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Ag multilayers submitted to thermal annealing
Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
Structural evolution and magnetic behavior of Co/Ag multilayers submitted to thermal annealing W.H. Flores!, S.R. Teixeira!,*, J. Geshev!, J.B.M. da Cunha!, P.J. Schilling", A. Traverse#, M.C. Martins Alves$ ! Instituto de Fı& sica-UFRGS, C.P. 15051, 91501-970 Porto Alegre, Brazil " Center for Advanced Microstructures and Devices, Baton Rouge, USA # LURE-Universite& Paris-Sud, 91405 Orsay Cedex, France $ Laborato& rio Nacional de Luz Sı& ncrotron, LNLS, C.P. 6192, 13081-970 Campinas, Brazil Received 13 September 1997; received in revised form 13 March 1998
Abstract We report on the structural evolution and the magnetic behavior of the Co/Ag multilayered films deposited in a UHV chamber at room temperature, submitted to 10 min thermal annealing at temperatures ranging from 100 to 600°C. The structural characterization was performed using X-ray diffraction and X-ray absorption spectroscopy techniques. Magnetoresistance and magnetization measurements were used to study the evolution and magnetic behavior of the samples. The results show that, besides the roughness at the interfaces and the structural disorder of the Co layers, the as-deposited sample has a compressive stress at Ag—Co interface originated by the difference between the surface energies of Ag and Co. After annealing at 400°C, there is a breakup of the layers accompanied by a relaxation of the stress and defects as well as partial crystallographic ordering of the Co clusters. The room-temperature magnetoresistance change from anisotropic magnetoresistance to giant magnetoresistance with a sharp maximum, reaching 4.5%. ( 1998 Elsevier Science B.V. All rights reserved. PACS: 74.25.Ha; 75.70.i; 75.70.Pa Keywords: Magnetic properties; Thin films; Multilayers
1. Introduction The observation of the giant magnetoresistance (GMR) effect in granular films [1,2], where the
* Corresponding author. Tel.: #55 51 316 6498/7111; fax: #55 51 319 1762; e-mail: [email protected].
magnetic material is formed by small precipitates embedded in a non-magnetic matrix, has stimulated remarkable interest because of its potential technological applications as magnetic sensors in the magnetic storage technology [3,4], and also by its interest in fundamental scientific research. The granular films can present a GMR effect even larger than in multilayers. They are often obtained by
0304-8853/98/$19.00 ( 1998 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 9 8 ) 0 0 1 4 2 - 5
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co-sputtering of two immiscible metallic elements [5], or by sputtering of metastable alloys [6]. The size of the magnetic particles is of great interest, as it determines to a large extent, the magnetic behavior of the granular alloy. Its control is dictated by substrate temperature during the deposition [7], or after deposition by subsequent annealing [6], until the optimum annealing temperature that maximizes the GMR effect is obtained. However, for granular alloy systems, the particle size, shape and temperature dominate the field dependence of the GMR effect, making difficult the use of these alloys in low-field sensitivity devices, because of the large distribution of particle sizes and shapes. It is also impossible to predict the optimum temperature because it depends on the deposition methods used to produce the alloys as well as on the sample components. Another important parameter that should be controlled during the deposition process of granular films is the concentration of the magnetic metal [8—10], which has an important influence on the GMR effect. At high magnetic metal concentration, the magnetic phase percolates throughout the material and the sample appears to be a ferromagnet. Recently, Hylton et al. [11] have proposed the preparation of discontinuous multilayers by annealing the sputtered multilayers in order to minimize the effects of crystal and shape anisotropy. With this procedure, it is possible to control the composition and to minimize the effects of the shape and the size of the magnetic particles on the GMR. In fact, the discontinuous multilayers will be composed of an island-like network of multidomain magnetic particles, within the magnetic layers separated by a continuous non-magnetic spacer. The interfaces between the ferromagnetic and non-magnetic layers are not necessarily flat and continuous, they have a roughness and contain some discontinuities. The ferromagnetic layers form a network consisting of connected ferromagnetic grains. Hence, a multidomain magnetic state within each layer is expected. The aim here is to achieve a complete insulation of the individual magnetic grains through successive annealing within a certain ordered layer structure. The role of annealing on the magnetic properties of the Co/Ag systems, has been investigated re-
cently [12—19]. In particular, the magnetic interactions between magnetic entities with two distinct magnetic phases, and the correlation of the structural parameters obtained by a detailed analysis of the hysteresis loops, have been described and discussed [15—18]. In this context, the presence of superparamagnetic and ferromagnetic particles in this system with inter-particle coupling leads to a complex behavior. Thus, the magnetic properties are dependent on the topological structure of the ferromagnetic component, especially on their sizes and structure. This paper adopts the same procedure suggested by Hylton et al. [11] to obtain a granular system. Starting from a dual electron beam deposited Co/Ag multilayer structure with thin magnetic layers, a breakup of the layers through successive annealing is performed to obtain a regular granular magnetic network embedded in a non-magnetic matrix. The main interest here, is to correlate the structural evolution with the magnetotransport and macroscopic magnetic behavior of the sample after each thermal treatment, to understand the drastic changes in the magnetoresistance effect of this system.
2. Experimental The multilayers [Ag(50 A_ )/Co(15 A_ )] , were 10 prepared at room temperature by alternate deposition in ultra-high vacuum on a Si(1 1 1) substrate covered with a thick SiO layer. The depositions 2 were performed using a dual electron beam Balzers evaporator. The base pressure in the evaporation chamber was 1.0]10~8 mbar. All samples were prepared with a deposition rate of 1 A_ /s, monitored by a quartz balance. The total thickness of the samples was 65 nm. The thickness of the Ag and Co layers were chosen to give an overall concentration of the approximately 20 at% for the ferromagnetic metal. The multilayers were annealed in a high vacuum furnace (10~7 mbar) for 10 min at temperatures ranging from 100 to 600°C. After each anneal, the samples were cooled down to room temperature for subsequent analyses. The structural characterization of the samples were made using X-ray diffraction (XRD) and
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
X-ray absorption spectroscopy (XAS), with X-ray absorption near edge (XANES) and extended X-ray absorption fine structure (EXAFS) modes. The XRD was performed in h—2h Bragg—Brentano geometry at room-temperature employing Cu Ka radiation. XAS experiments at the Co K-edge were performed on line D42 at the DCI storage ring, LURE-Orsay, running at an energy of 1.85 GeV and a current of 300 mA, and at LSU Center for Advanced Microstructures and Devices (CAMD). At the D42 line, the X-rays were monochromatized with a Si(3 3 1) channel cut single crystal with an energy resolution of 1 eV at the Co K-edge. The incident beam was collimated by slits and its intensity was measured by an ionization chamber. XANES and EXAFS spectra of the samples and model compounds were recorded in the total electron yield [20] due to the presence of the SiO 2 substrate on the multilayers. The XANES data were collected in scans with an energy step of 0.5 eV and two scans were added to obtain a good signalto-noise ratio. At CAMD, the synchrotron X-ray beam was monochromatized using the LNLS double-crystal monochromator equipped with Si(1 1 1) crystals [21]. All data were collected in total electron yield (TEY) at room temperature [22]. The EXAFS spectra were analyzed using standard methods [23]. It involves a background subtraction by means of a polynomial function normalized to the height of the absorption edge step. The various neighboring shells were obtained by a Fourier transform of the EXAFS signal from 3 to 12.5 A_ using a Kaiser window with a coefficient of 2.5. By an inverse Fourier transform into k space, the EXAFS oscillations corresponding to only one neighbor shell were obtained. Structural parameters are obtained from least squares fitting using experimental phase and amplitude functions deduced from a model compound. We have used phase shift and amplitudes extracted from a Co/SiO /Si (1 1 1) sample annealed at 600°C. 2 The magnetic properties were investigated at room temperature using a conventional vibrating sample magnetometer (VSM) and alternate gradient magnetometer (AGM), with the applied field always parallel to the film surface. The mag-
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netoresistence (MR) was measured at room temperature using a standard four-point configuration with in-plane current between 0.1 and 1 mA. These measurements were performed with the magnetic field parallel and perpendicular to the current. The fractional variation of MR, *R/R, was taken as R(H)!R(H"0)/R(H"0), where R(H"0) is the value in the initial, H"0, state; here R is the resistance and H is the applied magnetic field.
