Low field magnetoresistance properties of (La0.75Sr0.25)1.05Mn0.95O3 polycrystalline thin films on...

Met. Mater. Int., Vol. 17, No. 6 (2011), pp. 1045~1053doi: 10.1007/s12540-011-6026-4 Published 27 December 2011

Low Field Magnetoresistance Properties of (La0.75Sr0.25)1.05Mn0.95O3Polycrystalline Thin Films on a-SiO2/Si Substrates Prepared

by ex-situ Solid Phase Crystallization

Young-Min Kang1, Alexander N. Ulyanov2, Sung-Yun Lee1, and Sang-Im Yoo1,*

1Department of Materials Science and Engineering and Research Institute of Advanced Materials,Seoul National University, Seoul 151-744, Korea

2Donetsk Physico-Technical Institute of National Academy of Sciences, Donetsk 83114, Ukraine

(received date: 23 November 2010 / accepted date: 18 July 2011)

We report the low field magnetoresistance (LFMR) properties of (La0.75Sr0.25)1.05Mn0.95O3(LSMO) films ona-SiO2/Si substrates, prepared by ex-situ solid phase crystallization of amorphous films deposited by dc-mag-netron sputtering at room temperature. The average grain size of the LSMO films was gradually increasedwith increasing annealing temperature (Tan) and film thickness. High Tan also caused the growth of an amorphousinter-diffusion layer between a-SiO2 and LSMO. The highest LFMR values of 16 and 1.0 % were achievedat 100 K, 1.2 kOe and 300 K, 0.5 kOe, respectively, from an LSMO film of 200 nm thickness annealedat 900 °C. In accordance with a modified brick layer model, grain boundary areal resistance gradually increasedwith increasing Tan and decreasing film thickness due to the penetration of the amorphous inter-diffused phaseinto the LSMO grain boundary. Improved LFMR values are attributable to modification of the LSMO grainboundary into a more effective spin-dependent scattering center.

Keywords: ferromagnetic materials, sputtering, grain boundary, magnetic properties, spintronics

1. INTRODUCTION

Perovskite manganites of general formulae Re1-xDxMnO3

(Re = rare earth and Y; D = alkaline-earth), known as colossalmagnetoresistance (CMR) materials, have drawn muchattention because of their abundant physical properties andpotential applications in spintronic and magnetoelectric devices[1]. Unfortunately, however, the intrinsic CMR effect, observedin single crystals and epitaxial films of the perovskitemanganites, is restricted only to very high magnetic fields ofover several Tesla and to a narrow temperature region nearTC [2-4]. For practical applications, however, high magne-toresistance (MR) performance should be achieved at muchlower magnetic field levels and over a wider temperature regionnear room temperature. Thus, the low field magnetoresistance(LFMR) effect, manifested as a sharp resistivity drop at lowfield magnetization process [5], has been extensively studiedfor La0.7Sr0.3MnO3 composition, which exhibits the highestCurie temperature (TC) among all manganites [1].

Previous studies [5-8] revealed that the LFMR effect couldbe attributed to spin-dependent scattering of transport electronsat grain boundaries; that is, transport of spin-polarized cur-

rent flowing across the grain boundaries depends on thedirection of the magnetization of adjacent grains. A largeLFMR effect at low temperature sharply decreases withincreasing temperature. Taking into account that a seriousdegradation of LFMR at high temperature is caused bypoorer ferromagnetism [9] and lower spin polarization [10]at grain boundaries than those within adjacent grains, manyresearch groups [11-23] have focused on grain boundaryengineering to improve LFMR at high temperature. Forinstance, enhanced LFMR properties have been reported forvarious composites between CMR manganites and insulat-ing materials. Reported insulating materials include SrTiO3

[11], Al2O3[12], ZnO [13], In2O3 [14], CeO2 [15], TiO2 [16],YSZ [17], ZrO2 [18], glass [19], etc. The highest LFMRvalue in each study was obtained by optimizing the amountof additives [11,12,15-19] or the processing conditions [13,14],implying that the LFMR value could be improved by modi-fying the grain boundary of ferromagnetic manganites into amore effective spin-dependent tunnel barrier. It was also reportedthat the LFMR property could be tuned by the nanometricgrain size modulation (10~30 nm) of La0.7Sr0.3MnO3, forwhich the LFMR performance could be explained by theeffectiveness of the spin-polarized tunneling at the grain sur-face/boundary, where field-induced strong freezing of sur-face spins occurs [22,23].

