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Engineered Interface of Magnetic Oxides

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Science  30 Jul 2004:
Vol. 305, Issue 5684, pp. 646-648
DOI: 10.1126/science.1098867

Abstract

Interface-selective probing of magnetism is a key issue for the design and realization of spin-electronic junction devices. Here, magnetization-induced second-harmonic generation was used to probe the local magnetic properties at the interface of the perovskite ferromagnet La0.6Sr0.4MnO3 with nonmagnetic insulating layers, as used in spin-tunnel junctions. We show that by grading the doping profile on an atomic scale at the interface, robust ferromagnetism can be realized around room temperature. The results should lead to improvements in the performance of spin-tunnel junctions.

In the development of semiconductor technology, a major role was played by the elucidation of electronic states at heterointerfaces such as positive-negative (p-n) junctions and metal/semiconductor interfaces. In the p-n junction, for example, a depletion layer and built-in electric field are formed at the interface so as to unify the Fermi level all over the junction. Analogously, transition-metal-oxide compounds, as emergent electronics materials with correlated electrons, often exhibit nontrivial and potentially useful electronic properties at heterointerfaces. In most cases, however, such interface properties cannot be explained in terms of conventional band pictures (1). Moreover, the charge state in the strongly correlated electron oxides is closely coupled with the spin and orbital degrees of freedom. Therefore, magnetic oxide interfaces provide a challenging arena to explore new multifunctional materials (2).

Interface magnetism is also important for the development of devices with strongly correlated electron oxides. For example, the perfectly spinpolarized ferromagnet La1-xSrxMnO3 (LSMO) could be one of the best candidates for magnetic random-access memory if we could fully use its potential in spin-tunnel junctions. However, in the actual junctions—for example, LSMO/SrTiO3 (STO)/LSMO—the tunnel magnetoresistance (TMR) has been much smaller than that expected from the half-metallic nature of LSMO, and the response has been diminished at temperatures far below Tc (3, 4). The possible origin of the decreased MR is severe deterioration of the ferromagnetism locally occurring near the LSMO/STO interfacial region (57). Therefore, selective and quantitative evaluation of interfacial spin state may provide a key to realizing high-performance magnetic devices.

Nonlinear magneto-optical effects can probe the interface magnetism (8, 9), as has recently been demonstrated for oxide “tricolor” superlattices, where ferromagnetic LSMO layers and insulating STO and LaAlO3 (LAO) layers were accumulated one by one in an ABCABC... structure. The STO/LSMO/LAO superlattice exhibits very large magnetization-induced secondharmonic generation (MSHG) and a resultant nonlinear magneto-optical Kerr effect (NOMOKE) (10, 11). This “tricolor” spin superlattice is regarded as an artificially constructed noncentrosymmetric ferromagnet, where symmetries of space inversion and time reversal are both broken simultaneously. The generic feature of such a noncentrosymmetric ferromagnet should be extended to various heterointerfaces of strongly correlated electron oxides.

We show that such an interface magnetism can indeed be probed by MSHG from single heterojunctions between LSMO and STO or LAO, and we have found a robust ferromagnetic interface of LSMO. It can be obtained by modulating the interface doping profile on an atomic scale, which may open up a new path to LSMO-based spin-injection devices such as TMR junctions operating at room temperature.

The schematic side view of the heterostructures used in the present study is depicted in Fig. 1A. For detecting the interface magnetization at a single magnetic interface, we prepared bilayer films composed of an insulating cap layer (STO or LAO) with a thickness of 5 unit cells (u.c.) (i.e., 2 nm) and a magnetic LSMO bottom-layer grown on STO(001) substrates (Fig. 1A, left). The LSMO layer is as thick as 300 u.c (120 nm). Because the optical absorption coefficients in LSMO for incident beam (λ = 800 nm) and SH light (λ = 400 nm) are about 1 × 105 cm–1, the contribution from the bottom film/ substrate interface is negligibly small (<1%). All the films were fabricated by pulsed laser deposition (PLD) in a layer-by-layer growth mode while observations were made of the intensity oscillation of the specular spot in reflection high-energy electron diffraction (RHEED) (Fig. 1B). The resultant film surfaces are atomically flat with 200-nm-wide terraces and 0.4-nm-high steps. We also note that there should be no detectable interface roughness or disorder, judging from the transition electron microscopy image and x-ray diffraction for the “tricolor” superlattice grown in a similar way (12). Therefore, we may conclude that exact control of interface atomic stacking can be realized. Four-circle x-ray diffraction measurements indicate that the in-plane lattice constants are identical with that of the substrate (STO), ensuring coherent epitaxy throughout the bilayer film. Magnetization of the bilayer film (Mfilm) was measured with a superconducting quantum interference device (SQUID) magnetometer, indicating in-plane magnetic anisotropy, well-defined Tc of 340 K, and saturated magnetization of 3.6 μB / Mn for all the samples.