3. Results and discussion 3.1. X-ray The XRD patterns in a Bragg—Brentano geometry for the as-deposited and annealed samples are shown in Fig. 1. The as-deposited sample, Fig. 1a, exhibits peaks attributable to Ag(1 1 1), Ag(2 0 0), Ag(2 2 0), Ag(3 1 1), and Ag(2 2 2) reflections, corresponding to FCC polycrystalline structure. The samples show a preferential orientation in the (1 1 1) direction. As can be seen, the relative intensities of the high index lines compared to the (1 1 1) line, are much lower than the relative intensities of the high index lines also compared to the (1 1 1) line of a powder spectrum, even considering that in the present case, the spectra show an overlap between the Co(1 1 1) and Ag(2 0 0) lines. Due to the roughness of the interfaces, no satellites are observed indicating that the multilayer is most probably discontinuous. The Co(1 1 1) (FCC) reflection cannot be distinguished because it coincides with the Ag(2 0 0) reflection. It is important to note that for the as-deposited sample, no other reflections corresponding to Co are visible, indicating that the Co layers have a very disordered structure. The Ag(1 1 1) reflection, as can be observed from Fig. 1a, is split into two overlapping lines, 2h" 37.85 and 38.57°, differing from the samples obtained by sputtering of metastable Co—Ag alloys [6—24], or co-sputtering of Ag and Co [5,8]. The two split lines can be identified by the shoulder at the low-angle part and by a more intense peak at the right side of the whole Ag(1 1 1) reflection. To see if the origin of the split line was under stress at the Ag/Co interfaces, an u tilt measurement for the
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Fig. 2. Stress measurements, u tilt, of the Ag(1 1 1) planes. Plot of d versus sin2 u. 111
Fig. 1. The XDR patterns for Co/Ag: (a) as-deposited sample, (b) and (c) samples annealed at 400 and 600°C, respectively.
two peaks was performed. The results displayed in Fig. 2 as d versus sin2 u, clearly show a com111 pressive stress of the Ag(1 1 1) split line. Using the well known expression p"aE/(1#l), where a is the slope of the fitted straight line, E the Young’s modulus, and l the Poisson ratio (for Ag, E"7.58]1010 N/m2 and l"0.38), the values obtained for the compressive stress were 23 and 5.5 GPa for the low- and high-angle peaks, respectively. The stress comes from the Co/Ag interfaces, probably due to the differences of the surface energies of Ag and Co (c "1.3 J/m2 and A' c "2.71 J/m2 [25,26]). When Co is deposited on C0 a Ag layer, there will be a formation of small Co islands leading to a stressed Co/Ag interface, because Co does not wet the Ag film. This stressed region corresponds to the shoulder that appear at the low-angle peak region of the Ag(1 1 1) reflec-
tion. On the other hand, as the thickness of the Ag layer increases, the stress decreases and this less stressed region corresponds to the high-angle peak. The Ag wets the Co surface more, so Ag/Co interface is expected to be weakly stressed. The oscillations occurring in d values around the fitting 111 straight line that appear in Fig. 2 are due to the texture effects in the Ag layers [27]. The high-angle peak at Ag(1 1 1) is slightly shifted to the right from the expected position of the bulk Ag(1 1 1) peak as is indicated in Fig. 1a. This can be due to a small incorporation of Co atoms forming a solid solution in Ag. Using the Veggard’s law [28], it was estimated that 7.5% of Co are substitutional in the Ag lattice near Co/Ag interfaces. Annealing at temperatures below 400°C, for 10 min, does not significantly change the XRD patterns, which are very similar to the ones for the as-deposited sample. After annealing at 400°C, for the same time (Fig. 1b) the Ag peaks narrow considerably, and although very small, a second peak begins to emerge at 2h around 52°, corresponding to Co(2 0 0) (FCC), as indicated by the arrow in the figure. For this annealing temperature there is a narrowing of the diffraction linewidths as a consequence of an ordering of the crystalline structure due to the release of defects and stresses. The shoulder at the left side of the Ag(1 1 1) peak almost disappears. Despite the XRD results showing a polycrystalline spectra, the Ag becomes highly S1 1 1T-textured as the annealing temperature
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
increases. This behavior can be observed by the substantial increase of the Ag(1 1 1) and (2 2 2) peaks as compared with the intensities of the Ag(2 2 0) and (3 1 1) reflections. Although bulk Co (HCP) is stable below the allotropic transformation temperature of about 420°C, there are no evidences of Co (HCP) phase in our samples, as FCC phase fine particles are usually stable at ambient temperature (see, e.g. Sato et al. [29] and the references therein). The ordering temperature of the Co cannot be precisely estimated due to the overlap of the Ag(2 0 0) and Co(1 1 1) peaks. Annealing at 600°C, for 10 min, causes narrowing of the peaks and growing up of Co(2 2 0) (FCC) reflection. The figure clearly shows Co (FCC) clusters embedded in a Ag(1 1 1) textured matrix. 3.2. XAS Fig. 3 shows the Co K-edge XANES data recorded for Co/Ag samples as a function of the annealing temperature. The XANES data of a bulk Co are given for comparison. The reference energy value (0 eV) corresponds to the first inflection point of the metallic Co edge (7709 eV).
Fig. 3. The XANES spectra from the Co edges for Co/Ag asdeposited, and for some annealed samples. The spectra of Co reference and simulations are shown for comparison.
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In pure Co, peaks labeled 1, 2 and 3 arise from multiple scattering effects of the photoelectron by the different shells being characteristic of the HCP symmetry. Simulations using a FEFF code [30,31] for Co (FCC) and (HCP) allow us to differentiate the two Co structures, as can be seen in Fig. 3. By comparison with XRD spectra we conclude that only Co (FCC) is present. The XANES profiles show an evolution with a clear difference between the features observed for the as-deposited sample and the one annealed at 200°C, which are very similar, and those annealed at 400 and 600°C. As can be seen from the figure, there is a decrease in intensity of the peak 1 while an increase in intensity is observed for the peaks 2 and 3, with a visible sharpening of the peak 3. This means that a structural modification had occurred around Co after annealing at 400°C in agreement with what is observed from XRD. This evolution comes from the release of defects and stresses, as observed in the X-ray spectra (Fig. 1), originating the ordering of the Co and Ag. It is interesting to note that the XANES spectra are quite different from that of the reference (Co foil). Fig. 4a shows the EXAFS signal for the as-deposited and annealed samples. Co foil data are included for reference as well. The spectra of the as-deposited and the annealed samples are quite different from the Co foil ones. As the annealing temperature increases the intensity increases reaching values even higher than that of the reference. The small oscillations appearing at the right hand side of the first EXAFS oscillation are due to the existence of small Co clusters. Details about this fine structure will be published elsewhere. Fig. 4b presents the k3 weighted Fourier transform (FT) of the Co K-edge EXAFS signals depicted in Fig. 4a. The figure also shows the simulation using FEFF code for Co (HCP) and (FCC). Comparing the FT obtained experimentally with the simulation we see that the Co present a FCC structure, in agreement with that observed by XRD. This difference is evidenced by comparing the intensities of the peaks 3 and 4 with the simulated spectrum corresponding to the Co (FCC), which are quite different from the FT obtained for the Co foil. The FT of Co metal exhibits four peaks located at 2.2, 3.4, 4.1 and 4.7 A_ (uncorrected from the phase shift). The first peak is
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W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
Fig. 5. The inverse Fourier transform of the first shell: open circles are the experimental data and full line, fitted curve.
Fig. 4. (a) The EXAFS measurements; (b) Fourier transforms of signals from the Co K-edge for Co/Ag as-deposited, and for some annealed samples. *k"3—12.5 A_ , Kaiser window q"2.5. The spectra of Co reference and simulations are shown for comparison.
related to the Co—Co near neighbor (NN) distance and the others are related to more distant neighbors. In contrast to XRD measurements, there are no evidences of NN Co—Ag contribution, as was
reported by Regnard et al. [14,19] for the sputter deposited AgCo films. To clarify this differences, low temperature EXAFS measurements will be performed in near future. The amplitudes of the FT peaks for the as-deposited and annealed at 200°C samples are small compared with the Co reference. This is an evidence for a reduction in the neighbor coordination for all the shells. This effect disappears for higher annealing temperatures where the intensity of the peaks is close to that of the Co foil, indicating an ordering of the Co structure in agreement with what was observed by XRD. In order to obtain quantitative estimations of the bond lengths and the coordination numbers around the cobalt atoms, the inverse Fourier transforms of the first shell have been fitted. A typical fitting is shown in Fig. 5. The backscattering amplitude and the phase shift values were deduced from the EXAFS spectrum of a Co/SiO /Si(1 1 1) sample 2 annealed at 600°C. The fit parameters are given in Table 1. The Fourier filtered spectra were fitted with a single shell of neighboring atoms, as only Co atoms are present in the first shell. The curve fitting indicates a reduction in NN coordination number, for the as-deposited and annealed samples at 200 and 400°C, see Table 1. For the sample treated at 600°C this number is 12 as for the bulk, indicating an ordered structure of the Co clusters. The small coordination number for as-deposited sample indicates a disordered structure of the Co, as observed
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
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Table 1 Parameters determined from magnetic measurement and by fitting the first shell in the Fourier-filtered EXAFS spectra for Co/Ag as a function of the annealing temperature, ¹ . The phase shift and amplitudes were extracted from the sample Co/SiO /Si(1 1 1) treated at A 2 600°C. M , saturation magnetization in units of emu/cm3; D, diameter of the particles in A_ ; N, coordination number; R, distances of 4 nearest neighbors in A_ , *p, Debye—Waller factor relative to the model compound in (A_ 2) Magn.
