulse-plating, two-pulse plating and reversepulse plating were used to produce relatively thick selfsupported electrodeposited Co–Cu alloys and Co–Cuy Cu multilayers under galvanostatic control from an electrolyte containing CoSO4 and CuSO4. It was found w17x that the maximum room-temperature GMR of these multilayer deposits was approximately 9% at 8 kOe and it decreased with increasing bilayer repeat number. This gradual deterioration of the GMR was ascribed to the increasing structural disorder developing with increasing
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total multilayer thickness (bilayer repeat number) as revealed from a preliminary structural investigation w17x by atomic force microscopy, X-ray diffraction (XRD) and transmission electron microscopy (TEM). According to these studies, the direct-current plating resulted in a Co95Cu5 alloy with nearly equal amounts of facecentred cubic (fcc) and hexagonal close-packed (hcp) phases, while all pulsed-current methods yielded Co– CuyCu multilayers which exhibited a dominant fcc structure and contained hcp crystals in an amount of less than 2% only. The appearance of satellite peaks on both sides of the (1 1 1) Bragg maximum in the large-angle XRD spectrum confirmed the formation of a multilayered structure for each pulsed electrodeposition mode. It was furthermore found that there was a strong (1 1 1) crystallographic orientation at the substrate side of the samples and a loss of the preferred orientation towards the solution side in correlation with the observed decrease in the GMR values. The loss in GMR was regarded as a consequence of the structural evolution as the specimen grew, but this assumption is to be verified by using a technique that can directly detect structural features. On the other hand, it is well known that the structural parameters measured with XRD can be regarded as mean values only. XRD has a very limited ability to detect local concentration inhomogeneities and is not sensitive enough to reveal the local lattice distortions, like stacking faults, twin boundaries and bending of the layers. Therefore, in order to better reveal the nature of the microstructure forming in these multilayers, and especially their evolution during growth, a more detailed TEM study has now been performed. This is a unique direct method to visualise the microstructure at the atomic level. The results of these TEM investigations including elemental analysis on the nanometer range will be presented in this paper in order to complete our previous findings w17x mainly on the same electrodeposited Co–CuyCu multilayer samples. 2. Experimental Two-pulse plating from a single electrolyte containing CoSO4 and CuSO4 was used to produce electrodeposited Co–CuyCu multilayers under galvanostatic control on a polycrystalline Ti substrate as described in more detail in Ref. w17x. The deposits were peeled off from the substrates by mechanical stripping. The present detailed TEM studies were performed on electrodeposited multilayers of the form wCo95Cu5(dCo–Cu)yCu(dCu)xN where the nominal magnetic (dCo–Cu) and non-magnetic (dCu) layer thicknesses were in the ranges dCo–Cus3.6–12 nm and dCus0.22– 1.8 nm and the bilayer repeat number N varied from 28 to 1500. Most of the multilayer samples were identical
with those of Ref. w17x. Among them, the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer had the largest room-temperature GMR of approximately 9% at 8 kOe and, for this reason, TEM results on this sample will be presented in more detail. For a better understanding of the initial stage of deposition, a wCo95Cu5(3.6 nm)y Cu(1.1 nm)xNs28 multilayer was also prepared under identical conditions for the present TEM studies. Due to the small bilayer repeat number, the total nominal thickness of this sample was approximately 130 nm only. Therefore, after removing the deposit from the substrate, it was immediately appropriate for a TEM study in planar view. The structural features of the multilayer wCo95Cu5(3.6 nm)yCu(0.33 nm)xNs1500 were also of special interest since this multilayer had the smallest Cu spacer layer thickness at which still a clear GMR behaviour could be observed. The TEM investigations of the multilayer structure were carried out with a Philips CM20 equipment operated at 200 kV and using the Fresnel technique. A Philips CM20 type transmission electron microscope with field emission gun completed with energy disperse X-ray (EDX) analysis was used to measure the crosssectional local composition changes in the multilayers. The cross-sectional TEM samples were prepared by an ion-milling procedure developed by Barna w20x. 3. Results and discussion 3.1. Evolution of structure during multilayer growth A typical cross-sectional TEM picture of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer is shown in Fig. 1 that exhibits the characteristic microstructure of electrodeposited Co–CuyCu multilayers. The substrate side of the multilayer film is nearly smooth at this small magnification (it is a replica of the substrate surface itself), whereas at the solution side of the sample, a characteristic saw-tooth pattern with a roughness of approximately 1 mm can be identified. At a larger magnification (Fig. 2a), it is clearly visible that the substrate side of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer sample consists of very fine crystallites (marked with arrows). This indicates that the growth process starts with the deposition of fine crystallites whose diameter ranges from 20 to 50 nm. A TEM selected area diffraction (SAD) pattern (Fig. 2b) taken at the substrate side (marked area in Fig. 2a) also reveals that the deposit formed in the vicinity of the substrate consists of very small hcp and fcc crystals. As the thickness of the sample reaches approximately 100 nm, the grain size abruptly increases, and these more ordered large grains keep growing until the end of deposition.
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Fig. 1. Cross-sectional TEM picture of the electrodeposited wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer with the nominal thicknesses and bilayer repeat N as specified.