*Corresponding author: [email protected]©KIM and Springer

1046 Young-Min Kang et al.

As just mentioned, in order to improve the LFMR effect ofCMR materials, grain boundary engineering is very crucialsince more effective spin-dependent tunneling barriers areachievable at the grain boundaries by optimizing the amountof additives or the processing conditions. In such an effort,we have carefully studied the microstructures, and the mag-netic and LFMR properties of polycrystalline LSMO filmsthat were prepared by an ex-situ solid phase crystallizationprocess [24,25] of amorphous precursor films deposited ona-SiO2/Si substrates. In addition, we were able to interpretthe measured LFMR data by applying a modified brick layermodel [26].

2. EXPERIMENTAL

La0.7Sr0.3MnO3 (LSMO) thin films were prepared usingtwo subsequent steps. The first step was to produce an amor-phous film by dc magnetron sputtering of a polycrystallineLa0.7Sr0.3MnO3 target on an a-SiO2/Si substrate at room tem-perature. The 2 inch diameter targets were prepared by theconventional solid-state reaction method using precursors ofLa2O3, SrCO3, and Mn2O3 powders of 99.9 % purity. Theamorphous SiO2 (a-SiO2) layers on Si (001) substrates wereformed by thermal oxidation at 1100 °C in a pure oxygenatmosphere. The base pressure in the vacuum chamber wasmaintained at 3 × 10–6 Torr. The dc power was 75 W. Thesecond step was to crystallize the as-deposited amorphousfilms into LSMO films by ex-situ annealing. Ex-situ anneal-ing was performed at various high temperatures rangingfrom 550 to 1000 °C for 2 h in a pure oxygen atmosphere.Samples were then slowly cooled with a cooling rate of 2 °C/min to suppress crack formation in LSMO films due to thedifference in thermal expansion coefficients between thesubstrate and the film. The film deposition rate was ~10 nm /min. Films with thicknesses ranging from 50 to 300 nm weremainly studied because crack formation was unavoidable forfilms thicker than 300 nm.

Phase analysis was performed by XRD (M18XHF-SRA,Mac Science). Microstructure was analyzed by scanning elec-tron microscopy (SEM) (JSM-6330F, JEOL) and transmissionelectron microscopy (TEM) (F20, FEI). Cation composi-tions of the as-deposited amorphous films were analyzed byInductively Coupled Plasma - Atomic Emission Spectrome-ter (ICP-AES). Compositional analysis for the annealedfilms was performed by Auger Electron Spectroscopy (AES)and STEM (Scanning Transmission Electron Microscopy)with EDS (Energy Dispersive X-ray Spectroscopy) and EELS(Electron Energy Loss Spectroscopy). Temperature and fielddependencies of magnetization were measured with a SQUIDmagnetometer (MPMS XL5, Quantum Design). The M-Tcurves were measured with a zero field-cooled warming pro-cedure under an applied field of 1 kOe, parallel to the filmplane. For the transport measurements, films with the dimen-

sions of 10.0 mm × 1.0 mm were used. Measurements of theLFMR properties were performed using the standard fourprobe method within the SQUID magnetometer. Both theelectrical current and the magnetic field were applied alongthe long dimension of the samples. Curie temperature of thefilms was determined by an extrapolation of the ferromag-netic magnetization curves to zero value. Average grain sizeand grain connectivity were evaluated by a random interceptmethod applicable to SEM micrographs.

3. RESULTS AND DISCUSSION

In order to produce as-deposited amorphous films forwhich the composition is close to that of the bulk target, i.e.,La0.7Sr0.3MnO3, we systematically varied sputtering gas (pureAr) pressure, P and oxygen content in the sputtering gas, R(%) = [O2] / [Ar+O2]100%. Dependence of the relative cat-ion concentrations of the as-deposited amorphous films on Pand R are shown in Fig. 1(a) and (b), respectively. As a ref-erence, the relative cation ratios of the bulk target are repre-sented as dotted lines. Curie temperatures (TC) of the filmspost-annealed at 900 °C for 2 h are also presented in Fig. 1.One can see that the cation ratios of amorphous films depos-

Fig. 1. Dependence of relative atomic concentration for as-depositedLSMO films on sputtering gas (pure Ar) pressure (a), and content ofoxygen in the total processing gas of P=200 mtorr, R = [O2] / ([Ar] +[O2])100 % (b). The relative atomic concentration of target is shownwith dotted lines. Curie temperature is presented at right axis for theeach film, annealed at 900 °C.