EXAFS
¹ A
M 4
D
N
R
*p2
Fit quality (]10~4)
As-dep. 200°C 400°C 600°C Co ref.
1480.64 1477.48 1383.53 1413.28 1466.72
— — 48.9 42.2 —
7.3 8 10.1 12 12
2.52 2.52 2.51 2.52 2.50
1.3]10~6 2.9]10~6 1.7]10~5 1.0]10~7 —
19 8.3 7.1 0.2 —
by XRD measurements. After annealing at 200 and 400°C, an increase of the coordination number occurs, but still some disorder persists. The effect of the annealing is to release the defects increasing the volume/surface ratio of the Co and Ag. The Debye—Waller factors are quite small indicating a small spread on distances. 3.3. Magnetization As can be seen from Fig. 6, the in-plane hysteresis loops change substantially with the annealing temperature. There is no substantial change in the shape of the curves for the as-deposited sample and those annealed at temperatures below 400°C, (Fig. 6a). It suggests a multi-domain structure at the individual Co layers and/or a exchange ferromagnetic coupling between different Co layers. In the case of the Co/Ag system, due to the rather thick Ag spacer, the interlayer interactions are weak, and the remanent magnetization of multilayers of Co/Ag composition with continuous Co layers is independent of the Ag layer thickness, as was shown by van Alphen and de Jonge [32]. The behavior of the coercivity, H , and the re# manent magnetization, M , as a function of the 3 annealing temperature are different from those observed in most granular systems [33—35]; H has # a strong dependence on the granular size [36]. In general, if the particle size is sufficiently large, the magnetization is changed mainly by domain wall motion. In this case, much less energy is necessary
Fig. 6. The in-plane hysteresis loops at room temperature for Co/Ag; (a) as-deposited; (b) annealed at 400°C and (c) annealed at 600°C for 10 min.
to change the magnetization than in a single-domain particle, and a small H is expected. For # smaller grains, namely single-domain particles, the magnetization is changed by spin rotation, requiring
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W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
larger energy than for domain wall motion, thus resulting in a large H . Further decrease of the grain # size causes H to decrease again due the thermal # activation effects, leading to the spontaneous rotation of the magnetic moments. In the present case, the dependence of H as # a function of the annealing temperature is as follows: annealing at temperatures below 400°C is not sufficient to destroy the layered structure, which maintains its ferromagnetic behavior. A multi-domain magnetic state within each layer is present and as a consequence, a low H is observed. For # annealing at 400°C (Fig. 6b), a breakup of the layers is obtained, resulting in a random distribution of sizes and moments of the magnetic particles. At this thermal treatment, the magnetic interaction and size differences between Co particles lead to a coercivity distribution resulting in a large inplane hysteresis and remanence, as is observed in the figure. The decrease of H after annealing at # 600°C (Fig. 6c) is attributed to further breakup of the layers, leading to the formation of very small granules. Although this treatment promotes the breakup of the layer structure higher than that for annealing at 400°C, it is not sufficient for complete isolation of the magnetic grains. A weak magnetic interaction between the grains is maintained, which is confirmed by the shape of the hysteresis loops. The sizes of the grains after annealing at 600°C are below the critical diameter, around 100 A_ [36] for single-domain Co particles, where the thermal activation effects are expected to appear. The saturation magnetization, M , shows a 4 roughly constant value up to 200°C. This behavior indicates a maintenance of the multilayer characteristics. For temperatures higher than 200°C, the M values start to decrease as a result of the degra4 dation of the multilayered structure, see Table 1. In addition to the microstructure characterization, the size of the superparamagnetic particles can be estimated by analyzing the experimental magnetization curves. In an assembly of very small particles at sufficiently high temperature, the magnetization vectors of the particles are thermally agitated over their potential barriers, permitting them to rotate to their equilibrium directions. Such particles are referred to as ‘unblocked’ or superparamagnetic (SPM) ones.
In some cases, despite the magnetization not reaching saturation, the hysteresis loops show M and H greater than zero. In a system consisting 3 # of SPM particles only, the hysteresis could be explained by the presence of ferromagnetic interparticle interactions. As the remanence reaches relatively large values, we consider our samples as consisting of two magnetic phases: (i) non-interacting SPM particles and (ii) ‘blocked’ (interacting and/or larger ferromagnetic) grains, which we will call FM particles. The average size (D) of the SPM particles and the saturation magnetization can be obtained by fitting the experimental hysteresis loops taking into account both FM and SPM contributions [18,37]. The magnetization M(H) for our samples can be written as M(H)"MFM(H)#MSPM(H).
(1)
The term MFM(H) gives the ferromagnetic fraction of the particles, which are assumed to have either cubic magnetocrystalline anisotropy or uniaxial anisotropy. For some relatively low magnetic fields, the FM particles have saturated and hence the first term equals its saturation value MFM independent 4 of the magnetic field H. Here it is accepted that MFM"M /0.866, which holds for a disordered sys4 3 tem of single-domain cubic anisotropy particles with negative first-order magnetocrystalline anisotropy constant K (the case of Co (FCC)) [38]. 1 Thus, for high magnetic fields (2) MSPM(H)"M(H)!MFM. 4 Under the assumption of weak interactions, the magnetization of a superparamagnetic system with uniform particle size can be described by the Langevin equation ¸(a)"coth(a)!1/a. In real granular systems it is necessary to consider a distribution of the particle size. Therefore, the MSPM(H) term should be described by a weighted superposition of Langevin functions
P A B
= kH ¸ f (») d», (3) k ¹ 0 B where k"M » is the magnetic moment of 4 a single-domain particle with saturation magnetization M and volume », and f (») is the particle 4 MSPM(H)"M 4
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
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size distribution. Assuming spherical particles of diameter D for simplicity, a log-normal particle size distribution with a width p,
A
B
1 (ln D!ln DM )2 f (D)" exp ! 2(ln p)2 J2p ln p
(4)
is often used, where »"pD3/6. Assuming certain value for M , MSPM(H), Eq. (2) can be fitted to 4 Eq. (3); therefore, f (D) and p can be obtained. The diameters of the Co particles thus obtained are listed in Table 1. Once MSPM(H) is obtained, the ferromagnetic part of the experimental magnetization can be fitted. Assuming that the FM particles are large enough so the magnetization reversals which occur due to a combination of thermal and field-induced effects can be neglected, one can use the model curve for a disordered fine particles systems with four easy magnetization axes, M/M "f (H/(2M / 4 4 DK D)) [39,40]. Along with M , it is sufficient to 1 3 know just one more parameter of the curve, the coercivity HFM, obtained from the curve M(H)! # MSPM(H). Thus, the total fitting demagnetization curve is made up of a superposition of both SPM and FM fitting curves. These curves along with the experimental data are shown in Fig. 7a for one representative sample (¹ "400°C). The agreement !//. between theory and experiment is evident from the figure, even in the low-field region. The good fit supports the assumption of weak interactions, also confirmed by the shape of the dMDC plot [41] constructed for the same sample, shown in Fig. 7b. The technique of dM plot is based on a comparison of the remanence curves, the isothermal remanent magnetization curve, M (H), and the DC demag3 netization remanence curve, M (H) [42]. In the $ case of no interaction and uniaxial anisotropy, dM(H) is zero for all values. On the other hand, Henkel plots, i.e. M (H) versus M (H), calculated $ 3 for the case of randomly oriented non-interacting spherical particles with cubic anisotropy (3 and 4 easy magnetization axes) are non-linear in a ‘positive’ sense with the curves concave downwards, leading to positive dMDC plots [41]. The shape of these plots depends on the initial demagnetization state as well. Our experimental plot for the sample
Fig. 7. (a) Experimental data and fitting curves; (b) reduced remanence curves for the sample annealed at 400°C; (c) dMDC plots for the as-deposited sample and samples annealed at 400 and 600°C.
annealed at 400°C has a shape very similar to the model one, except it has slightly lower values in the non-zero range, indicating weak negative interactions or demagnetizing interactions, in the sense that the interactions have the effect to stabilize the demagnetized state, for example by formation of flux closure structures. Fig. 7c shows the dMDC plots for the as-deposited sample and for the ones annealed at 400 and 600°C. A decrease of the negative interactions for the last sample is observed, and the shape of its dMDC plot is similar to the one for the non-interacting case. This result is in agreement with the idea that the annealing at 600°C promotes the breakup of Co platelets, decreasing their sizes and reducing the magnetic interaction, thus some SPM particles are formed.