Fig. 3a shows a characteristic scanning TEM picture of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer near the substrate. It can be clearly seen that there are marked structural differences between the initial polycrystalline region deposited immediately on the substrate and the next region consisting of relatively large crystallites growing in a columnar form. A line-scan analysis along the work-line marked in Fig. 3a which is taken parallel to the substrate in the initial polycrystalline region reveals very large fluctuations in the Co and Cu concentrations (Fig. 3b). The chemical analysis indicates the frequent alternation of crystals with either very low or fairly high Cu-content. Although it is not directly evidenced from the data shown, one can assume that the crystals with a lower Cu concentration exhibit an hcp structure and those with a higher Cu-content form the fcc structure. This conclusion can be drawn from the equilibrium phase diagram of the Co–Cu system w21x: whereas hcp-Co cannot dissolve any Cu, in the high-temperature fcc-Co structure the solubility of Cu can be as high as 19.7 at.% (at 1367 8C). Taking into account that the thickness of the finegrained initial deposit is very small compared to the total thickness of the sample, the total amount of the hcp phase is estimated to be below 2%. This is in good agreement with the result of the earlier X-ray investigation w17x.
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The line-scan analyses performed parallel to the substrate in the large-grain region (at distances of ;0.15 and 0.3 mm from the substrate side, see Fig. 3c and d, respectively) detected much smoother concentration profiles. Comparing the results shown in Fig. 3c and d, we can see that the Cu-content increases with the growth of the film in the first 300 nm, reaching the average value measured with a large window-mode on the total cross-section of the thinned sample. The average value of the Cu concentration was found to be 42 at.% for this sample from the EDX analysis. Despite the fact that the same tendency was detected by the XRD method w17x, namely the lattice parameter measured at the solution side is larger, which means a larger Cu concentration than at the substrate side, the average Cu concentration measured by EDX is below the concentration value estimated from the lattice parameter resulting from Vegard’s law. This deviation in the concentration values obtained by two different methods shows clearly that in the case of a multilayered sample the application of the Vegard’s law overestimates the value obtained with a direct method. This is because the multilayered alloy is not a homogeneous solid solution but it has a larger amount of inner stress as a consequence of the interface lattice miss-match than the case of the replacement of Co atoms by Cu ones in a solid solution. To study the initial process taking place during the deposition, very thin samples have been produced which were transparent for the electron beam, and it was possible to investigate them in plane-view without any thinning. The selected wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs28 multilayer can be considered as a sample representing the very early stage of the deposition only. A typical plane-view TEM micrograph recorded for this sample is shown in Fig. 4a which reveals a fine polycrystalline microstructure. This suggests that the initial stage of deposition occurs via a nucleation and growth mechanism resulting in an island-like structure. The corresponding SAD pattern (Fig. 4b) can be indexed according to the same structures as in Fig. 2a, i.e. both fcc and hcp crystallites are present. Most of the small grains are characterised by a quasi-hexagonal shape (Fig. 4a), corresponding to the (1 1 1)fcc and (0 0 1)hcp orientation, in agreement with the preferred orientation detected by XRD w17x. A high density of Moire patterns is observable in the grains of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs28 multilayer on the higher-magnification TEM pictures (Fig. 4c). The parallel Moire fringes were found to exhibit two characteristic periodicities of 4.0 and 2.0 nm in the areas marked with A and B, respectively. By taking into account the preferred orientation relationship of the hcp and fcc crystallites, it can be concluded that they arise from the superposition of the hcp Co (1 0 0) lattice plane and the fcc (Co–Cu) (1 1 1) lattice plane in the
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Fig. 2. (a) The higher-magnification cross-sectional TEM picture taken on the same sample shown in Fig. 1 reveals the microstructural details at the substrate side; (b) the SAD pattern taken on the area marked in (a) reveals that the fine polycrystals have either an fcc or a hcp crystal structure of roughly equal amount.
case of A-type fringes, whereas from that of the hcp Co (1 0 1) lattice plane and the fcc (Co–Cu) (1 1 1) lattice plane in the case of B-type Moire fringes. With completely relaxed layers, the relationship 1ydMoires1ydhcpCoy1ydfcc(Co–Cu) would give Moire fringes of a period of dMoires4.8 nm for the A-type and 2.5 nm for the B-type. Here we used dhcpCo(1 0 0)s0.2165 nm, dhcpCo(1 0 1)s0.1910 nm for pure Co and dfcc(Co–Cu)(1 1 1)s0.2067 nm as obtained from the results of XRD measurements on the given sample. Similarly large deviation from the observed values were obtained when calculating the dMoire values from the pure fcc-Co(1 1 1) and fcc-Cu(1 1 1) metal lattice plane distances. These large differences in the observed and expected values of the Moire fringe periodicity indicate an extraordinarily high level of unrelaxed stresses in these small crystallites at the beginning of the deposition process that may effectively prevent the formation of large grains in the initial stage of multilayer deposition. In other words, the small size of the grains can be considered as arising mainly due to the high level of internal stresses.
3.2. Concentration inhomogeneities across the thickness of the multilayered samples Transmission line-scan EDX microanalysis revealed concentration fluctuations across the thickness of multilayers as shown in Fig. 5a and b where the result of a line-scan measurement taken on the wCo95Cu5(3.6 nm)y Cu(1.1 nm)xNs1325 sample can be seen. In agreement with the results obtained by the line-scan analysis taken parallel with the substrate (Fig. 3), these perpendicular line-scan analysis data also showed that the initial deposition starts with a higher Co:Cu ratio. For more bilayer repeats, the Cu concentration achieves the average value, but towards the solution side of the multilayer, the Cu concentration increases up to above 50% (Fig. 5b). Taking into account the increase of the surface roughness of the electrodeposited multilayer during deposition w17,19x, it can be understood that the exchange reaction (namely CoqCu2qsCuqCo2q) can take place with higher efficiency after the deposition of a large number of bilayers than at the beginning of the deposition. Hence, the Cu-content of the magnetic layer apparently increases as the deposition proceeds.