Low Field Magnetoresistance Properties of (La0.75Sr0.25)1.05Mn0.95O3 Polycrystalline Thin Films on a-SiO2/Si Substrates Prepared 1047

ited by dc magnetron sputtering are very sensitive to P andR. On the basis of these results, the sputtering gas conditionnecessary to prepare amorphous precursor films for furtherstudy was determined as P = 200 mT of a pure Ar atmosphere(R = 0) since the cation ratio (La : Sr : Mn = 0.79 : 0.26 : 0.95)of the as-deposited film in this condition was the closest tothat of target (La : Sr : Mn = 0.7 : 0.3 :1.0) among all as-depos-ited films; also, the highest Curie temperature (TC) was obtainedamong all post-annealed films.

Figure 2(a) shows the XRD patterns of as-deposited andpost-annealed films with ~300 nm thickness. It is obviousthat the as-deposited films are completely amorphous. Clearcrystalline LSMO peaks without any second phase peaks areobserved for the films annealed in the temperature region of600-950°C in pure oxygen atmosphere. We thus designate thecomposition of the annealed LSMO films as (La0.75Sr0.25)1.05

Mn 0.95O3 since the cation ratio of amorphous films is La : Sr: Mn = 0.79 : 0.26 : 0.95. The peak intensity of the LSMOphase increases with increasing annealing temperature (Tan)up to 850 °C, but it decreases with further increasing of Tan.Unidentified second phases begin to appear when the anneal-ing temperature is 1000 °C.

SEM micrographs of all the LSMO films are shown inFig. 3. Both plan and cross-sectional views are represented.The grain structure is observable for the LSMO films annealedat temperatures higher than 550°C. An amorphous-like phase isstill observable for the films annealed at 550 and 600 °C,shown in Fig. 3(b) and (c), respectively. With careful obser-vation of the film morphology in the cross-sectional SEMmicrographs, we can discern that crystallization occurs by arandom nucleation and growth throughout the whole film.

The average grain size of the films gradually increases withincreasing Tan from 800 to 1000 °C. While the films annealedin the temperature region of 700-850 °C show round-shapedgrains, the films annealed at temperatures higher than 900 °Cshow faceted grains. In accordance with the XRD patternsshown in Fig. 2, in which it can be seen that the secondphase appears when Tan is higher than 950 °C, the secondphases are also observed in the cross-sectional image of thefilm annealed at 1000 °C, which is manifested as a tiltedcolumnar structure grown on the top of an a-SiO2 layer [Fig.3(f)]. It is worthwhile to note that two very distinguished lay-ers are observed in the annealed films, as is marked in thecross-sectional images. The upper grain-structured layer is apolycrystalline LSMO and the lower layer is an intermixedphase, located between LSMO and SiO2. The thickness ofthe intermixed layer increases with Tan. A similar inter-diffu-

Fig. 2. (a) X-ray diffraction patterns of as-deposited and annealed at550-1000 C films with 300 nm thickness and (b) 50 nm-thick filmsannealed at 900 °C. Asterisks mark the peaks of unidentified secondphases.

Fig. 3. SEM plan view and cross-section images of ~300 nm-thickLSMO films: as-deposited (a), annealed at 550, 600, 700 800, 850,900, 950, and 1000 °C (b), (c), (d), (e), (f), (g), (h), and (i), respectively.


sion process between LSMO and SiO2 layers was alsoreported for the sol-gel-processed LSMO system [27]. Forthe films annealed over 800 °C, average grain size and inter-mixed layer thickness, determined by SEM analysis, are pre-sented in Tables 1 and 2. The error bar of the thickness andgrain size is about ± 5 % of the values.