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W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
3.4. Magnetoresistance The MR for the as-deposited sample as shown in Fig. 8 displays a feature of anisotropic magnetoresistance effect (AMR): the resistivity increases when the magnetic field is longitudinal with respect to the current flow (positive MR), and decrease for the transverse magnetic field (negative MR) [43]. This behavior of AMR is dependent on the direction of the spontaneous magnetization, being due to the domain wall motion, and is characteristic for ferromagnetic metals [44]. This anisotropy in MR is caused by the shape of magnetic layers, their discontinuities, and by the shape anisotropy of the individual magnetic plates in each layer. The AMR persists for samples annealed at temperatures below 400°C. The values observed can be explained by supposing that there exists either a small portion of domain in adjacent layers that couple ferromagnetically or plates with randomly distributed magnetization, or by the coexistence of both states. After annealing at 400°C the magnetoresistance changes from AMR to GMR with a sharp maximum of the resistance at H , #
Fig. 8. The room temperature AMR for Co/Ag as-deposited sample. In-plane field with respect to the current: (a) transverse configuration; (b) longitudinal configuration. Details of the curves are showed in the inset.
reaching an amplitude of 4.5%, as can be seen from Fig. 9a. This increase in MR must be a consequence of scattering from non-aligned ferromagnetic grains and due to a distribution of the coercivities of the magnetic particles [45], as mentioned earlier. At this point the sample is formed by a regular arrangement of planes containing grains with random distribution of moments and coercivities. It is important to note that despite the high value of the GMR for this sample, there is also a second contribution to the MR curve. This contribution, represented by the tail at high fields in the MR curve, is originated by scattering of the conduction electrons in spin fluctuations due to small clusters of Co atoms. After thermal treatment at 600°C (Fig. 9a) the value of MR decreases. The lower value of MR, as compared with the sample annealed at 400°C, can be ascribed to the reduction of the interfacial spindependent scattering with the ordering of the Co clusters, the decrease in size of these clusters, and the reduction of the particle coercivity distribution, in agreement with what was observed in the hysteresis curves (Fig. 6). For these annealing temperatures, the GMR results are well correlated with the measured hysteresis loops. It is worth noting that the shape of the *R/R peak for the sample annealed at 600°C is quite different from the sharp maximum represented by the sample treated at 400°C. The sample now is characterized by a broad size distribution of the Co grains embedded in a Ag matrix. This is in accordance with the results obtained by XRD, XAS and magnetometry. When *R/R is plotted against the reduced magnetization M/M [2] (Fig. 9b and Fig. 9c), a devi4 ation from a parabolic behavior in the low-field region is observed for the samples annealed at 400 and 600°C, respectively. This result, which is characterized by a flat top, is often mentioned as a proof of the existence of magnetic interactions among particles [15,46]. However, for the sample annealed at 600°C, the flat top is reduced while approaching the parabolic curve, indicating a partial suppression of the inter-particle interactions, as compared with 400°C thermal treatment. This behavior is in accordance with the results depicted in Fig. 7c. The annealing at 600°C promotes a breakup of the
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
27
platelets reducing their sizes and the particle coercivity distribution. As a consequence, there is an enhancement of the superparamagnetic character of the sample, as depicted in Fig. 6c. The narrowing of the particle coercivities distribution leads to the decrease in GMR for this sample.
4. Conclusions
Fig. 9. (a) The room temperature GMR for Co/Ag samples annealed at 400 and 600°C, for 10 min. Magnetoresistance versus reduced magnetization M/M for same samples; (b) 400°C; 4 (c) 600°C. Open circles are experimental data; solid curves are parabolic fittings.
The synthesis of the ferromagnetic granular metal can be produced in a well controlled manner starting from multilayer structure with a subsequent breakup of the layers through thermal annealing. For the as-deposited sample the Co layers are highly disordered and discontinuous with only a short-range order according to XRD and XAS measurements. The Co/Ag interfaces have compressive stress due to the difference of the surfaces energies between the Ag and Co. This stress is responsible for the formation of islands of Co at the interfaces, which contributes to the roughness and discontinuity of the Co layers. There is also a small incorporation of Co into the Ag lattice as observed by XRD measurements. The magnetic measurements show a ferromagnetic coupling between the SPM grains forming the Co layers and/or presence of larger ferromagnetic particles. This ferromagnetic network leads to the AMR effect common for most of ferromagnetic systems. Annealing at 200°C does not significantly change the structural and magnetic behavior of the sample, although some alteration can be observed at the FT curves, in the region corresponding to the next nearest neighbor, from 3 to 5.5 A_ . This alteration is most probably due to enhancement of the discontinuity with a consequent decrease of the volume surface ratio of the Co platelets. According to XRD and XAS measurements, the ordering of Co begins around 400°C with a relaxation of stresses and defects. Only Co (FCC) is observed. For this annealing temperature there is a great modification in the magnetic behavior of the sample. The magnetoresistance changes from AMR to GMR, with an enhancement in amplitude, as a consequence of the breakdown of the Co
28
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
layers. The sample now can be seen as a planar arrangement of Co grains regularly distributed through the Ag matrix, with a distribution of the coercivities of the magnetic particles. After annealing at 600°C, there is a broadening of the size distribution of Co granules and a narrowing of the coercivity distribution. The latter explains the decrease in the GMR amplitude. The Co cluster grains are more homogeneously dispersed through the Ag matrix than after annealing at 400°C. This behavior promotes an enhancement of the superparamagnetic character with a reduction of the magnetic interaction, as observed by dMDC plots and *R/R versus M/M . 4 Acknowledgements The authors wish to acknowledge the partial support of this work by CNPq, CAPES, FINEP, FAPERGS and FAPESP Brazilian financial agencies.
References [1] A. Berkowitz, A.P. Young, J.R. Mitchell, S. Zang, M.J. Carey, F.E. Spada, F.T. Parder, A. Hutten, G. Thomas, Phys. Rev. Lett. 68 (1992) 3745. [2] J.G. Xiao, J.S. Jiang, C.L. Chien, Phys. Rev. Lett. 68 (1992) 3749. [3] B. Abeles, in: R. Wolfe (Ed.), Applied Solid State Science: Advances in Material and Device Research, Academic Press, New York, 1976, p. 1. [4] T.L. Hylton, Appl. Phys. Lett. 62 (1993) 2431. [5] A. Berkowitz, J.R. Mitchell, M.J. Carey, A.P. Young, D. Rao, A. Starr, S. Zhang, F.E. Spada, F.T. Parker, A. Hutten, G. Thomas, J. Appl. Phys. 73 (1993) 5320. [6] J.Q. Xiao, J.S. Jiang, C.L. Chien, Phys. Rev. B 46 (1992) 9266. [7] C.L. Chien, J. Appl. Phys. 69 (1991) 5267. [8] J.H. Du, Q. Li, L.C. Wang, H. Sang, S.Y. Zhang, Y.W. Du, D. Feng, J. Phys.: Condens. Matter 7 (1995) 9425. [9] J.A. Barnard, A. Waknis, M. Tan, E. Haftek, M.R. Parker, M.L. Watson, J. Magn. Magn. Mater. 114 (1992) 230. [10] S.H. Liou, S. Malhotra, Z.S. Shan, D.J. Sellmyer, S. Nafis, J.A. Woollam, C.P. Reed, R.J. De Angelis, G.M. Chow, J. Appl. Phys. 70 (1991) 5882. [11] T.L. Hylton, K.R. Coffey, M.A. Parker, J.K. Howard, Sci. 261 (1993) 1021. [12] L.F. Schelp, G. Tosin, M. Carara, M.N. Baibich, A.A. Gomes, J.E. Schmidt, Appl. Phys. Lett. 61 (1992) 1858.