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Fig. 3. (a) Characteristic scanning TEM image taken near the substrate side of the wCo95 Cu5 (3.6 nm)yCu(1.1 nm)xNs1325 multilayer; (b) the trace of the line-scan analysis along the line drawn parallel to the substrate in the initial polycrystalline band as marked in (a); (c) and (d) present the results of two further line-scans taken parallel to the line in (a) at a distance of approximately 150 nm and 300 nm, respectively, from the substrate side surface.
High Co duty cycle (i.e. deposition with high tCo y (tCoqtCu) ratio) results in an increased migration in the neighbourhood of the cathode and, hence, it represents a higher contribution to the diffusion-controlled transport of the Cu2q species. This phenomenon also contributes to the concentration variation of Cu in the Co layer with the thickness of the Co layer. According to earlier observations w22x, Ni–Co–CuyCu multilayers with thicker Cu layers exhibit a smaller copper concentration in the ferromagnetic layers than the samples with thin Cu layers. The reason specified in Ref. w22x as being
responsible for this effect is the diffusion-controlled concentration shift of the electrolyte during deposition. We rather think that the Cu concentration dependence can be attributed to the change in migration with duty cycle. Most of the line-scans taken on this sample show another effect, namely, there is a ‘semi-macroscopic’ and a nearly periodic concentration fluctuation around the average value along the concentration profile recorded perpendicular to the substrate. The length of periodicity is approximately 3 mm. The origin of this
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Fig. 4. (a) Plane-view TEM image and (b) SAD pattern of a wCo95 Cu5 (3.6 nm)yCu(1.1 nm)xNs28 multilayer sample revealing the features of the initial stage of deposition (it is noted that no thinning was necessary for this sample and the diffraction pattern was identical with that in Fig. 2a); (c) higher-magnification TEM picture.
phenomenon could be ascribed to the concentration changes of the electrolyte near the growing surface during the deposition.
Another effect revealed by the TEM pictures of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer is that the thickness of the Co–Cu and Cu layers changes also
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Fig. 5. (a) Scanning TEM picture taken on the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer; (b) the EDX microanalysis results along the line marked in (a) from the top to the bottom across the total thickness of the multilayer.
Fig. 6. Cross-sectional TEM pictures of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer (a) at the bottom (substrate side) and (b) near to the top (solution side) of the multilayer.
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Fig. 7. Cross-sectional TEM picture of the wCo95Cu5(9.4 nm)yCu(1.1 nm)xNs1325 multilayer at the top of the sample where the canting of the layer planes is the most characteristic.
during the deposition. At the substrate side, the repeat period has nearly the same value as the nominal one, but towards the solution side the thickness of the Co– Cu layers decreases remarkably during growth (compare the layer thicknesses of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer shown in Fig. 6a and b), resulting in a higher average Cu concentration, in accordance with the observation of the EDX measurements. The change of the layer thicknesses (repeat period) is influenced by another phenomenon, namely, by the layer tilting that occurs as the layer growth proceeds. Due to the gradual development of this distortion, it is the most remarkable at the solution side of the samples (see the cross-sectional TEM picture of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer in Fig. 7) as already observed also for electrodeposited Ni81Cu19 yCu multilayers w23x. This change in the layer orientation causes a surface area enhancement of the growth front. Corre-
spondingly, the actual current density decreases to the same extent as the active cathode area increases. The diminished current density leads to a decrease in the thickness of the layers and hence it also results in a smaller-than-expected repeat period at the solution side of the sample. At the substrate side where the layers are nearly parallel to the substrate (Fig. 6a), the repeat period was found to be the same as the nominal one, but later it decreased remarkably (Fig. 6b and Fig. 7). This is the reason why the average repeat period detected by XRD w17x was smaller than that calculated with Faraday’s law (i.e. the nominal layer thickness). In some cases, this type of change in the repeat period across the thick multilayered film makes the detection of the layered structure by XRD impossible. The effect of the Cu spacer layer thickness on GMR was investigated previously w17x and it was found that the smallest nominal Cu layer thickness is 0.33 nm at
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Fig. 8. Cross-sectional TEM picture of the wCo95 Cu5 (3.6 nm)yCu(0.33 nm)xNs1500 multilayer which had the thinnest Cu spacer layer still exhibiting GMR.
which the sample still shows a GMR behaviour. For this reason, the structural study of the wCo95Cu5(3.6 nm)yCu(0.33 nm)xNs1500 multilayer was especially important. The TEM picture taken on this sample (Fig. 8) shows that the layered structure is not continuous everywhere but there are breaks marked by arrows. These breaks in the thinnest Cu spacer layer indicate the island-like nucleation and growth process during the deposition.
The inhomogeneities formed during deposition after some hundreds of repeat period could be made responsible for the decrease of GMR with increasing total thickness of the multilayered samples w17x. It has been established that decreasing the thickness of Cu spacer layers down to 0.33 nm (the sample with thinnest Cu spacer layers still exhibiting GMR), the layered structure is no more continuous everywhere, indicating an island-like nucleation and growth process during the deposition of each layer.
