Toward All-Oxide Magnetic Tunnel Junctions: Epitaxial Growth of

Article pubs.acs.org/crystal

Toward All-Oxide Magnetic Tunnel Junctions: Epitaxial Growth of SrRuO3/CoFe2O4/La2/3Sr1/3MnO3 Trilayers Ana M. Sánchez,*,† Laura Ä kas̈ lompolo,‡ Qi Hang Qin,‡ and Sebastiaan van Dijken*,‡ †

Department of Physics, University of Warwick, Coventry CV4 7AL, U.K. NanoSpin, Department of Applied Physics, Aalto University School of Science, P.O. Box 15100, FI-00076 Aalto, Finland

‡

ABSTRACT: Epitaxial growth of SrRuO3/CoFe2O4/La2/3Sr1/3MnO3 trilayers on SrTiO3 (001) substrates has been successfully achieved using pulsed laser deposition. This trilayer configuration, which consists of two conducting ferromagnetic oxides separated by a thin insulating ferrite film, is a promising candidate for all-oxide magnetic tunnel junctions. Structural analyses carried out using transmission electron microscopy and X-ray diffraction demonstrate a remarkable continuation of the in-plane and out-plane crystallographic relations across the entire structure. Magnetic measurements on SrRuO3/CoFe2O4/ La2/3Sr1/3MnO3 trilayers indicate independent magnetic switching in an external magnetic field, which is one of the prerequisites for large tunneling magnetoresistance effects.

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INTRODUCTION Quantum mechanical tunneling of electrons through thin insulating barriers forms the basis of many solid-state devices, including resonant tunneling diodes, resonant tunneling transistors, superconducting quantum interference devices (SQUIDs), and magnetic tunnel junctions. Epitaxial growth of single-crystal ferromagnetic, ferroelectric, and multiferroic oxide layers may enable the fabrication of all-oxide tunneling structures with novel functionalities.1 The integration of complex oxides into multilayer structures is, however, challenging, and obstacles to epitaxial growth including large lattice mismatch and differences in structural symmetry need to be addressed. Here, we report on the growth of oxide trilayers that combine three classes of magnetic oxides and hold a potential for all-oxide magnetic tunnel junctions with efficient spin filtering properties. The structures consist of a mixed valence manganite La2/3Sr1/3MnO3 (LSMO) at the bottom, a conductive ferromagnetic oxide SrRuO3 (SRO) at the top, and an insulating ferrite CoFe2O4 (CFO) in the middle. In the envisioned all-oxide magnetic tunnel junctions, the LSMO and SRO layers are used as magnetic electrodes and a thin CFO film is utilized as spin-filtering tunnel barrier. The mixed valence manganite LSMO has been intensively studied because it exhibits the colossal magnetoresistance effect.2−4 Moreover, theoretical calculations and experiments have shown that the LSMO conduction band is fully polarized at the Fermi level.5−7 This half-metallic nature in combination with abrupt magnetic switching characteristics in small magnetic fields makes LSMO a strong candidate for exploratory studies on magnetotransport phenomena in magnetic oxide heterostructures. The Curie temperature of bulk LSMO is ∼370 K. SRO is a conductive ferromagnetic oxide with high chemical stability, low resistivity, and good lattice match to other perovskite oxides, and, hence, it has become a popular © 2012 American Chemical Society

material in oxide electronic devices. In particular, SRO films have been used as a thin barrier layer in superconducting Josephson junctions,8 as electrodes in ferroelectric capacitors,9,10 and as a source of spin-polarized current in magnetic tunnel junctions.11,12 SRO exhibits negative spin polarization, a relatively low Curie temperature of ∼150 K, a large uniaxial magnetocrystalline anisotropy, and a large coercivity. Finally, insulating CFO is a hard magnetic material with a high degree of magnetic anisotropy and magnetostriction, and a coercive field of several tenths of Teslas.13 It has been investigated for several applications including magnetostrictive sensors14 and microwave devices,15 and has also been considered as a important component for multiferroic composites.16 The unique magnetic properties of CFO can be understood from its electronic and structural configuration in which the divalent cobalt and trivalent iron cations are located on the octahedral and tetrahedral sites, respectively (inverse spinel structure). The Curie temperature of CFO is ∼790 K.17 In magnetic tunnel junctions, exchange splitting of the conduction band in CFO provides distinctive tunnel barrier heights for majority and minority spin electrons. In combination with a ferromagnetic electrode, this spin-filtering property produces a tunneling magnetoresistance effect.18−21 While the successful integration of various magnetic oxides in trilayer structures opens the door to the design of new magnetic and magnetotransport properties, their growth is challenged by significant differences in lattice structure and symmetry. Bulk LSMO possesses a rhombohedral (R3̅c) structure at room temperature with a corresponding (nonprimitive) pseudocubic lattice parameter of aLSMO = 3.870 Å Received: October 26, 2011 Revised: December 21, 2011 Published: January 4, 2012 954