[13] A.D.C. Viegas, J. Geshev, L.F. Schelp, J.E. Schmidt, J. Appl. Phys. 82 (1997) 2466. [14] J.R. Regnard, C. Revenant-Brizard, B. Dieny, B. Mevel, J. Mimault, Mater. Res. Proc. 400 (1996) 329. [15] J.F. Gregg, S.M. Thompson, S.J. Dawson, K. Ounadjela, C.R. Staddon, J. Hamman, C. Fermon, G. Saux, K. O’Grady, Phys. Rev. B 49 (1994) 1065. [16] J.C. Lodder, P. de Haan, H. van Kranenburg, J. Magn. Magn. Mater. 128 (1993) 219. [17] D.J. Kubinski, H. Holloway, J. Appl. Phys. 77 (1995) 782. [18] A.D. Viegas, J. Geshev, L.S. Dorneles, J.E. Schmidt, M. Knobel, J. Appl. Phys. 82 (1997) 3047. [19] J.R. Regnard, J. Juanhuix, C. Revenant-Brizard, B. Dieny, B. Mevel, J. Mimault, O. Proux, Solid State Commun. 97 (1996) 419. [20] G. Tourillon, E. Dartyge, A. Fontaine, M. Lemonnier, F. Bartol, Phys. Lett. A 121 (1987) 251. [21] M.C. Correa, H. Tolentino, A. Craievich, C. Cusatis, Rev. Sci. Instrum. 63 (1992) 896. [22] F.W. Lytle, R.B. Greegor, D.R. Sandstrom, E.C. Marques, J. Wong, C.L. Spiro, G.P. Huffman, F.E. Huggins, NIM 226 (1984) 542. [23] A. Michalowicz, Logiciels pour la Chimie, vol. 102, Ed. Societe´ Franc7 aise de Chimie, Paris, 1991. [24] C.L. Chien, J.Q. Xiao, J.S. Jiang, J. Appl. Phys. 73 (1993) 5309. [25] E.A. Brandes, Smithells Metals Reference Book, 6th ed., Butterworths, London, 1983. [26] L.I. Maisser, R. Glang (Eds.), Handbook of Thin Film Technology, McGraw-Hill, New York, 1983. [27] R.H. Marion, J.B. Cohen, Advances in X-Ray Analysis, vol. 15, Plenum Press, New York, 1975, p. 466. [28] C.S. Barret, T.B. Massalski, Structure of Metals, 3rd ed., McGraw-Hill, New York, 1966, p. 357. [29] H. Sato, O. Kitakami, T. Sakurai, Y. Shimada, Y. Otani, K. Fukamichi, J. Appl. Phys. 81 (1997) 1858. [30] J.J. Rehr, S.I. Zabinsky, R.C. Albers, Phys. Rev. Lett. 69 (1992) 3397. [31] J.J. Rehr, Jpn. J. Appl. Phys. 32 (1993) 8. [32] E.A.M. van Alphen, W.J.M. de Jonge, Phys. Rev. B 51 (13) (1995) 8182. [33] J.R. Childress, C.L. Chien, M. Nathan, Appl. Phys. Lett. 56 (1990) 95. [34] Y.X. Zhang, S.H. Liou, R.J. DeAngelis, K.W. Lee, C.P. Reed, A. Nazareth, J. Appl. Phys. 69 (1991) 5273. [35] J.R. Childress, C.L. Chien, J. Appl. Phys. 70 (1991) 5885. [36] F.E. Luborsky, J. Appl. Phys. 32 (1961) 171S. [37] B.J. Hickey, M.A. Howson, S.O. Musa, G.J. Tomka, B.D. Rainford, N. Wiser, J. Magn. Magn. Mater. 147 (1995) 253. [38] R. Gans, Ann. Phys. 15 (1932) 28. [39] J. Geshev, O. Popov, V. Masheva, M. Mikhov, J. Magn. Magn. Mater. 92 (1990) 185. [40] I. Joffe, R. Heuberger, Phil. Magn. 314 (1974) 1051. [41] J. Geshev, M. Mikhov, J. Magn. Magn. Mater. 104—107 (1992) 1569. [42] P.E. Kelly, K. O’Grady, P.I. Mayo, R.W. Chantrell, IEEE Trans. Magn. 2 (1989) 3881.
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[45] E. Kolb, M. Mulloy, C. Dupas, M. Galtier, D. Renard, J.P. Renard, F. Trigui, E. Ve´lu, J. Magn. Magn. Mater. 148 (1995) 315. [46] P. Allia, M. Knobel, P. Tiberto, F. Vinai, Phys. Rev. B (1995) 15398.
Structural evolution and magnetic behavior of Co/Ag multilayers submitted to thermal annealing W.H. Flores!, S.R. Teixeira!,*, J. Geshev!, J.B.M. da Cunha!, P.J. Schilling", A. Traverse#, M.C. Martins Alves$ ! Instituto de Fı& sica-UFRGS, C.P. 15051, 91501-970 Porto Alegre, Brazil " Center for Advanced Microstructures and Devices, Baton Rouge, USA # LURE-Universite& Paris-Sud, 91405 Orsay Cedex, France $ Laborato& rio Nacional de Luz Sı& ncrotron, LNLS, C.P. 6192, 13081-970 Campinas, Brazil Received 13 September 1997; received in revised form 13 March 1998
Abstract We report on the structural evolution and the magnetic behavior of the Co/Ag multilayered films deposited in a UHV chamber at room temperature, submitted to 10 min thermal annealing at temperatures ranging from 100 to 600°C. The structural characterization was performed using X-ray diffraction and X-ray absorption spectroscopy techniques. Magnetoresistance and magnetization measurements were used to study the evolution and magnetic behavior of the samples. The results show that, besides the roughness at the interfaces and the structural disorder of the Co layers, the as-deposited sample has a compressive stress at Ag—Co interface originated by the difference between the surface energies of Ag and Co. After annealing at 400°C, there is a breakup of the layers accompanied by a relaxation of the stress and defects as well as partial crystallographic ordering of the Co clusters. The room-temperature magnetoresistance change from anisotropic magnetoresistance to giant magnetoresistance with a sharp maximum, reaching 4.5%. ( 1998 Elsevier Science B.V. All rights reserved. PACS: 74.25.Ha; 75.70.i; 75.70.Pa Keywords: Magnetic properties; Thin films; Multilayers
1. Introduction The observation of the giant magnetoresistance (GMR) effect in granular films [1,2], where the
* Corresponding author. Tel.: #55 51 316 6498/7111; fax: #55 51 319 1762; e-mail: [email protected].
magnetic material is formed by small precipitates embedded in a non-magnetic matrix, has stimulated remarkable interest because of its potential technological applications as magnetic sensors in the magnetic storage technology [3,4], and also by its interest in fundamental scientific research. The granular films can present a GMR effect even larger than in multilayers. They are often obtained by
0304-8853/98/$19.00 ( 1998 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 9 8 ) 0 0 1 4 2 - 5
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co-sputtering of two immiscible metallic elements [5], or by sputtering of metastable alloys [6]. The size of the magnetic particles is of great interest, as it determines to a large extent, the magnetic behavior of the granular alloy. Its control is dictated by substrate temperature during the deposition [7], or after deposition by subsequent annealing [6], until the optimum annealing temperature that maximizes the GMR effect is obtained. However, for granular alloy systems, the particle size, shape and temperature dominate the field dependence of the GMR effect, making difficult the use of these alloys in low-field sensitivity devices, because of the large distribution of particle sizes and shapes. It is also impossible to predict the optimum temperature because it depends on the deposition methods used to produce the alloys as well as on the sample components. Another important parameter that should be controlled during the deposition process of granular films is the concentration of the magnetic metal [8—10], which has an important influence on the GMR effect. At high magnetic metal concentration, the magnetic phase percolates throughout the material and the sample appears to be a ferromagnet. Recently, Hylton et al. [11] have proposed the preparation of discontinuous multilayers by annealing the sputtered multilayers in order to minimize the effects of crystal and shape anisotropy. With this procedure, it is possible to control the composition and to minimize the effects of the shape and the size of the magnetic particles on the GMR. In fact, the discontinuous multilayers will be composed of an island-like network of multidomain magnetic particles, within the magnetic layers separated by a continuous non-magnetic spacer. The interfaces between the ferromagnetic and non-magnetic layers are not necessarily flat and continuous, they have a roughness and contain some discontinuities. The ferromagnetic layers form a network consisting of connected ferromagnetic grains. Hence, a multidomain magnetic state within each layer is expected. The aim here is to achieve a complete insulation of the individual magnetic grains through successive annealing within a certain ordered layer structure. The role of annealing on the magnetic properties of the Co/Ag systems, has been investigated re-
cently [12—19]. In particular, the magnetic interactions between magnetic entities with two distinct magnetic phases, and the correlation of the structural parameters obtained by a detailed analysis of the hysteresis loops, have been described and discussed [15—18]. In this context, the presence of superparamagnetic and ferromagnetic particles in this system with inter-particle coupling leads to a complex behavior. Thus, the magnetic properties are dependent on the topological structure of the ferromagnetic component, especially on their sizes and structure. This paper adopts the same procedure suggested by Hylton et al. [11] to obtain a granular system. Starting from a dual electron beam deposited Co/Ag multilayer structure with thin magnetic layers, a breakup of the layers through successive annealing is performed to obtain a regular granular magnetic network embedded in a non-magnetic matrix. The main interest here, is to correlate the structural evolution with the magnetotransport and macroscopic magnetic behavior of the sample after each thermal treatment, to understand the drastic changes in the magnetoresistance effect of this system.