4. Conclusions Acknowledgments TEM measurements combined by EDX analysis revealed the structural details of multilayers, including the initial stage of deposition. It has been established that the deposition process starts with the formation of very small crystallites having either hcp or fcc crystal structure. These nanometer-sized crystallites have a high level of internal stress. As the deposition proceeds, the size of crystallites increases forming a characteristic columnar structure with an fcc superlattice. An EDX analysis showed that the average Cu concentration is lower near the substrate and increases during deposition. This increase is not monotonous across the whole sample thickness but it has a quasi ‘macroscopic’ fluctuation. During the growth process, the tilting of the layer planes becomes more remarkable, forming a characteristic saw-tooth shape at the top of the sample. This layer-bending causes a surface area enhancement of the growth front, decreasing the actual current density, resulting in the observed change of the repeat period.
This work was supported by the Hungarian Scientific Research Fund (OTKA) through grant F 032046. We have benefited from a joint research grant supported by the Hungarian–German (D-36y97) Intergovernmental Science and Technology Co-operation Programme. The allocation of the Philips CM20 equipment for the TEM studies at the Research Institute for Technical Physics and Materials Science of the Hungarian Academy of Sciences, Budapest, is also gratefully acknowledged. References w1x M.N. Baibich, J.M. Broto, A. Fert, F. Nguyen van Dau, F. Petroff, P. Etienne, G. Creuzet, A. Friederich, J. Chazelas, Phys. Rev. Lett. 61 (1988) 2472. w2x G. Binasch, P. Grunberg, ¨ F. Saurenbach, W. Zinn, Phys. Rev. B 39 (1989) 4828. w3x S.S.P. Parkin, R. Bhadra, K.P. Roche, Phys. Rev. Lett. 66 (1991) 2152.
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w4x S.S.P. Parkin, Z.G. Li, D.J. Smith, Appl. Phys. Lett. 58 (1991) 2710. w5x D.H. Mosca, F. Petroff, A. Fert, P.A. Schroeder, W.P. Pratt Jr., R. Laloee, J. Magn. Magn. Mater. 94 (1991) L1. w6x S.K. Lenczowski, C. Schonenberger, ¨ M.A.M. Gijs, W.J.M. de Jonge, J. Magn. Magn. Mater. 148 (1995) 455. w7x P. Nallet, E. Chassaing, M.G. Walls, M.J. Hytch, J. Appl. Phys. 79 (1996) 6884. w8x E. Chassaing, A. Morrone, J.E. Schmidt, J. Electrochem. Soc. 146 (1999) 1794. w9x Y. Jyoko, S. Kashiwabara, Y. Hayashi, J. Electrochem. Soc. 144 (1997) L5. w10x A. Dinia, K. Rahmouni, G. Schmerber, H. Bouanani, F. Cherkaoui, A. Berrada, Mater. Res. Soc. Symp. Proc. 475 (1997) 611. w11x H. El Fanity, K. Rahmouni, M. Bouanani, A. Dinia, G. ´ Cziraki, ´ ´ F. Cherkaoui, A. Schmerber, C. Meny, P. Panissod, A. Berrada, Thin Solid Films 318 (1998) 227. w12x Y. Ueda, N. Hataya, H. Zaman, J. Magn. Magn. Mater. 156 (1996) 350. w13x Y. Ueda, N. Kikuchi, S. Ikeda, T. Houga, J. Magn. Magn. Mater. 198–199 (1999) 740.
w14x M. Shima, L. Salamanca-Riba, T.P. Moffat, R.D. McMichael, L.J. Schwarzendruber, J. Appl. Phys. 84 (1998) 1504. w15x M. Shima, L. Salamanca-Riba, T.P. Moffat, R.D. McMichael, J. Magn. Magn. Mater. 198–199 (1999) 52. w16x K. Attenborough, H. Boere, J. De Boeck, G. Borghs, J.-P. Celis, Appl. Phys. Lett. 74 (1999) 2205. ´ Cziraki, w17x L. Peter, ´ ´ ´ A. L. Pogany, Z. Kupay, I. Bakonyi, M. Uhlemann, M. Herrich, B. Arnold, T. Bauer, K. Wetzig, J. Electrochem. Soc. 148 (2001) C168. w18x M. Shima, L.G. Salamanca-Riba, R.D. McMichael, T.P. Moffat, J. Electrochem. Soc. 148 (2001) C518. ´ Cziraki, w19x L. Peter, ´ ´ J. Padar, ´ ´ J. Toth, ´ I. Bakonyi, J. Z. Kupay, A. Phys. Chem. B 105 (2001) 10867. ´ Barna, Mater. Res. Soc. Symp. Proc. 254 (1991) 3. w20x A. w21x T.B. Massalski, Binary Alloy Phase Diagrams, ASM International, Materials Park, OH, USA, 1990. w22x S.Z. Hua, D.S. Lashmore, L. Salamanca-Riba, W. Schwarzacher, L.J. Swartzendruber, R.D. McMichael, L.H. Bennett, R. Hart, J. Appl. Phys. 76 (1994) 6519. ´ Cziraki, w23x A. ´ I. Gerocs, ´´ B. Fogarassy, B. Arnold, M. Reibold, K. ´ ´ ´ I. Bakonyi, Z. Metallkde. 88 (1997) Wetzig, E. Toth-Kadar, 781.