The microstructures of 300, 200, 100, and 50 nm films,annealed at 900 °C for 2 h, are shown in Fig. 3(g), and Fig.4(a)-(c), respectively. As the film thickness decreases from300 to 100 nm, the grain size decreases. For simplicity, weuse the sample ID of F300-800 to represent the LSMO filmwith a thickness of ~300 nm at the crystallization tempera-ture of 800 °C. While the F300-900 (Fig. 3(g)) and F200-900[Fig. 4(a)] films show clear grain structures, the grain struc-ture of the F100-900 film becomes less clear. The morphol-ogy of the F50-900 film is much different from those ofF200-900 and F300-900, which might be attributed to acomplete formation of the intermixed layer (Fig. 4(c)). Thedifference between the grain-structured and intermixed layeris also evident in the XRD analysis. As can be seen in Fig.2(b), the F50-900 film is amorphous. The gradual decreaseof XRD peak intensity with Tan (Tan > 850 °C) (Fig. 2(a)) is

also attributable to the growth of amorphous intermixedphase with increasing Tan.

Compositional analysis was performed on the F300-900sample by both AES (Auger Electron Spectroscopy) andTEM-EDS (Energy Dispersive X-ray Spectroscopy). Figure5 shows the AES-depth profile of cations from the film topsurface to the bottom. The AES distinguishes the “pure” Sifrom the substrate and the Si from the SiO2 oxide. Figures 6(a)and (b) show the cross-sectional STEM Z-contrast images ofthe F300-900 sample and the TEM-EDS line profiles of ele-ments (from bottom to top) corresponding to the line markedin Fig. 6(a). From the depth profiles of AES (Fig. 5) andTEM-EDS (Fig. 6(b)), one can see that a severe inter-diffu-sion occurs between SiO2 and the as-deposited amorphousfilm during the ex-situ annealing process; the intermixedlayer is the glass-like Si-Sr-Mn-La-O phase and the upperlayer, with a very small amount of Si, is the granular LSMOphase. In addition, pile-up of the Sr component at the bottomof the intermixed layer is also observed.

A TEM bright field image showing an interface betweenthe intermixed layer and two adjacent LSMO grains is pre-sented in Fig. 7. Fast Fourier transformation (FFT) patterns,

Table 1. Film thickness (tf), intermixed layer thickness (tim), average grain size (d), Curie temperature (TC), resistivity (ρ) (calculated with excluding intermixed layer), grain connectivity ( f ), grain boundary areal resistance (rAGB), and LFMR at 100 K and 300 K for LSMO films

annealed at 800-1000 °C

Films parameters and propertiesFilms

F300-800 F300-850 F300-900 F300-950 F300-1000tf (nm) 280 273 300 287 294tim (nm) 40 67 100 163 210d (nm) 43.0 68.5 83.3 94.2 117.2TC (K) 275 355 363 366 370

ρ (Ωcm) at 100 K − 0.173 0.255 0.229 0.306ρ (Ωcm) at 300 K − 0.298 0.445 0.395 0.318

f − 0.960 0.965 0.991 0.942rGB (µΩ·cm2) at 100 K − 1.14 2.05 2.14 3.38LFMR (%) at 100 K − 13.8 11.8 12.2 7.3LFMR (%) at 300 K − 0.72 0.82 0.81 0.60

Table 2. Curie temperature (TC), intermixed layer thickness(tim), average grain size (d), resistivity excluding intermixed layer (ρ), grain connectivity ( f ), areal grain boundary resistance (rGB), and LFMR at 100 K and 300 K of LSMO thin films annealed at 900 °C with thickness of 50, 100,

200, and 300 nm

Films parameters and propertiesFilms

F50-900 F100-900 F200-900 F300-900tim (nm) 50 64 94 100d (nm) − 45.7 58.8 83.3TC (K) − 364 370 366

ρ (Ωcm) at 100 K Non -conductive 0.856 0.481 0.255ρ (Ωcm) at 300 K Non -conductive 0.838 0.763 0.445

f − 0.954 0.962 0.965rGB (µΩ·cm2) at 100 K − 3.73 2.72 2.05LFMR (%) at 100 K − 12.6 15.8 11.8LFMR (%) at 300 K − 0.60 0.97 0.82


marked with a square box in the image, are also shown. Wecan observe clear FFT patterns of crystalline phases withinthe two LSMO grains (Fig. 7(b) and (c)) and that of theamorphous phase in the intermixed layer (Fig. 7(e)), whichis in accordance with the XRD analysis results. The grainboundary region shown in the bright field image appears tobe an amorphous-like phase. Its FFT pattern also shows apattern of amorphous-rich phase with small amount of crys-talline phase, detected as the dim dot spots (Fig. 7(d)). Theseanalysis results support the idea that the intermixed amor-phous phase might diffuse along the grain boundary of theLSMO during the annealing process at high temperature.