2. Experimental The multilayers [Ag(50 A_ )/Co(15 A_ )] , were 10 prepared at room temperature by alternate deposition in ultra-high vacuum on a Si(1 1 1) substrate covered with a thick SiO layer. The depositions 2 were performed using a dual electron beam Balzers evaporator. The base pressure in the evaporation chamber was 1.0]10~8 mbar. All samples were prepared with a deposition rate of 1 A_ /s, monitored by a quartz balance. The total thickness of the samples was 65 nm. The thickness of the Ag and Co layers were chosen to give an overall concentration of the approximately 20 at% for the ferromagnetic metal. The multilayers were annealed in a high vacuum furnace (10~7 mbar) for 10 min at temperatures ranging from 100 to 600°C. After each anneal, the samples were cooled down to room temperature for subsequent analyses. The structural characterization of the samples were made using X-ray diffraction (XRD) and
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
X-ray absorption spectroscopy (XAS), with X-ray absorption near edge (XANES) and extended X-ray absorption fine structure (EXAFS) modes. The XRD was performed in h—2h Bragg—Brentano geometry at room-temperature employing Cu Ka radiation. XAS experiments at the Co K-edge were performed on line D42 at the DCI storage ring, LURE-Orsay, running at an energy of 1.85 GeV and a current of 300 mA, and at LSU Center for Advanced Microstructures and Devices (CAMD). At the D42 line, the X-rays were monochromatized with a Si(3 3 1) channel cut single crystal with an energy resolution of 1 eV at the Co K-edge. The incident beam was collimated by slits and its intensity was measured by an ionization chamber. XANES and EXAFS spectra of the samples and model compounds were recorded in the total electron yield [20] due to the presence of the SiO 2 substrate on the multilayers. The XANES data were collected in scans with an energy step of 0.5 eV and two scans were added to obtain a good signalto-noise ratio. At CAMD, the synchrotron X-ray beam was monochromatized using the LNLS double-crystal monochromator equipped with Si(1 1 1) crystals [21]. All data were collected in total electron yield (TEY) at room temperature [22]. The EXAFS spectra were analyzed using standard methods [23]. It involves a background subtraction by means of a polynomial function normalized to the height of the absorption edge step. The various neighboring shells were obtained by a Fourier transform of the EXAFS signal from 3 to 12.5 A_ using a Kaiser window with a coefficient of 2.5. By an inverse Fourier transform into k space, the EXAFS oscillations corresponding to only one neighbor shell were obtained. Structural parameters are obtained from least squares fitting using experimental phase and amplitude functions deduced from a model compound. We have used phase shift and amplitudes extracted from a Co/SiO /Si (1 1 1) sample annealed at 600°C. 2 The magnetic properties were investigated at room temperature using a conventional vibrating sample magnetometer (VSM) and alternate gradient magnetometer (AGM), with the applied field always parallel to the film surface. The mag-
19
netoresistence (MR) was measured at room temperature using a standard four-point configuration with in-plane current between 0.1 and 1 mA. These measurements were performed with the magnetic field parallel and perpendicular to the current. The fractional variation of MR, *R/R, was taken as R(H)!R(H"0)/R(H"0), where R(H"0) is the value in the initial, H"0, state; here R is the resistance and H is the applied magnetic field.
3. Results and discussion 3.1. X-ray The XRD patterns in a Bragg—Brentano geometry for the as-deposited and annealed samples are shown in Fig. 1. The as-deposited sample, Fig. 1a, exhibits peaks attributable to Ag(1 1 1), Ag(2 0 0), Ag(2 2 0), Ag(3 1 1), and Ag(2 2 2) reflections, corresponding to FCC polycrystalline structure. The samples show a preferential orientation in the (1 1 1) direction. As can be seen, the relative intensities of the high index lines compared to the (1 1 1) line, are much lower than the relative intensities of the high index lines also compared to the (1 1 1) line of a powder spectrum, even considering that in the present case, the spectra show an overlap between the Co(1 1 1) and Ag(2 0 0) lines. Due to the roughness of the interfaces, no satellites are observed indicating that the multilayer is most probably discontinuous. The Co(1 1 1) (FCC) reflection cannot be distinguished because it coincides with the Ag(2 0 0) reflection. It is important to note that for the as-deposited sample, no other reflections corresponding to Co are visible, indicating that the Co layers have a very disordered structure. The Ag(1 1 1) reflection, as can be observed from Fig. 1a, is split into two overlapping lines, 2h" 37.85 and 38.57°, differing from the samples obtained by sputtering of metastable Co—Ag alloys [6—24], or co-sputtering of Ag and Co [5,8]. The two split lines can be identified by the shoulder at the low-angle part and by a more intense peak at the right side of the whole Ag(1 1 1) reflection. To see if the origin of the split line was under stress at the Ag/Co interfaces, an u tilt measurement for the
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Fig. 2. Stress measurements, u tilt, of the Ag(1 1 1) planes. Plot of d versus sin2 u. 111
Fig. 1. The XDR patterns for Co/Ag: (a) as-deposited sample, (b) and (c) samples annealed at 400 and 600°C, respectively.
two peaks was performed. The results displayed in Fig. 2 as d versus sin2 u, clearly show a com111 pressive stress of the Ag(1 1 1) split line. Using the well known expression p"aE/(1#l), where a is the slope of the fitted straight line, E the Young’s modulus, and l the Poisson ratio (for Ag, E"7.58]1010 N/m2 and l"0.38), the values obtained for the compressive stress were 23 and 5.5 GPa for the low- and high-angle peaks, respectively. The stress comes from the Co/Ag interfaces, probably due to the differences of the surface energies of Ag and Co (c "1.3 J/m2 and A' c "2.71 J/m2 [25,26]). When Co is deposited on C0 a Ag layer, there will be a formation of small Co islands leading to a stressed Co/Ag interface, because Co does not wet the Ag film. This stressed region corresponds to the shoulder that appear at the low-angle peak region of the Ag(1 1 1) reflec-
tion. On the other hand, as the thickness of the Ag layer increases, the stress decreases and this less stressed region corresponds to the high-angle peak. The Ag wets the Co surface more, so Ag/Co interface is expected to be weakly stressed. The oscillations occurring in d values around the fitting 111 straight line that appear in Fig. 2 are due to the texture effects in the Ag layers [27]. The high-angle peak at Ag(1 1 1) is slightly shifted to the right from the expected position of the bulk Ag(1 1 1) peak as is indicated in Fig. 1a. This can be due to a small incorporation of Co atoms forming a solid solution in Ag. Using the Veggard’s law [28], it was estimated that 7.5% of Co are substitutional in the Ag lattice near Co/Ag interfaces. Annealing at temperatures below 400°C, for 10 min, does not significantly change the XRD patterns, which are very similar to the ones for the as-deposited sample. After annealing at 400°C, for the same time (Fig. 1b) the Ag peaks narrow considerably, and although very small, a second peak begins to emerge at 2h around 52°, corresponding to Co(2 0 0) (FCC), as indicated by the arrow in the figure. For this annealing temperature there is a narrowing of the diffraction linewidths as a consequence of an ordering of the crystalline structure due to the release of defects and stresses. The shoulder at the left side of the Ag(1 1 1) peak almost disappears. Despite the XRD results showing a polycrystalline spectra, the Ag becomes highly S1 1 1T-textured as the annealing temperature
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
increases. This behavior can be observed by the substantial increase of the Ag(1 1 1) and (2 2 2) peaks as compared with the intensities of the Ag(2 2 0) and (3 1 1) reflections. Although bulk Co (HCP) is stable below the allotropic transformation temperature of about 420°C, there are no evidences of Co (HCP) phase in our samples, as FCC phase fine particles are usually stable at ambient temperature (see, e.g. Sato et al. [29] and the references therein). The ordering temperature of the Co cannot be precisely estimated due to the overlap of the Ag(2 0 0) and Co(1 1 1) peaks. Annealing at 600°C, for 10 min, causes narrowing of the peaks and growing up of Co(2 2 0) (FCC) reflection. The figure clearly shows Co (FCC) clusters embedded in a Ag(1 1 1) textured matrix. 3.2. XAS Fig. 3 shows the Co K-edge XANES data recorded for Co/Ag samples as a function of the annealing temperature. The XANES data of a bulk Co are given for comparison. The reference energy value (0 eV) corresponds to the first inflection point of the metallic Co edge (7709 eV).
Fig. 3. The XANES spectra from the Co edges for Co/Ag asdeposited, and for some annealed samples. The spectra of Co reference and simulations are shown for comparison.