Structural evolution during growth of electrodeposited Co–CuyCu multilayers with giant magnetoresistance ´ Cziraki ´ a, L. Peter ´ b, B. Arnoldc, J. Thomasc, H.D. Bauerc, K. Wetzigc, I. Bakonyib,* A. a
¨ ¨ University, P.O. Box 32, H-1518 Budapest, Hungary Department of Solid State Physics, Eotvos Research Institute for Solid State Physics and Optics, Hungarian Academy of Sciences, P.O. Box 49, H-1525 Budapest, Hungary c ¨ Festkorperanalytik ¨ ¨ Festkorper-und ¨ Institut fur und Strukturforschung, Institut fur Werkstofforschung, Helmholtzstrasse 20, D-01069 Dresden, Germany b
Received 19 March 2002; received in revised form 11 October 2002; accepted 24 November 2002
Abstract The maximum room-temperature giant magnetoresistance (GMR) of electrodeposited Co–CuyCu multilayers produced during this work was approximately 9% at 8 kOe, and it was found to decrease with increasing bilayer repeat number. A transmission electron microscopy study has revealed the fine details of the microstructure formed during growth. At the beginning of the deposition very small, nano-sized crystallites formed with both hexagonal close-packed (hcp) and face-centred cubic (fcc) crystal structures containing a high level of internal stress. The Cu-content of these small crystallites was found to depend strongly on their crystal structure (fcc or hcp). After this initial polycrystalline region, the size of crystallites increases, forming an fcc superlattice with increasing average Cu concentration at the first hundreds of repeat periods. This increase is not monotonous across the whole sample thickness. As another effect, the bending of layer planes becomes more remarkable as the growth progresses. The above inhomogeneities formed during the deposition of hundreds of bilayers could be responsible for the decrease in GMR with increasing total thickness of the multilayered samples. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Multilayers; Structural properties; Transmission electron microscopy; Giant magnetoresistance
1. Introduction The discovery of giant magnetoresistance (GMR) observed in magnetic multilayers w1,2x has opened up a new material class for magnetic devices, such as magnetic read out heads, magnetic sensors and other information storage and retrieval devices. Exhibiting very large GMR (up to 50%) at room temperature w3–5x, the Co–Cu system is one of the most important among the taylored magneticynon-magnetic multilayers. A simple and cheap way to produce relatively thick self-supporting multilayered films in the Co–Cu system is electrodeposition. Nevertheless, as a result of the nature of the deposition method, the magnetic layers cannot consist of pure Co but rather a Co-rich Co–Cu alloy with *Corresponding author. Tel.: q36-1-392-2628; fax: q36-1-3922215. E-mail address: [email protected] (I. Bakonyi).
magnetic properties very similar to those of pure Co. The GMR values measured on these electrodeposited Co–CuyCu multilayered films have remained below 20% w6–19x, regardless of whether they were deposited with the control of potential, current or both alternating. In our recent studies of electrodeposited Co–CuyCu multilayers w17,19x, the aim was to clarify the deposition parameters that govern the GMR value. Direct-current plating, pulse-plating, two-pulse plating and reversepulse plating were used to produce relatively thick selfsupported electrodeposited Co–Cu alloys and Co–Cuy Cu multilayers under galvanostatic control from an electrolyte containing CoSO4 and CuSO4. It was found w17x that the maximum room-temperature GMR of these multilayer deposits was approximately 9% at 8 kOe and it decreased with increasing bilayer repeat number. This gradual deterioration of the GMR was ascribed to the increasing structural disorder developing with increasing
0040-6090/03/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 0 - 6 0 9 0 Ž 0 2 . 0 1 1 2 6 - 4
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total multilayer thickness (bilayer repeat number) as revealed from a preliminary structural investigation w17x by atomic force microscopy, X-ray diffraction (XRD) and transmission electron microscopy (TEM). According to these studies, the direct-current plating resulted in a Co95Cu5 alloy with nearly equal amounts of facecentred cubic (fcc) and hexagonal close-packed (hcp) phases, while all pulsed-current methods yielded Co– CuyCu multilayers which exhibited a dominant fcc structure and contained hcp crystals in an amount of less than 2% only. The appearance of satellite peaks on both sides of the (1 1 1) Bragg maximum in the large-angle XRD spectrum confirmed the formation of a multilayered structure for each pulsed electrodeposition mode. It was furthermore found that there was a strong (1 1 1) crystallographic orientation at the substrate side of the samples and a loss of the preferred orientation towards the solution side in correlation with the observed decrease in the GMR values. The loss in GMR was regarded as a consequence of the structural evolution as the specimen grew, but this assumption is to be verified by using a technique that can directly detect structural features. On the other hand, it is well known that the structural parameters measured with XRD can be regarded as mean values only. XRD has a very limited ability to detect local concentration inhomogeneities and is not sensitive enough to reveal the local lattice distortions, like stacking faults, twin boundaries and bending of the layers. Therefore, in order to better reveal the nature of the microstructure forming in these multilayers, and especially their evolution during growth, a more detailed TEM study has now been performed. This is a unique direct method to visualise the microstructure at the atomic level. The results of these TEM investigations including elemental analysis on the nanometer range will be presented in this paper in order to complete our previous findings w17x mainly on the same electrodeposited Co–CuyCu multilayer samples. 