Fig. 1.

(A) Atomic stacking sequences for perovskite-type magnetic oxide interfaces. (Left) LSMO film capped with a 5-u.c. (2-nm) layer of LaAlO3 (LAO) or SrTiO3 (STO). (Right) A 2-u.c. layer of undoped LMO is inserted between STO and LSMO layers to grade the hole concentration in the manganite layer near the interface. In these films, the top surface is terminated by the MnO2 or BO2 (B = Ti or Al) layer (18). The arrows depict possible arrangements of magnetic moments at the respective MnO2 layers, showing considerable spin canting near the interface (left) and robust ferromagnetism (right). (B) (Left) RHEED oscillation during the growth of the STO/LMO/LSMO heterostruture. (Right) Atomic force microscopy image for the top surface of STO/LMO/LSMO. Atomically flat terraces with the single-u.c. steps (0.4 nm high) are clearly seen.

Figure 2A illustrates the experimental configuration for the MSHG measurements (13). The external or built-in fields induce magnetization (M), electric polarization (P), and toroidal moment (TP × M), which gives MSHG (Fig. 2A, bottom). The SH light intensity (I), measured in a magnetic field (H) of ± 0.05 T, was plotted against the analyzer angle (θ), as shown in Fig. 2C. MSHG corresponds to the s(y)-polarized component (θ= 90°) in the SH light. A phase shift in the I-θ curves for ± H gives twice the NOMOKE angle, from which we can deduce the intensity of MSHG. We define the square root of the MSH intensity as the interface magnetization (Mint), which is plotted against T in Fig. 2B. In the upper panel, the Mint-T plot is normalized with Mint(50 K) and compared with the Mfilm-T curve.

Fig. 2.

(A) (Top) Schematic of experimental configuration for the measurement of MSHG. Polarization of incident light and applied magnetic field (H) are along y and x, respectively. The reflected SH light undergoes Kerr rotation. (Bottom) Configuration of various electric and magnetic moments. P (//z) stands for electric polarization originating from noncentrosymmetric atomic stacking along the z direction. M (//x) shows magnetization of LSMO along the film plane induced by a magnetic field H (//x). T (//y) indicates a toroidal moment, which is generated as T = P × M. (B) (Bottom) Temperature (T) dependence of magnetization at the interface (Mint), defined as the square root of MSHG intensity (s-in s-out SH light). The applied magnetic field is 0.05 T. Red, blue, and green symbols indicate the data for STO/LSMO, LAO/LSMO, and STO/LMO/LSMO interfaces, respectively. (Top) Mint(T) normalized with Mint (50 K), and T dependence of magnetization of the films (Mfilm) measured at 0.05 T with a SQUID magnetometer (a black line). (C) SH intensities plotted against analyzer angle (θ) at 50 K, 150 K, and 250 K for +H and –H. θ denotes 0° and 90° for p- and s-polarized SH light, respectively. The θ dependence can be fitted to the formula |Pz cosθ ± Py sinθ|2. Phase shift induced by the H reversal corresponds to twice the NOMOKE angle, which gives the magnitude of Mint.

The MSHG data shown in Fig. 2C indicate a large variation of magnetic properties between STO/LSMO and LAO/LSMO interfaces. The STO/LSMO interface hardly shows MSHG or NOMOKE even at 50 K, which directly demonstrates the existence of a dead layer at the interface and explains the inferior TMR response found in LSMO/STO/LSMO junctions (5). At the LAO/LSMO interface, by contrast, large MSHG is observed, and the ferromagnetism survives even at 250 K, although the onset T of the MSHG signal as the interface Tc is not so high or well defined as the Tc of the bulky thick films. The present observation is consistent with the finding of the studies on transport properties of the related superlattices that the LAO and STO form robust and susceptible interfaces with LSMO, respectively (10, 11).