21
In pure Co, peaks labeled 1, 2 and 3 arise from multiple scattering effects of the photoelectron by the different shells being characteristic of the HCP symmetry. Simulations using a FEFF code [30,31] for Co (FCC) and (HCP) allow us to differentiate the two Co structures, as can be seen in Fig. 3. By comparison with XRD spectra we conclude that only Co (FCC) is present. The XANES profiles show an evolution with a clear difference between the features observed for the as-deposited sample and the one annealed at 200°C, which are very similar, and those annealed at 400 and 600°C. As can be seen from the figure, there is a decrease in intensity of the peak 1 while an increase in intensity is observed for the peaks 2 and 3, with a visible sharpening of the peak 3. This means that a structural modification had occurred around Co after annealing at 400°C in agreement with what is observed from XRD. This evolution comes from the release of defects and stresses, as observed in the X-ray spectra (Fig. 1), originating the ordering of the Co and Ag. It is interesting to note that the XANES spectra are quite different from that of the reference (Co foil). Fig. 4a shows the EXAFS signal for the as-deposited and annealed samples. Co foil data are included for reference as well. The spectra of the as-deposited and the annealed samples are quite different from the Co foil ones. As the annealing temperature increases the intensity increases reaching values even higher than that of the reference. The small oscillations appearing at the right hand side of the first EXAFS oscillation are due to the existence of small Co clusters. Details about this fine structure will be published elsewhere. Fig. 4b presents the k3 weighted Fourier transform (FT) of the Co K-edge EXAFS signals depicted in Fig. 4a. The figure also shows the simulation using FEFF code for Co (HCP) and (FCC). Comparing the FT obtained experimentally with the simulation we see that the Co present a FCC structure, in agreement with that observed by XRD. This difference is evidenced by comparing the intensities of the peaks 3 and 4 with the simulated spectrum corresponding to the Co (FCC), which are quite different from the FT obtained for the Co foil. The FT of Co metal exhibits four peaks located at 2.2, 3.4, 4.1 and 4.7 A_ (uncorrected from the phase shift). The first peak is
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W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
Fig. 5. The inverse Fourier transform of the first shell: open circles are the experimental data and full line, fitted curve.
Fig. 4. (a) The EXAFS measurements; (b) Fourier transforms of signals from the Co K-edge for Co/Ag as-deposited, and for some annealed samples. *k"3—12.5 A_ , Kaiser window q"2.5. The spectra of Co reference and simulations are shown for comparison.
related to the Co—Co near neighbor (NN) distance and the others are related to more distant neighbors. In contrast to XRD measurements, there are no evidences of NN Co—Ag contribution, as was
reported by Regnard et al. [14,19] for the sputter deposited AgCo films. To clarify this differences, low temperature EXAFS measurements will be performed in near future. The amplitudes of the FT peaks for the as-deposited and annealed at 200°C samples are small compared with the Co reference. This is an evidence for a reduction in the neighbor coordination for all the shells. This effect disappears for higher annealing temperatures where the intensity of the peaks is close to that of the Co foil, indicating an ordering of the Co structure in agreement with what was observed by XRD. In order to obtain quantitative estimations of the bond lengths and the coordination numbers around the cobalt atoms, the inverse Fourier transforms of the first shell have been fitted. A typical fitting is shown in Fig. 5. The backscattering amplitude and the phase shift values were deduced from the EXAFS spectrum of a Co/SiO /Si(1 1 1) sample 2 annealed at 600°C. The fit parameters are given in Table 1. The Fourier filtered spectra were fitted with a single shell of neighboring atoms, as only Co atoms are present in the first shell. The curve fitting indicates a reduction in NN coordination number, for the as-deposited and annealed samples at 200 and 400°C, see Table 1. For the sample treated at 600°C this number is 12 as for the bulk, indicating an ordered structure of the Co clusters. The small coordination number for as-deposited sample indicates a disordered structure of the Co, as observed
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
23
Table 1 Parameters determined from magnetic measurement and by fitting the first shell in the Fourier-filtered EXAFS spectra for Co/Ag as a function of the annealing temperature, ¹ . The phase shift and amplitudes were extracted from the sample Co/SiO /Si(1 1 1) treated at A 2 600°C. M , saturation magnetization in units of emu/cm3; D, diameter of the particles in A_ ; N, coordination number; R, distances of 4 nearest neighbors in A_ , *p, Debye—Waller factor relative to the model compound in (A_ 2) Magn.
EXAFS
¹ A
M 4
D
N
R
*p2
Fit quality (]10~4)
As-dep. 200°C 400°C 600°C Co ref.
1480.64 1477.48 1383.53 1413.28 1466.72
— — 48.9 42.2 —
7.3 8 10.1 12 12
2.52 2.52 2.51 2.52 2.50
1.3]10~6 2.9]10~6 1.7]10~5 1.0]10~7 —
19 8.3 7.1 0.2 —
by XRD measurements. After annealing at 200 and 400°C, an increase of the coordination number occurs, but still some disorder persists. The effect of the annealing is to release the defects increasing the volume/surface ratio of the Co and Ag. The Debye—Waller factors are quite small indicating a small spread on distances. 3.3. Magnetization As can be seen from Fig. 6, the in-plane hysteresis loops change substantially with the annealing temperature. There is no substantial change in the shape of the curves for the as-deposited sample and those annealed at temperatures below 400°C, (Fig. 6a). It suggests a multi-domain structure at the individual Co layers and/or a exchange ferromagnetic coupling between different Co layers. In the case of the Co/Ag system, due to the rather thick Ag spacer, the interlayer interactions are weak, and the remanent magnetization of multilayers of Co/Ag composition with continuous Co layers is independent of the Ag layer thickness, as was shown by van Alphen and de Jonge [32]. The behavior of the coercivity, H , and the re# manent magnetization, M , as a function of the 3 annealing temperature are different from those observed in most granular systems [33—35]; H has # a strong dependence on the granular size [36]. In general, if the particle size is sufficiently large, the magnetization is changed mainly by domain wall motion. In this case, much less energy is necessary
Fig. 6. The in-plane hysteresis loops at room temperature for Co/Ag; (a) as-deposited; (b) annealed at 400°C and (c) annealed at 600°C for 10 min.
to change the magnetization than in a single-domain particle, and a small H is expected. For # smaller grains, namely single-domain particles, the magnetization is changed by spin rotation, requiring
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W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
larger energy than for domain wall motion, thus resulting in a large H . Further decrease of the grain # size causes H to decrease again due the thermal # activation effects, leading to the spontaneous rotation of the magnetic moments. In the present case, the dependence of H as # a function of the annealing temperature is as follows: annealing at temperatures below 400°C is not sufficient to destroy the layered structure, which maintains its ferromagnetic behavior. A multi-domain magnetic state within each layer is present and as a consequence, a low H is observed. For # annealing at 400°C (Fig. 6b), a breakup of the layers is obtained, resulting in a random distribution of sizes and moments of the magnetic particles. At this thermal treatment, the magnetic interaction and size differences between Co particles lead to a coercivity distribution resulting in a large inplane hysteresis and remanence, as is observed in the figure. The decrease of H after annealing at # 600°C (Fig. 6c) is attributed to further breakup of the layers, leading to the formation of very small granules. Although this treatment promotes the breakup of the layer structure higher than that for annealing at 400°C, it is not sufficient for complete isolation of the magnetic grains. A weak magnetic interaction between the grains is maintained, which is confirmed by the shape of the hysteresis loops. The sizes of the grains after annealing at 600°C are below the critical diameter, around 100 A_ [36] for single-domain Co particles, where the thermal activation effects are expected to appear. The saturation magnetization, M , shows a 4 roughly constant value up to 200°C. This behavior indicates a maintenance of the multilayer characteristics. For temperatures higher than 200°C, the M values start to decrease as a result of the degra4 dation of the multilayered structure, see Table 1. In addition to the microstructure characterization, the size of the superparamagnetic particles can be estimated by analyzing the experimental magnetization curves. In an assembly of very small particles at sufficiently high temperature, the magnetization vectors of the particles are thermally agitated over their potential barriers, permitting them to rotate to their equilibrium directions. Such particles are referred to as ‘unblocked’ or superparamagnetic (SPM) ones.
In some cases, despite the magnetization not reaching saturation, the hysteresis loops show M and H greater than zero. In a system consisting 3 # of SPM particles only, the hysteresis could be explained by the presence of ferromagnetic interparticle interactions. As the remanence reaches relatively large values, we consider our samples as consisting of two magnetic phases: (i) non-interacting SPM particles and (ii) ‘blocked’ (interacting and/or larger ferromagnetic) grains, which we will call FM particles. The average size (D) of the SPM particles and the saturation magnetization can be obtained by fitting the experimental hysteresis loops taking into account both FM and SPM contributions [18,37]. The magnetization M(H) for our samples can be written as M(H)"MFM(H)#MSPM(H).