dx.doi.org/10.1021/cg201418q | Cryst. Growth Des. 2012, 12, 954−959

Crystal Growth & Design

Article

and an angle between axes of α = 89.74°. The unit cell of CFO has a face-centered cubic structure (Fd3m̅ ) with a lattice parameter of aCFO = 8.39 Å. Finally, SRO crystallizes in the GdFeO3-type orthorhombic distorted perovskite structure (Pnma) with lattice constants of aSRO = 5.54 Å, bSRO = 5.57 Å, and cSRO = 7.86 Å. The structure of SRO can be described as slightly distorted peudo-cubic (aSRO = 3.93 Å)22 since the orthorhombic distortion is close to unity and the (110) and (002) plane spacings are degenerate. In this work, we analyze the crystal structure and magnetic properties of SRO/CFO/LSMO trilayers on SrTiO3 (STO) substrates with a (001) crystal orientation. STO possesses an undistorted cubic perovskite structure (Pm3̅m) and aSTO = 3.905 Å. Hence, the STO crystal lattice matches well with the bottom LSMO and top SRO layers, but the mismatch between CFO and STO is significant ((0.5aCFO − aSTO)/aSTO = 7.4%).

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Figure 1. XRD θ − 2θ scans for single, bilayer and trilayer oxide structures on STO. The thickness of the single films is 60, 40, and 70 nm for LSMO, CFO, and SRO, respectively. In the bilayer and trilayer structures the LSMO and SRO layers are 30 and 20 nm thick. The vertical dashed lines indicate the position of the (002) and (004) reflections for bulk LSMO, CFO, and SRO.

EXPERIMENTAL SECTION

The trilayer heterostructures were grown by pulsed laser deposition (PLD) on STO (001). Prior to multilayer film growth, the STO substrates were first etched in buffered HF for 30 s and subsequently annealed in oxygen atmosphere at 950 °C for 1 h. This process, which is described in more detail in refs 23−25, resulted in fully TiO2terminated surfaces as confirmed by atomic force microscopy images showing regular terraces with straight and one unit cell high step edges. The three magnetic oxide layers were deposited from different targets on a carousel holder. The LSMO films were grown at 800 °C in an oxygen partial pressure of 0.5 mbar. After deposition the LSMO films were annealed at 800 °C in 0.65 mbar oxygen for 15 min. The CFO and SRO films were deposited at 650 °C in an oxygen partial pressure of 0.1 mbar. The laser (KrF excimer, 248 nm) was operated at 4 Hz and its fluence was 2.5 J/cm2. This resulted in a deposition rate of 9 nm/min, 3 nm/min, and 10 nm/min for LSMO, CFO, and SRO, respectively. For comparison, single films of LSMO, CFO, and SRO were also grown. Structural characterization was conducted using transmission electron microscopy (TEM) and X-ray diffraction (XRD). The TEM specimens were prepared using standard routes for cross-section (XTEM) and planar-view (PVTEM) geometries, taken along the ⟨110⟩STO and [001] STO directions, respectively. TEM analysis was performed using JEOL 2000FX and 2100 TEMs operating at 200 kV. Magnetic hysteresis curves were measured using a SQUID magnetometer.

To analyze the microstructure in more detail, TEM measurements were conducted on cross-section and planarview specimens. Figure 2a shows a representative cross-section TEM image of the SRO/CFO/LSMO trilayer on STO. It is clear that the structural integrity decreases toward the top of the oxide heterostructure, with occasional discontinuities observed in the uppermost SRO layer. Nanodiffraction patterns of each layer and the STO substrate are presented in the right panels of Figure 2a. The patterns were obtained using a 5 nm spot size on the cross-section specimen, a convergence angle of ∼1.5 mrad, and a camera length of 40 cm. Selected area electron diffraction (SAED) patterns were also measured on PVTEM specimens along the growth direction. Figure 2b shows a PVTEM image that was taken under bright field conditions near the [001] zone axis and the panels on the right display the corresponding SAED patterns from different sample positions. Pattern b.3 was measured on an area that contains the entire SRO/CFO/LSMO/STO multilayer structure. The perfect alignment of the diffraction spots clearly indicates that the epitaxial relations are maintained across all oxide layers. Detailed analyses of the diffraction patterns, both in XTEM and PVTEM, reveal the following epitaxial relationships:

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RESULTS AND DISCUSSION Figure 1 summarizes the XRD results for single, bilayer, and trilayer oxide structures on STO (001). The single LSMO and SRO films are strained and epitaxial with an out-of-plane lattice parameter of cLSMO = 3.847 Å and cSRO = 3.958 Å. CFO, on the other hand, is relaxed and its lattice parameter equals the CFO bulk value (cCFO = 8.39 Å). The growth of 20 nm SRO directly on top of a 30 nm LSMO film further shifts the SRO (002) reflection to smaller 2θ angles in SRO/LSMO bilayers. As the position of the LSMO (002) reflection does not change significantly, this clearly indicates that the SRO lattice relaxes toward its bulk value in the topmost layers of the 70 nm thick SRO single film. In the 20 nm SRO/ 30 nm LSMO bilayer, on the other hand, both films are epitaxial and fully strained. Insertion of a thin CFO layer between the LSMO and SRO films drastically alters the microstructure of the SRO top layer. Just 3 nm of CFO completely eliminates the SRO (002) reflection, indicating an enhancement of structural disorder. An increase of the CFO layer thickness to 10 nm does not drastically alter the X-ray diffraction pattern any further.

001 SRO 001 CFO 202̅ 1̅ LSMO 001 STO 110 SRO 110 CFO 10 1̅ 1 LSMO 110 STO Pattern b.2 in Figure 2b corresponds to a very thin area of the specimen where only CFO and SRO layers are present. In addition to strong diffraction spots, extra spots arise due to double diffraction, which are caused by the superposition of two materials with different lattice parameters. The spots are slightly elongated along arcs around the central beam, and this effect is stronger for the SRO reflections. The same phenomenon is observed more clearly in pattern b.1, which was measured on an area that only contained SRO. The elongation of the spots indicates in-plane misorientations (twists) between different subgrains in the SRO layer, which are probably caused by the large lattice mismatch between CFO and LSMO ((0.5aCFO − aLSMO)/aLSMO = 8.4%), followed by a further large mismatch between SRO and CFO ((aSRO − 0.5aCFO)/0.5aCFO = −6.3%). 955



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Figure 2. Epitaxial relationships in 20 nm SRO/10 nm CFO/30 nm LSMO trilayer on STO (001). (a) Bright field 002 STO image of the SRO/CFO/ LSMO/STO heterostructure. The right panels contain nanodiffraction patterns that correspond to the different oxide layers and the STO substrate. The diffraction patterns are rotated clockwise by about 30° with respect to the TEM image. (b) PVTEM image along the [001] direction. The right panels show experimental SAED patterns for three different sample areas: b.1 SRO only, b.2 SRO + CFO, and b.3 entire oxide heterostructure.

Figure 3. HRTEM image of (a) LSMO/STO and (b) SRO/CFO/LSMO interfaces.

A series of high resolution (HR)TEM images were obtained from the cross-section specimen to examine the interface structure, which is pivotal for magnetic tunnel junctions. Figure 3a shows a HRTEM image of LSMO/STO, indicating clear epitaxy and a sharp, flat interface. Figure 3b shows both CFO/ LSMO and SRO/CFO interfaces. Although epitaxial growth is maintained, atomic steps on the LSMO film are observed; that is, the CFO/LSMO interface is not completely flat. A subgrain structure can be seen in the CFO film, with small twists and tilts with respect to the LSMO layer below. This may be caused by the high mismatch between CFO and LSMO in conjunction with interfacial roughness. The upper SRO/CFO interface is rougher than the CFO/LSMO interface. Grooves can be seen where the CFO subgrain boundaries meet the upper interface, and variations in orientation of the SRO film are correspondingly larger than that of the underlying layers. These observations are consistent with the arc-shaped spots in patterns b.1 and b.2 of Figure 2b and the disappearance of the SRO (002) reflection in the XRD data of Figure 1. The interfacial defects were also examined by PVTEM. Figure 4a shows a bright field image, recorded under two beam conditions with a 220STO reflection, revealing a widely spaced misfit dislocation network along the [010] and [100] directions at the LSMO/STO interface. Although less common, some misfit dislocations along ⟨110⟩ are also present. This type of dislocation network is typical for the initial stages of lattice