Figure 3 exhibits the magnetic field dependence of Mint as deduced by the square root of MSHG intensity, measured at T = 50 K. This “interface M-H curve” enables us to decompose Mint into the spontaneous magnetization and field-induced part. To clearly distinguish the two components, Mint is fitted with the spontaneous magnetization plus the H-linear function and normalized with Mint(7 T) (Fig. 3, top). The magnetic field of 7 T increases the spontaneous magnetization by 200% and 70% for STO/LSMO and LAO/LSMO interfaces, respectively. This means that the antiferromagnetic spin canting occurs at the STO/LSMO interface, while the ferromagnetic spin arrangement is much less canted by LAO.

Fig. 3.

Magnetic field (H) dependence of the square root of MSHG intensity, representing the interface magnetism (Mint), measured at 50 K. Mint (bottom) and normalized Mint [Mint(H)/Mint(7 T)] (top). STO/LSMO, LAO/LSMO, STO/LMO/LSMO interfaces are displayed with red, blue, and green symbols, respectively. The solid lines are the fit, assuming spontaneous magnetization (intersection to the vertical axis) and H linear increase in Mint from the canted states by an external field. A black line in the top panel shows the H dependence of Mfilm at 5 K.

In bulk crystals, LSMO (x = 0.4) has the highest Tc of 370 K among the La1-xSrxMnO3 series (14), but the LSMO (x = 0.4) film grown on STO is on the verge of an A-type (layered) antiferromagnet with reduced Tc(340 K) as a result of the tensile strain (15). This instability toward the antiferromagnetic phase explains why the ferromagnetic spin ordering in LSMO (x = 0.4) is so easily reduced at the interface. Moreover, STO produces a dead layer in the adjacent LSMO more aggressively than does LAO, because the STO layer works as a hole-donating layer. It has been considered that the valence-mismatched interface composed of the stacking sequence -TiO2-SrO-MnO2-La0.6Sr0.4O-induces the charge transfer and that the overdoped LSMO is dominated by the antiferromagnetic spin canting (57, 10, 16). Therefore, the underdoped LSMO (preferably x < 0.3) layer may enable us not only to stabilize the ferromagnetism but also to compensate the charge transfer from STO at the interface. However, both high Tc and conductivity are lost by decreasing x, where we confront a dilemma.

To enhance the interface magnetization without sacrificing the ferromagnetic and metallic characters in the film, we propose the compositionally graded LSMO interfaces, where the bulk LSMO electrode has the optimal doping level of x = 0.4, and x decreases gradually to 0 toward the insulating layer. This idea bears some analogy to the procedure to improve the superconductivity grain boundary junction (17), in which partial Ca substitution on Y sites in YBa2Cu3O7-x enhances the critical supercurrent density (Jc), but at the same time suppresses Tc. In our study, the doping profile was controlled on an atomic scale. The 2-u.c. layer of LaMnO3 (LMO) was inserted as a locally underdoped layer between STO (5 u.c.) and LSMO (x = 0.4; 300 u.c.) layers (Fig. 1A, right). We anticipated that the LMO layer, which is originally an antiferromagnet with TN = 140 K, should be hole-donated by LSMO(0.4) and STO layers and hence can compensate for the interface effects. This compositionally modulated interface shows large MSHG and NOMOKE, which exceed by far the STO/LSMO direct interface and are comparable to those of the LAO/LSMO interface (Fig. 2B). In accord with this observation, the spin-tunnel junction equipped with such an atomically engineered interface, [LSMO(0.4)/LMO (2 u.c.)]/STO (2 nm)/[LMO (2 u.c.)/LSMO(0.4)], shows an improved (comparable) device performance as compared with the direct-interface junction, LSMO/STO/LSMO (LSMO/LAO/LSMO) (13). As shown in Fig. 3, Mint for this STO/LMO/LSMO interface is almost H-independent above 0.1 T, thus confirming that the nearly bulk-like ferromagnetism is maintained even in the vicinity of STO. The normalized Mint(T) value of STO/LMO/LSMO reaches 0.5 at 250 K, which is close to a value (0.7) for normalized magnetization of bulk film at the same temperature and is much larger than the value (<0.2) for the direct interface of STO/LSMO (x = 0.4).

Supporting Online Material

www.sciencemag.org/cgi/content/full/305/5684/646/DC1

Materials and Methods

Fig. S1

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