(1)
The term MFM(H) gives the ferromagnetic fraction of the particles, which are assumed to have either cubic magnetocrystalline anisotropy or uniaxial anisotropy. For some relatively low magnetic fields, the FM particles have saturated and hence the first term equals its saturation value MFM independent 4 of the magnetic field H. Here it is accepted that MFM"M /0.866, which holds for a disordered sys4 3 tem of single-domain cubic anisotropy particles with negative first-order magnetocrystalline anisotropy constant K (the case of Co (FCC)) [38]. 1 Thus, for high magnetic fields (2) MSPM(H)"M(H)!MFM. 4 Under the assumption of weak interactions, the magnetization of a superparamagnetic system with uniform particle size can be described by the Langevin equation ¸(a)"coth(a)!1/a. In real granular systems it is necessary to consider a distribution of the particle size. Therefore, the MSPM(H) term should be described by a weighted superposition of Langevin functions
P A B
= kH ¸ f (») d», (3) k ¹ 0 B where k"M » is the magnetic moment of 4 a single-domain particle with saturation magnetization M and volume », and f (») is the particle 4 MSPM(H)"M 4
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
25
size distribution. Assuming spherical particles of diameter D for simplicity, a log-normal particle size distribution with a width p,
A
B
1 (ln D!ln DM )2 f (D)" exp ! 2(ln p)2 J2p ln p
(4)
is often used, where »"pD3/6. Assuming certain value for M , MSPM(H), Eq. (2) can be fitted to 4 Eq. (3); therefore, f (D) and p can be obtained. The diameters of the Co particles thus obtained are listed in Table 1. Once MSPM(H) is obtained, the ferromagnetic part of the experimental magnetization can be fitted. Assuming that the FM particles are large enough so the magnetization reversals which occur due to a combination of thermal and field-induced effects can be neglected, one can use the model curve for a disordered fine particles systems with four easy magnetization axes, M/M "f (H/(2M / 4 4 DK D)) [39,40]. Along with M , it is sufficient to 1 3 know just one more parameter of the curve, the coercivity HFM, obtained from the curve M(H)! # MSPM(H). Thus, the total fitting demagnetization curve is made up of a superposition of both SPM and FM fitting curves. These curves along with the experimental data are shown in Fig. 7a for one representative sample (¹ "400°C). The agreement !//. between theory and experiment is evident from the figure, even in the low-field region. The good fit supports the assumption of weak interactions, also confirmed by the shape of the dMDC plot [41] constructed for the same sample, shown in Fig. 7b. The technique of dM plot is based on a comparison of the remanence curves, the isothermal remanent magnetization curve, M (H), and the DC demag3 netization remanence curve, M (H) [42]. In the $ case of no interaction and uniaxial anisotropy, dM(H) is zero for all values. On the other hand, Henkel plots, i.e. M (H) versus M (H), calculated $ 3 for the case of randomly oriented non-interacting spherical particles with cubic anisotropy (3 and 4 easy magnetization axes) are non-linear in a ‘positive’ sense with the curves concave downwards, leading to positive dMDC plots [41]. The shape of these plots depends on the initial demagnetization state as well. Our experimental plot for the sample
Fig. 7. (a) Experimental data and fitting curves; (b) reduced remanence curves for the sample annealed at 400°C; (c) dMDC plots for the as-deposited sample and samples annealed at 400 and 600°C.
annealed at 400°C has a shape very similar to the model one, except it has slightly lower values in the non-zero range, indicating weak negative interactions or demagnetizing interactions, in the sense that the interactions have the effect to stabilize the demagnetized state, for example by formation of flux closure structures. Fig. 7c shows the dMDC plots for the as-deposited sample and for the ones annealed at 400 and 600°C. A decrease of the negative interactions for the last sample is observed, and the shape of its dMDC plot is similar to the one for the non-interacting case. This result is in agreement with the idea that the annealing at 600°C promotes the breakup of Co platelets, decreasing their sizes and reducing the magnetic interaction, thus some SPM particles are formed.
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W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
3.4. Magnetoresistance The MR for the as-deposited sample as shown in Fig. 8 displays a feature of anisotropic magnetoresistance effect (AMR): the resistivity increases when the magnetic field is longitudinal with respect to the current flow (positive MR), and decrease for the transverse magnetic field (negative MR) [43]. This behavior of AMR is dependent on the direction of the spontaneous magnetization, being due to the domain wall motion, and is characteristic for ferromagnetic metals [44]. This anisotropy in MR is caused by the shape of magnetic layers, their discontinuities, and by the shape anisotropy of the individual magnetic plates in each layer. The AMR persists for samples annealed at temperatures below 400°C. The values observed can be explained by supposing that there exists either a small portion of domain in adjacent layers that couple ferromagnetically or plates with randomly distributed magnetization, or by the coexistence of both states. After annealing at 400°C the magnetoresistance changes from AMR to GMR with a sharp maximum of the resistance at H , #
Fig. 8. The room temperature AMR for Co/Ag as-deposited sample. In-plane field with respect to the current: (a) transverse configuration; (b) longitudinal configuration. Details of the curves are showed in the inset.
reaching an amplitude of 4.5%, as can be seen from Fig. 9a. This increase in MR must be a consequence of scattering from non-aligned ferromagnetic grains and due to a distribution of the coercivities of the magnetic particles [45], as mentioned earlier. At this point the sample is formed by a regular arrangement of planes containing grains with random distribution of moments and coercivities. It is important to note that despite the high value of the GMR for this sample, there is also a second contribution to the MR curve. This contribution, represented by the tail at high fields in the MR curve, is originated by scattering of the conduction electrons in spin fluctuations due to small clusters of Co atoms. After thermal treatment at 600°C (Fig. 9a) the value of MR decreases. The lower value of MR, as compared with the sample annealed at 400°C, can be ascribed to the reduction of the interfacial spindependent scattering with the ordering of the Co clusters, the decrease in size of these clusters, and the reduction of the particle coercivity distribution, in agreement with what was observed in the hysteresis curves (Fig. 6). For these annealing temperatures, the GMR results are well correlated with the measured hysteresis loops. It is worth noting that the shape of the *R/R peak for the sample annealed at 600°C is quite different from the sharp maximum represented by the sample treated at 400°C. The sample now is characterized by a broad size distribution of the Co grains embedded in a Ag matrix. This is in accordance with the results obtained by XRD, XAS and magnetometry. When *R/R is plotted against the reduced magnetization M/M [2] (Fig. 9b and Fig. 9c), a devi4 ation from a parabolic behavior in the low-field region is observed for the samples annealed at 400 and 600°C, respectively. This result, which is characterized by a flat top, is often mentioned as a proof of the existence of magnetic interactions among particles [15,46]. However, for the sample annealed at 600°C, the flat top is reduced while approaching the parabolic curve, indicating a partial suppression of the inter-particle interactions, as compared with 400°C thermal treatment. This behavior is in accordance with the results depicted in Fig. 7c. The annealing at 600°C promotes a breakup of the
W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
27
platelets reducing their sizes and the particle coercivity distribution. As a consequence, there is an enhancement of the superparamagnetic character of the sample, as depicted in Fig. 6c. The narrowing of the particle coercivities distribution leads to the decrease in GMR for this sample.
4. Conclusions
Fig. 9. (a) The room temperature GMR for Co/Ag samples annealed at 400 and 600°C, for 10 min. Magnetoresistance versus reduced magnetization M/M for same samples; (b) 400°C; 4 (c) 600°C. Open circles are experimental data; solid curves are parabolic fittings.
The synthesis of the ferromagnetic granular metal can be produced in a well controlled manner starting from multilayer structure with a subsequent breakup of the layers through thermal annealing. For the as-deposited sample the Co layers are highly disordered and discontinuous with only a short-range order according to XRD and XAS measurements. The Co/Ag interfaces have compressive stress due to the difference of the surfaces energies between the Ag and Co. This stress is responsible for the formation of islands of Co at the interfaces, which contributes to the roughness and discontinuity of the Co layers. There is also a small incorporation of Co into the Ag lattice as observed by XRD measurements. The magnetic measurements show a ferromagnetic coupling between the SPM grains forming the Co layers and/or presence of larger ferromagnetic particles. This ferromagnetic network leads to the AMR effect common for most of ferromagnetic systems. Annealing at 200°C does not significantly change the structural and magnetic behavior of the sample, although some alteration can be observed at the FT curves, in the region corresponding to the next nearest neighbor, from 3 to 5.5 A_ . This alteration is most probably due to enhancement of the discontinuity with a consequent decrease of the volume surface ratio of the Co platelets. According to XRD and XAS measurements, the ordering of Co begins around 400°C with a relaxation of stresses and defects. Only Co (FCC) is observed. For this annealing temperature there is a great modification in the magnetic behavior of the sample. The magnetoresistance changes from AMR to GMR, with an enhancement in amplitude, as a consequence of the breakdown of the Co
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W.H. Flores et al. / Journal of Magnetism and Magnetic Materials 188 (1998) 17—29
layers. The sample now can be seen as a planar arrangement of Co grains regularly distributed through the Ag matrix, with a distribution of the coercivities of the magnetic particles. After annealing at 600°C, there is a broadening of the size distribution of Co granules and a narrowing of the coercivity distribution. The latter explains the decrease in the GMR amplitude. The Co cluster grains are more homogeneously dispersed through the Ag matrix than after annealing at 400°C. This behavior promotes an enhancement of the superparamagnetic character with a reduction of the magnetic interaction, as observed by dMDC plots and *R/R versus M/M . 4 Acknowledgements The authors wish to acknowledge the partial support of this work by CNPq, CAPES, FINEP, FAPERGS and FAPESP Brazilian financial agencies.
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