2. Experimental Two-pulse plating from a single electrolyte containing CoSO4 and CuSO4 was used to produce electrodeposited Co–CuyCu multilayers under galvanostatic control on a polycrystalline Ti substrate as described in more detail in Ref. w17x. The deposits were peeled off from the substrates by mechanical stripping. The present detailed TEM studies were performed on electrodeposited multilayers of the form wCo95Cu5(dCo–Cu)yCu(dCu)xN where the nominal magnetic (dCo–Cu) and non-magnetic (dCu) layer thicknesses were in the ranges dCo–Cus3.6–12 nm and dCus0.22– 1.8 nm and the bilayer repeat number N varied from 28 to 1500. Most of the multilayer samples were identical
with those of Ref. w17x. Among them, the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer had the largest room-temperature GMR of approximately 9% at 8 kOe and, for this reason, TEM results on this sample will be presented in more detail. For a better understanding of the initial stage of deposition, a wCo95Cu5(3.6 nm)y Cu(1.1 nm)xNs28 multilayer was also prepared under identical conditions for the present TEM studies. Due to the small bilayer repeat number, the total nominal thickness of this sample was approximately 130 nm only. Therefore, after removing the deposit from the substrate, it was immediately appropriate for a TEM study in planar view. The structural features of the multilayer wCo95Cu5(3.6 nm)yCu(0.33 nm)xNs1500 were also of special interest since this multilayer had the smallest Cu spacer layer thickness at which still a clear GMR behaviour could be observed. The TEM investigations of the multilayer structure were carried out with a Philips CM20 equipment operated at 200 kV and using the Fresnel technique. A Philips CM20 type transmission electron microscope with field emission gun completed with energy disperse X-ray (EDX) analysis was used to measure the crosssectional local composition changes in the multilayers. The cross-sectional TEM samples were prepared by an ion-milling procedure developed by Barna w20x. 3. Results and discussion 3.1. Evolution of structure during multilayer growth A typical cross-sectional TEM picture of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer is shown in Fig. 1 that exhibits the characteristic microstructure of electrodeposited Co–CuyCu multilayers. The substrate side of the multilayer film is nearly smooth at this small magnification (it is a replica of the substrate surface itself), whereas at the solution side of the sample, a characteristic saw-tooth pattern with a roughness of approximately 1 mm can be identified. At a larger magnification (Fig. 2a), it is clearly visible that the substrate side of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer sample consists of very fine crystallites (marked with arrows). This indicates that the growth process starts with the deposition of fine crystallites whose diameter ranges from 20 to 50 nm. A TEM selected area diffraction (SAD) pattern (Fig. 2b) taken at the substrate side (marked area in Fig. 2a) also reveals that the deposit formed in the vicinity of the substrate consists of very small hcp and fcc crystals. As the thickness of the sample reaches approximately 100 nm, the grain size abruptly increases, and these more ordered large grains keep growing until the end of deposition.
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Fig. 1. Cross-sectional TEM picture of the electrodeposited wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer with the nominal thicknesses and bilayer repeat N as specified.
Fig. 3a shows a characteristic scanning TEM picture of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer near the substrate. It can be clearly seen that there are marked structural differences between the initial polycrystalline region deposited immediately on the substrate and the next region consisting of relatively large crystallites growing in a columnar form. A line-scan analysis along the work-line marked in Fig. 3a which is taken parallel to the substrate in the initial polycrystalline region reveals very large fluctuations in the Co and Cu concentrations (Fig. 3b). The chemical analysis indicates the frequent alternation of crystals with either very low or fairly high Cu-content. Although it is not directly evidenced from the data shown, one can assume that the crystals with a lower Cu concentration exhibit an hcp structure and those with a higher Cu-content form the fcc structure. This conclusion can be drawn from the equilibrium phase diagram of the Co–Cu system w21x: whereas hcp-Co cannot dissolve any Cu, in the high-temperature fcc-Co structure the solubility of Cu can be as high as 19.7 at.% (at 1367 8C). Taking into account that the thickness of the finegrained initial deposit is very small compared to the total thickness of the sample, the total amount of the hcp phase is estimated to be below 2%. This is in good agreement with the result of the earlier X-ray investigation w17x.