Figure 4. Bright field 220 STO images of a PVTEM specimen revealing (a) misfit dislocation networks at the LSMO/STO interface and (b) twin domains in the LSMO layer.

relaxation in heterostructures with a small mismatch that have exceeded the Matthews-Blakeslee critical thickness (hc).26,27 There is a noticeable difference in the density of misfit dislocations along the two perpendicular ⟨001⟩ directions. This asymmetry can be attributable to a preferential alignment of the rhombohedral LSMO structure on the STO substrate or, simply, a statistical variation of the dislocation density from place to place. We note that the presence of misfit dislocations is not inconsistent with the presence of strain as seen in the XRD measurements. The misfit dislocation density observed at the LSMO/STO interface is not enough to relax the layer. In low-misfit systems, the stress is normally accommodated by a combination of elastic and plastic strains which minimize the 956



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For heteroepitaxial systems with a large lattice mismatch (|f | > 4−5%), the critical thickness becomes less than one monolayer, and as a result a geometrical misfit dislocation network relaxes most of the lattice strain. Generally, such very closely spaced arrays of misfit dislocations are not easily observed by conventional TEM. In our trilayer system, the formation of two misfit dislocation networks is anticipated, one at the bottom CFO/LSMO interface and another at the top SRO/ CFO interface. Both networks overlap in the PVTEM images and as a result complex moiré fringes are present (Figure 5). Magnetic hysteresis curves of epitaxial SRO/CFO/LSMO trilayers and single LSMO, CFO, and SRO films are compared in Figure 6. Here, we focus on measurements at two temperatures, namely, 150 and 5 K. At 150 K, the SRO film is paramagnetic and, hence, it does not contribute to the magnetic signal. At this temperature, magnetic switching in the single LSMO film is abrupt, whereas the magnetization in the single CFO film reverses much more gradually and only saturates in a large magnetic field of ∼5 T. The hysteresis curve of the oxide trilayer also contains these two components. The inset in the main panel shows the hysteresis curve of the trilayer structure after subtraction of the low-field LSMO contribution. The resemblance between the shape of this reconstructed curve and the hysteresis of the single CFO film on STO as well as their similar saturation magnetization (240 kA/m for CFO in the trilayer versus 230 kA/m for CFO on STO) clearly indicates that the LSMO bottom layer does not drastically change the magnetic properties of the CFO film. Moreover, this convincingly illustrates that the magnetization of the LSMO and CFO layers switch independently in the epitaxial oxide trilayer. This latter point is important for the realization of large magnetoresistance effects in all-oxide magnetic tunnel junctions. The magnetic hysteresis curves of the same samples at 5 K confirm these observations. At this temperature, the SRO layer is also ferromagnetic. The single SRO film on STO exhibits perpendicular magnetic anisotropy as illustrated by the closed hysteresis curve that is measured with the magnetic field along the film plane (IP) and the near-square hysteresis that is obtained for out-of-plane (OOP) applied fields. This is in agreement with earlier reports on magnetic anisotropy in this system,33,34 which might be ascribed to a distortion of the

total energy of the system,26,27 and a film thickness of ∼4 hc is necessary for significant strain relaxation.28 For LSMO/STO 4 hc ≈ 150 nm in the Matthews-Blakeslee model. However, the relaxation process in LSMO/STO may be quite different from metals and semiconductors. For this epitaxial system, other studies have found that strain relaxation only occurs above a LSMO film thickness of about 100 nm in the form of columnar growth.29 In addition to a dislocation network, a “tweed” structure is also observed in Figure 4a. This tweed structure, which is shown more clearly in Figure 4b, is formed by twin domains in the LSMO layer.30−32 The rhombohedral symmetry of the LSMO leads to the formation of a number of energetically equivalent variants (domains) which originate from the four possible alignments of the unique ⟨111⟩ axes. The twin domains, as well as the misfit dislocations, run preferentially along one of the ⟨100⟩ directions in this part of the specimen, which may indicate some interaction between them. Therefore, the LSMO/STO lattice mismatch (−0.9%) and the angular mismatch (0.26°) are accommodated by the formation of a dislocation network and geometrical twinning. PVTEM images using bright field conditions near the [001] zone axis reveal moiré fringes associated with misfit dislocations at the SRO/CFO/LSMO interfaces as illustrated in Figure 5.