The temperature and field dependencies of the magnetiza-tion of LSMO films with various thicknesses, annealed at900 °C, are shown in Fig. 8. The F50-900 sample shows avery broad ferromagnetic-paramagnetic (FM-PM) phasetransition and a very low value of saturation magnetization,MS. The low MS value can be attributed to the very weakmagnetism of the amorphous intermixed layer. As the film

thickness increases, TC increases and its transition becomesshaper because ferromagnetic LSMO is formed on the inter-mixed layer and the relative volume of the magneticallydead intermixed layer decreases. For the thick film of F900-900, the TC decreases and the FM-PM transition becomesbroad again. This is probably due to an oxygen deficiencyfor thick films at the given annealing temperature and time.Temperature and field dependencies of the magnetization of300 nm films on annealing temperature are shown in Fig. 9.When Tan is higher than 800 °C, TC increases above roomtemperature and thus becomes close to the value of the bulksample (TC ~370 K). The magnetization values, M (emu/cc),of the films annealed at 850−1000 °C, obtained by excludingthe intermixed layer thickness for its calculation, are similarto one another (the Fig. 9(a)–(c)), while the saturation mag-netization (MS), calculated by including the whole LSMOfilm thickness, decreases with Tan because the relative vol-ume of the magnetically dead intermixed layer increases(Fig. 9(d), (e)).

The low field magnetoresistance (LFMR) properties of thefilms are presented in Figs. 10-11. The MR (%) was definedas (RH – RH=1200 Oe ) / RH = 1200 Oe × 100 % and (RH – RH=500 Oe ) / RH

Fig. 4. SEM plan view and cross-section images of LSMO filmsannealed at 900 °C with thickness of 200 (a), 100 (b), and 50 nm (c).

Fig. 5. AES thickness profiles of cations concentration for F300-900sample.

Fig. 6. (a) Cross-sectional STEM Z-contrast image of F300-900 sam-ple and (b) TEM-EDS line profiles of the elements of the film, corre-sponding to a line marked in Fig. (a)

Fig. 7. (a) TEM HR image of an interface between the intermixedlayer and two adjacent LSMO grains, and (b)–(d) Fast Fourier Trans-formation (FFT) patterns of the marked square area in A, B, C, D,respectively.


= 500 Oe × 100 % for the measurements performed at 100 and300 K, respectively. The LFMR properties of the LSMOfilms with 300 nm thickness, annealed at 850–1000 °C, areshown in Fig. 10. The MR curves exhibit hysteresis with twopeaks, corresponding to increasing and decreasing field,because the MR behavior is closely related to the magnetiza-tion process caused by the magnetic domain motion. Thus,the MR peak positions show a good agreement with thecoercive field (HC) (see Fig. 9). For the F300-850 film, rela-tively high LFMR values of 13.8 and 0.72 % are obtained at100 and 300 K, respectively. The dependence of the LFMRon film thickness can be seen in Fig. 11. Among the LSMOfilms studied here, the highest MR values of 15.8 % (~16 %)at 100 K and 1.2 kOe, and 0.97 % (~1.0 %) at 300 K and 0.5kOe, respectively, are obtained from the F200-900 sample.These LFMR values are also relatively high compared withthose of polycrystalline LSMO-based films previously reported[12,14,28-30] (see Table 3), although it was possible toobtain much better LFMR performance from the bridge-pat-terned LSMO films with artificial grain boundaries [31]. TheMR value of the F100-900 sample is higher than that of theF300-900 sample at 100 K but becomes lower at 300 K. Thereason for this is that the F100-900 sample has weaker ferro-magnetism than the thicker films at 300 K [Fig. 8(a)]. Actu-ally, the ratio of magnetization value, obtained at 300 K and100 K, M300K/M100K, is about 0.7 for the F200-900 and F300-