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The line-scan analyses performed parallel to the substrate in the large-grain region (at distances of ;0.15 and 0.3 mm from the substrate side, see Fig. 3c and d, respectively) detected much smoother concentration profiles. Comparing the results shown in Fig. 3c and d, we can see that the Cu-content increases with the growth of the film in the first 300 nm, reaching the average value measured with a large window-mode on the total cross-section of the thinned sample. The average value of the Cu concentration was found to be 42 at.% for this sample from the EDX analysis. Despite the fact that the same tendency was detected by the XRD method w17x, namely the lattice parameter measured at the solution side is larger, which means a larger Cu concentration than at the substrate side, the average Cu concentration measured by EDX is below the concentration value estimated from the lattice parameter resulting from Vegard’s law. This deviation in the concentration values obtained by two different methods shows clearly that in the case of a multilayered sample the application of the Vegard’s law overestimates the value obtained with a direct method. This is because the multilayered alloy is not a homogeneous solid solution but it has a larger amount of inner stress as a consequence of the interface lattice miss-match than the case of the replacement of Co atoms by Cu ones in a solid solution. To study the initial process taking place during the deposition, very thin samples have been produced which were transparent for the electron beam, and it was possible to investigate them in plane-view without any thinning. The selected wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs28 multilayer can be considered as a sample representing the very early stage of the deposition only. A typical plane-view TEM micrograph recorded for this sample is shown in Fig. 4a which reveals a fine polycrystalline microstructure. This suggests that the initial stage of deposition occurs via a nucleation and growth mechanism resulting in an island-like structure. The corresponding SAD pattern (Fig. 4b) can be indexed according to the same structures as in Fig. 2a, i.e. both fcc and hcp crystallites are present. Most of the small grains are characterised by a quasi-hexagonal shape (Fig. 4a), corresponding to the (1 1 1)fcc and (0 0 1)hcp orientation, in agreement with the preferred orientation detected by XRD w17x. A high density of Moire patterns is observable in the grains of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs28 multilayer on the higher-magnification TEM pictures (Fig. 4c). The parallel Moire fringes were found to exhibit two characteristic periodicities of 4.0 and 2.0 nm in the areas marked with A and B, respectively. By taking into account the preferred orientation relationship of the hcp and fcc crystallites, it can be concluded that they arise from the superposition of the hcp Co (1 0 0) lattice plane and the fcc (Co–Cu) (1 1 1) lattice plane in the
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Fig. 2. (a) The higher-magnification cross-sectional TEM picture taken on the same sample shown in Fig. 1 reveals the microstructural details at the substrate side; (b) the SAD pattern taken on the area marked in (a) reveals that the fine polycrystals have either an fcc or a hcp crystal structure of roughly equal amount.
case of A-type fringes, whereas from that of the hcp Co (1 0 1) lattice plane and the fcc (Co–Cu) (1 1 1) lattice plane in the case of B-type Moire fringes. With completely relaxed layers, the relationship 1ydMoires1ydhcpCoy1ydfcc(Co–Cu) would give Moire fringes of a period of dMoires4.8 nm for the A-type and 2.5 nm for the B-type. Here we used dhcpCo(1 0 0)s0.2165 nm, dhcpCo(1 0 1)s0.1910 nm for pure Co and dfcc(Co–Cu)(1 1 1)s0.2067 nm as obtained from the results of XRD measurements on the given sample. Similarly large deviation from the observed values were obtained when calculating the dMoire values from the pure fcc-Co(1 1 1) and fcc-Cu(1 1 1) metal lattice plane distances. These large differences in the observed and expected values of the Moire fringe periodicity indicate an extraordinarily high level of unrelaxed stresses in these small crystallites at the beginning of the deposition process that may effectively prevent the formation of large grains in the initial stage of multilayer deposition. In other words, the small size of the grains can be considered as arising mainly due to the high level of internal stresses.
3.2. Concentration inhomogeneities across the thickness of the multilayered samples Transmission line-scan EDX microanalysis revealed concentration fluctuations across the thickness of multilayers as shown in Fig. 5a and b where the result of a line-scan measurement taken on the wCo95Cu5(3.6 nm)y Cu(1.1 nm)xNs1325 sample can be seen. In agreement with the results obtained by the line-scan analysis taken parallel with the substrate (Fig. 3), these perpendicular line-scan analysis data also showed that the initial deposition starts with a higher Co:Cu ratio. For more bilayer repeats, the Cu concentration achieves the average value, but towards the solution side of the multilayer, the Cu concentration increases up to above 50% (Fig. 5b). Taking into account the increase of the surface roughness of the electrodeposited multilayer during deposition w17,19x, it can be understood that the exchange reaction (namely CoqCu2qsCuqCo2q) can take place with higher efficiency after the deposition of a large number of bilayers than at the beginning of the deposition. Hence, the Cu-content of the magnetic layer apparently increases as the deposition proceeds.
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Fig. 3. (a) Characteristic scanning TEM image taken near the substrate side of the wCo95 Cu5 (3.6 nm)yCu(1.1 nm)xNs1325 multilayer; (b) the trace of the line-scan analysis along the line drawn parallel to the substrate in the initial polycrystalline band as marked in (a); (c) and (d) present the results of two further line-scans taken parallel to the line in (a) at a distance of approximately 150 nm and 300 nm, respectively, from the substrate side surface.
High Co duty cycle (i.e. deposition with high tCo y (tCoqtCu) ratio) results in an increased migration in the neighbourhood of the cathode and, hence, it represents a higher contribution to the diffusion-controlled transport of the Cu2q species. This phenomenon also contributes to the concentration variation of Cu in the Co layer with the thickness of the Co layer. According to earlier observations w22x, Ni–Co–CuyCu multilayers with thicker Cu layers exhibit a smaller copper concentration in the ferromagnetic layers than the samples with thin Cu layers. The reason specified in Ref. w22x as being
responsible for this effect is the diffusion-controlled concentration shift of the electrolyte during deposition. We rather think that the Cu concentration dependence can be attributed to the change in migration with duty cycle. Most of the line-scans taken on this sample show another effect, namely, there is a ‘semi-macroscopic’ and a nearly periodic concentration fluctuation around the average value along the concentration profile recorded perpendicular to the substrate. The length of periodicity is approximately 3 mm. The origin of this
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Fig. 4. (a) Plane-view TEM image and (b) SAD pattern of a wCo95 Cu5 (3.6 nm)yCu(1.1 nm)xNs28 multilayer sample revealing the features of the initial stage of deposition (it is noted that no thinning was necessary for this sample and the diffraction pattern was identical with that in Fig. 2a); (c) higher-magnification TEM picture.
phenomenon could be ascribed to the concentration changes of the electrolyte near the growing surface during the deposition.