Figure 5. Bright field near the [001] zone axis revealing the moiré pattern corresponding to the highly mismatch interfaces in the trilayer. The inset on the top-left corner corresponds to a magnified area clearly demonstrating the moiré fringes.

Figure 6. Magnetic hysteresis curves for single LSMO, CFO, and SRO films and a SRO/10 nm CFO/LSMO trilayer on STO (001) at 150 K (a) and 5 K (b). The insets in the main panels show the magnetic hysteresis curves of the oxide trilayer after subtraction of the low-field LSMO contribution. Note that the magnetization of single films is plotted as kA/m (magnetic moment per sample volume), while the magnetization of the trilayer is expressed in mA (magnetic moment per sample area). Hysteresis curves measured with in-plane (IP) and out-of-plane (OOP) magnetic field are indicated by solid black and red lines, respectively. 957



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orthorhombic unit cell in thin films.35,36 With the SRO layer being ferromagnetic at 5 K, the magnetization curve of the oxide trilayer now contains three components. Abrupt switching in small applied magnetic fields is again due to magnetization reversal in the LSMO layer. When this contribution is subtracted from the magnetization curve of the entire trilayer, a combination of the CFO and SRO signals remains (inset in main panel). The remanent magnetization of the reconstructed curve (280 kA/m) is only slightly larger than that of the single CFO film on STO (200 kA/m). This suggests that the SRO film in the oxide trilayer exhibits small remanence and, thus, that the perpendicular anisotropy of the single SRO film on STO is mostly preserved in the oxide trilayer. Moreover, the magnetization of the trilayer does not saturate in a 7 T field, which indicates that large magnetic fields are required to rotate the SRO magnetization into the film plane. From these results, we conclude that magnetic interlayer coupling in the epitaxial heterostructures is negligible at 5 K and that the magnetization of the three layers reverse independently from each other.

interesting discussions and Antti Hakola and Timo Kajava for assistance with the PLD system.

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CONCLUSIONS We have demonstrated that thin films of LSMO, CFO, and SRO can successfully be grown into epitaxial trilayers on STO (001) substrates. Despite the large lattice mismatch between the CFO film and the bottom LSMO layer (8.4%) and the subsequent mismatch between the lattice of the SRO top layer and CFO (−6.3%), full epitaxial relations are maintained across the entire heterostructure. The differences in crystal structure and lattice symmetry, however, give rise to the formation of several defect structures. First, a widely spaced misfit dislocation network appears at the LSMO/STO interface. Second, twin domains are observed in the LSMO layer due to the rhombohedral symmetry of the LSMO lattice. Finally, the differences in lattice parameter at both trilayer interfaces lead to dense dislocation networks. This in turn produces a subgrain structure with mosaic spread in the CFO and SRO layers and an enhancement of the interface roughness. The magnetization of the LSMO, CFO, and SRO layers are not strongly coupled in the epitaxial oxide trilayers. As a consequence, the layers switch independently in an applied magnetic field. This magnetic feature and the ability to grow epitaxial oxides trilayers open the door to the realization of all-oxide magnetic tunnel junctions.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*(A.M.S.) Phone: +44 24765 74395. E-mail: a.m.sanchez@ warwick.ac.uk. (S.v.D.) Phone: +358-50-3160969. E-mail: [email protected]. Web: http://tfy.tkk.fi/ nanospin/.

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ACKNOWLEDGMENTS A.M.S. thanks the Science City Research Alliance and the HEFCE Strategic Development Fund for funding support. The Gatan Orius digital TEM camera used in this research was funded by Birmingham Science City: Creating and Characterizing Next Generation Advanced Materials, with support from Advantage West Midlands and partly funded by the European Regional Development Fund. L.Ä. acknowledges financial support from the Väisälä Foundation. This work was partly supported by the Academy of Finland (Grant Nos. 127731 and 252301). The authors would like to thank Richard Beanland for 958



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(35) Gan, Q.; Rao, R. A.; Eom, C. B.; Wu, L.; Tsui, F. J. Appl. Phys. 1999, 85, 5297−5299. (36) Kan, D.; Shimakawa, Y. Cryst. Growth Des. 2011, 11, 5483− 5487.

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Toward All-Oxide Magnetic Tunnel Junctions: Epitaxial Growth of

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