900 samples, but it is about 0.4 for the F100-900 sample.The maintenance of ferromagnetism at room temperature isstrongly required for MR sensor applications because theLFMR is based on a spin polarized current in a ferromag-netic half metal. Resistivity and LFMR of the films, mea-sured at 100 K and 300 K, are presented in Tables 1 and 2. Itis important to note that the F50-900 sample is nonconduc-tive (at least, the resistivity is over 300 Ωcm). The high resis-tivity of this film is most probably caused by its amorphousintermixed layer, which was identified by SEM [Fig. 4(c)]and XRD [Fig. 2(b)]. The thickness of the intermixed layerwas excluded when deducing the resistivity of the topLSMO layer because the contribution of the intermixed layerto the film conductivity would be negligibly small.

To describe the electrical transport property of granularmaterials, the brick layer model (BLM) has been widelyused [26,32-37]. According to the BLM, samples can be rep-resented by the cubic grains of size d, connected through thegrain boundaries (GBs) with thickness of δ. For the validityof the BLM, δ should be much less than d, and resistivity of

Fig. 8. (a) Temperature dependence of magnetization of LSMO filmsannealed at 900 °C with thickness of 50-900 nm, and magnetic hyster-esis loops of the films annealed at 900 °C measured at 100 K (b). M(emu/cc) was calculated with including whole LSMO films thickness.

Fig. 9. (a) Temperature dependencies of magnetization, obtained at H= 1.0 kOe, of 300nm-thick films annealed at 700, 800, 850, 900, 950,1000 °C, and magnetic hysteresis loops of the films annealed at 850,900, 950, 1000 °C and measured at 100 K (b), (d) and 300 K (c),(e). InFig (b), (c), M (emu/cc) was calculated with excluding the intermixedlayer thickness.


GB, ρGB, should be much higher than that of the grain, ρG. Inthis case, the current flows only across the GBs (series path),but not along the GBs (parallel or side-wall path). Then, thesample resistivity can be written as

ρ = ρG + ρGB · (δ/d). (1)

Although this series-BLM model has many limitations[34-37] in describing a real material system, it is quite wellapplicable to the present ferromagnetic manganite filmssince the average grain size (70-120 nm) of our films is largeenough to yield strong metallic ferromagnetism within grainsand grains have relatively narrow distribution in their shapeand size.

The grain resistivity at 100 K, reported for the single crys-tals and epitaxial LSMO thin films, is ~10−4 Ωcm [4,5], whileit is ~10−1 Ωcm [30,38] for polycrystalline LSMO thin films.Thus, it is obvious that the resistivity of the LSMO films ismainly due to the GB contribution. Introducing the grainboundary areal resistance, rGB =ρGB·δ, the sample resistivityin Eq. (1) can be rewritten as

ρ = ρG + rGB / d. (2)

The second summand is much higher than the first one andEq. 2 can be rewritten as

ρ ≅ rGB / d, (3)

and grain boundary areal resistance can be expressed as

rGB ≅ ρ · d, (4)

which shows the properties of the separated grain bound-ary. Also, taking into account the connectivity ( f ) betweenthe grains, the GB areal resistance can be presented as

rGB ≅ f · ρ · d. (5)

Using the measured resistivity of sample ρ, and grain sizeand grain connectivity (obtained by random intercepts meth-ods, applied to the film SEM images), and Eq. (5), the GBareal resistance, rGB, was evaluated as listed in Tables 1 and2. The film resistivity, obtained at 100 K, increases (exclud-ing the F300-950 sample) with increasing Tan. The deviationfrom the increasing tendency, shown by the F300-950 film,is probably caused by its grain connectivity, which is higherthan that in other films. In reality, the rGB, calculated by tak-ing into account the grain connectivity ( f) and grain size (d),increases continuously with Tan. In this system, an increase inrGB can be caused by the increased disorder of spin align-ment due to the compositional and structural imperfections

Fig. 10. Low field magnetoresistance of LSMO films, annealed at 850-1000 °C, measured with switching field at 100 K (a) and 300 K (b).