Another effect revealed by the TEM pictures of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer is that the thickness of the Co–Cu and Cu layers changes also
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Fig. 5. (a) Scanning TEM picture taken on the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer; (b) the EDX microanalysis results along the line marked in (a) from the top to the bottom across the total thickness of the multilayer.
Fig. 6. Cross-sectional TEM pictures of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer (a) at the bottom (substrate side) and (b) near to the top (solution side) of the multilayer.
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Fig. 7. Cross-sectional TEM picture of the wCo95Cu5(9.4 nm)yCu(1.1 nm)xNs1325 multilayer at the top of the sample where the canting of the layer planes is the most characteristic.
during the deposition. At the substrate side, the repeat period has nearly the same value as the nominal one, but towards the solution side the thickness of the Co– Cu layers decreases remarkably during growth (compare the layer thicknesses of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer shown in Fig. 6a and b), resulting in a higher average Cu concentration, in accordance with the observation of the EDX measurements. The change of the layer thicknesses (repeat period) is influenced by another phenomenon, namely, by the layer tilting that occurs as the layer growth proceeds. Due to the gradual development of this distortion, it is the most remarkable at the solution side of the samples (see the cross-sectional TEM picture of the wCo95Cu5(3.6 nm)yCu(1.1 nm)xNs1325 multilayer in Fig. 7) as already observed also for electrodeposited Ni81Cu19 yCu multilayers w23x. This change in the layer orientation causes a surface area enhancement of the growth front. Corre-
spondingly, the actual current density decreases to the same extent as the active cathode area increases. The diminished current density leads to a decrease in the thickness of the layers and hence it also results in a smaller-than-expected repeat period at the solution side of the sample. At the substrate side where the layers are nearly parallel to the substrate (Fig. 6a), the repeat period was found to be the same as the nominal one, but later it decreased remarkably (Fig. 6b and Fig. 7). This is the reason why the average repeat period detected by XRD w17x was smaller than that calculated with Faraday’s law (i.e. the nominal layer thickness). In some cases, this type of change in the repeat period across the thick multilayered film makes the detection of the layered structure by XRD impossible. The effect of the Cu spacer layer thickness on GMR was investigated previously w17x and it was found that the smallest nominal Cu layer thickness is 0.33 nm at
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Fig. 8. Cross-sectional TEM picture of the wCo95 Cu5 (3.6 nm)yCu(0.33 nm)xNs1500 multilayer which had the thinnest Cu spacer layer still exhibiting GMR.
which the sample still shows a GMR behaviour. For this reason, the structural study of the wCo95Cu5(3.6 nm)yCu(0.33 nm)xNs1500 multilayer was especially important. The TEM picture taken on this sample (Fig. 8) shows that the layered structure is not continuous everywhere but there are breaks marked by arrows. These breaks in the thinnest Cu spacer layer indicate the island-like nucleation and growth process during the deposition.
The inhomogeneities formed during deposition after some hundreds of repeat period could be made responsible for the decrease of GMR with increasing total thickness of the multilayered samples w17x. It has been established that decreasing the thickness of Cu spacer layers down to 0.33 nm (the sample with thinnest Cu spacer layers still exhibiting GMR), the layered structure is no more continuous everywhere, indicating an island-like nucleation and growth process during the deposition of each layer.
4. Conclusions Acknowledgments TEM measurements combined by EDX analysis revealed the structural details of multilayers, including the initial stage of deposition. It has been established that the deposition process starts with the formation of very small crystallites having either hcp or fcc crystal structure. These nanometer-sized crystallites have a high level of internal stress. As the deposition proceeds, the size of crystallites increases forming a characteristic columnar structure with an fcc superlattice. An EDX analysis showed that the average Cu concentration is lower near the substrate and increases during deposition. This increase is not monotonous across the whole sample thickness but it has a quasi ‘macroscopic’ fluctuation. During the growth process, the tilting of the layer planes becomes more remarkable, forming a characteristic saw-tooth shape at the top of the sample. This layer-bending causes a surface area enhancement of the growth front, decreasing the actual current density, resulting in the observed change of the repeat period.
This work was supported by the Hungarian Scientific Research Fund (OTKA) through grant F 032046. We have benefited from a joint research grant supported by the Hungarian–German (D-36y97) Intergovernmental Science and Technology Co-operation Programme. The allocation of the Philips CM20 equipment for the TEM studies at the Research Institute for Technical Physics and Materials Science of the Hungarian Academy of Sciences, Budapest, is also gratefully acknowledged. References w1x M.N. Baibich, J.M. Broto, A. Fert, F. Nguyen van Dau, F. Petroff, P. Etienne, G. Creuzet, A. Friederich, J. Chazelas, Phys. Rev. Lett. 61 (1988) 2472. w2x G. Binasch, P. Grunberg, ¨ F. Saurenbach, W. Zinn, Phys. Rev. B 39 (1989) 4828. w3x S.S.P. Parkin, R. Bhadra, K.P. Roche, Phys. Rev. Lett. 66 (1991) 2152.
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