Fig. 11. Low field magnetoresistance of LSMO films, annealed at900 °C with thickness of 100, 200, 300 nm, measured with switchingfield at 100 K (a) and 300 K (b).


or to the increase of the thickness of the disordered layers atthe grain surface/boundary region. Taking into account thathigher values of rGB are obtained for samples with higher rel-ative thicknesses of the intermixed layer (see Tables 1 and 2)and the results of the TEM analysis performed on the grainboundary region (Fig. 7), it can be understood that the increasein both ρ and rGB of the films is mainly caused by the diffu-sion of the amorphous intermixed compound (Si-La-Sr-Mn-O) from the intermixed layer to the GBs. The LFMR valueof the ~300 nm thick films, measured at 100 K, decreaseswith increasing rGB (Table 1). For the films annealed at 900 °Cwith various thicknesses, the highest LFMR was achievedfrom the F200-900 sample with intermediate rGB (Table 2). Itis difficult to establish a direct relationship between theLFMR value and rGB for these films because no clear ten-dency of LFMR variation with rGB can be found in Tables 1and 2. However, under the consideration that LFMR iscaused by spin-dependent tunneling at the GBs, the presentresults can be interpreted in the following way: the change ofGB properties through the diffusion of intermixed elementseither makes the GBs spin flipping centers, degrading theLFMR effect, or makes them effective spin-dependent scat-tering centers, causing LFMR enhancement, depending onthe processing conditions by which the GB tunnel barrier isformed. This phenomenon is very similar to that in previousresults reported for the composite systems between CMRmanganites and insulating materials [11,15-17], in which thehighest LFMR was obtainable from the composite with aproper amount of insulting barrier because the enhancementof LFMR is due to the tuning of GBs, allowing them to playroles as the most effective spin dependent tunneling barriers.It is also reported that the LSMO-5 % In2O3 composite filmsshow either enhanced or degraded LFMR effect dependingon the annealing temperature [14]. In this respect, it can beseen that both the degree of GB amorphization, due the inter-diffusion effect, and annealing temperature are closelyrelated to the modification of the GB property on which the

LFMR depends. Therefore, the enhanced LFMR of LSMOfilms in this study is attributable to the well-directed control-ling of the GB properties.

4. CONCLUSION

The enhanced LFMR properties of polycrystalline LSMOfilms with the composition of (La0.75Sr0.25)1.05Mn0.95O3 can beobtained on a-SiO2/Si substrates by ex-situ solid phase crys-tallization of as-sputtered amorphous films. Microstructure,Curie temperature, magnetization and LFMR properties ofthe films are tunable by controlling the annealing tempera-ture and the film thickness. An intermixed layer between theLSMO and the a-SiO2 layers is observed and its thickness isfound to increase with annealing temperature. Analyses ofthe microstructure, and the magnetic and low field magne-toresistance (LFMR) properties of the films, combined withthe brick layer model, show that the grain boundary proper-ties of LSMO films can be positively modified by the diffu-sion of intermixed elements from the inter-reacted layerbetween the as-deposited amorphous LSMO and the a-SiO2

layers, and hence the grain boundaries of the LSMO film canbe turned into more effective spin-dependent tunneling bar-riers, leading to an enhancement of LFMR. Further improve-ment in LFMR may be possible by optimizing the processingconditions and film thickness.

ACKNOWLEDGMENT

This work was supported by the National Research Foun-dation of Korea(NRF) grant funded by the Korea govern-ment(MEST) (No. 20090076203).

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Table 3. Low field magnetoresistance (LFMR) performances of selected LSMO films prepared by different methods Composition Substrate Fabrication method T (K) H (Oe) LFMR (%) Ref.

La0.7Sr0.3MnO3 Si(111) PLD 86 3000 8 12La0.7Sr0.3MnO3+10 % Al2O3 Si(111) PLD 86 3000 15 12

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La0.7Sr0.3MnO3 SiO2/Si(001) dc-sputter 77 1200 12 29La2/3Sr1/3MnO3-d Si(001) dc-sputter 100 1200 7 30

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Low field magnetoresistance properties of (La0.75Sr0.25)1.05Mn0.95O3 polycrystalline thin films on...

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