We investigate the structural and magnetic properties of La0.7Sr0.3CrO3(LSCO)/La0.7Sr0.3MnO3 (LSMO) heterostructures grown on (001)-oriented MgO by molecular beam epitaxy. Due to the large film-substrate lattice mismatch, strain relaxation is found to occur within the first 2–3 unit cells (uc) of the film as evidenced by reflection high energy electron diffraction and high-resolution synchrotron x-ray reciprocal space mapping. We find that the presence of the LSCO spacer and capping layers leads to ferromagnetism in ultra-thin LSMO layers with thicknesses of the order of 2 uc with the magnetic easy axis oriented in the film plane. Net magnetic moments of 1.4 and 2.4 μB/Mn are measured for (2 uc LSCO/2 uc LSMO) and (2 uc LSCO/4 uc LSMO) superlattices, respectively, by superconducting quantum interference device magnetometry. The effective magnetic anisotropy of the relaxed (2 uc LSCO/4 uc LSMO) heterostructure is found to be an order of magnitude higher than bulk LSMO highlighting the critical role of interfacial interactions in tuning magnetic anisotropy at complex oxide interfaces.

Doped rare-earth manganites (A1xBxMnO3 where A is a rare-earth ion and B is an alkaline-earth ion) exhibit a wide range of interesting physical properties including tunable magnetic phases and metal–insulator transitions, colossal magnetoresistivity, and half metallicity.1,2 For bulk La1xSrxMnO3 (LSMO) with 30% Sr doping, a ferromagnetic metallic state exists below 360 K.3,4 Due to the strong coupling of the lattice, spin, and electronic degrees of freedom, the transport and magnetic properties of LSMO thin films have been tuned by epitaxial growth on closely lattice-matched single crystal substrates such as LaAlO3, SrTiO3, DyScO3, and La0.18Sr0.82)(Al0.59Ta0.41)O3 (LSAT).5–7 Epitaxial strain provides an effective route to control magnetic anisotropy (MA) in LSMO thin films with important implications for the design of novel spin-based devices.8–11 The magnetic easy axis for LSMO films under tensile strain on SrTiO3 lies in-plane along the [110] pseudocubic axis while compressively strained films on LaAlO3 exhibit a strong perpendicular magnetic anisotropy (PMA).8 As the film thickness is reduced to the dimensions on the order of a unit cell, surface and interfacial contributions to magnetic anisotropy energy are non-negligible and compete with strain-mediated contributions to magnetocrystalline anisotropy.12,13

An interface-induced magnetic anisotropy has been reported for SrIrO3/La1xSrxMnO3 superlattices, where an increase in the perpendicular magnetic anisotropy (PMA) is correlated with an increase in oxygen octahedral rotations about an in-plane axis as a function of the Sr content.14 A PMA has also been reported for tensile-strained La1xSrxCoO3/La1xSrxMnO3 bilayers. For the relaxed La0.7Sr0.3CrO3 (LSCO)/LSMO heterostructures, the in-plane and out-of-plane rotations are expected to be bulk-like and equivalent along the orthogonal pseudo-cubic axes. Recent results indicate that interface-induced preferential occupations of the Mn 3d orbitals can result in spin-reorientation transitions. Song et al. show that an interplay between ferromagnetic interfacial Mn–O–Co exchange, charge transfer, and a preferential occupation of out-of-plane pointing Mn 3d z2r2 orbitals are associated with a PMA in tensiley strained LaSrCoO3/LaSrMnO3 bilayers.13 By tuning the Co–Mn charge transfer and the interfacial orbital polarization via ionic-liquid gating, a reorientation of the magnetic easy axis in the in-plane direction is achieved. Additional reports on compressively strained LaSrCoO3/LSMO/LSrCoO3 trilayers show an in-plane easy axis in contrast to a PMA observed for single layer compressively strained LSMO on LaAlO3 due to off-center Mn displacements, which favor a preferential occupation of the in-plane Mn 3d x2y2 orbitals.15 

For uncapped stoichiometric LSMO films, below a critical thickness of 4–10 nm, ferromagnetism is suppressed, limiting potential applications in thin film devices and a direct decoupling of the various contributions to magnetic anisotropy.5,16 The suppression of ferromagnetism in thin LSMO films has been attributed to structural distortions and chemical and electronic reconstructions arising from interfacial polar discontinuities,17–20 oxygen vacancies,21 cation disorder,22 orbital reconstructions,23 and distortions of the oxygen octahedra due to the interfacial structural coupling.24–27 Recent reports on coherently strained LSMO layers in LSMO/SrRuO3 superlattices28 and LSMO layers capped with La0.7Sr0.3CrO3 (LSCO) have evidenced ferromagnetism in LSMO layers as thin as 2 unit cells (0.8 nm).6,29–31

In this letter, we investigate the structural and magnetic properties of strain-free LSCO/LSMO heterostructures grown on (001)-oriented single crystal MgO substrates.32–34 Bulk LSMO and LSCO have pseudocubic lattice constants of 3.88 and 3.86 Å, respectively. The lattice mismatch of LSMO and LSCO with cubic MgO (c = 4.212 Å) is 8.5%. Due to the large tensile lattice mismatch between the MgO substrate and the LSMO and LSCO layers, strain relaxation is found to occur within the first 2–3 unit cells of the film as evidenced by in situ reflection high energy electron diffraction (RHEED) measurements and ex situ x-ray diffraction. A paramagnetic–ferromagnetic transition is observed for heterostructures with 2 unit cell (uc) thick LSMO layers indicative of the removal of magnetically dead LSMO layers. The relaxed states of the ultra-thin LSMO films allow us to elucidate the role confinement places in the magnetic anisotropy in this system. The magnetic easy axis is found to lie in the plane of the film due to the dominant contributions of LSCO/LSMO interfacial interactions and confinement to the total magnetic anisotropy.

Trilayer 3 uc LSCO/3 uc LSMO/3uc LSCO films [referred to as (3/3/3)] and [2ucLSCO/NucLSMO]x superlattices (where x is the number of repeats) were fabricated by molecular beam epitaxy (MBE) on (001)-oriented MgO substrates. The LSMO layer thickness, N, was varied from 2 to 6 uc and the superlattices were capped with 2 uc LSCO. Prior to growth, the MgO substrates were etched in buffered hydrofluoric acid and annealed in a tube furnace at 1200°C to achieve atomically flat surfaces. The films were grown at a substrate temperature determined by an optical pyrometer of 850±25°C using an oxygen plasma source with an oxygen partial pressure of 3×106 Torr. After growth, the films were cooled to room temperature at a rate of 5°C/min at the growth oxygen pressure to minimize the formation of oxygen vacancies. In situ reflection high energy electron diffraction (RHEED) was used to monitor the film thickness and crystallinity during growth.

To determine the strain states of the samples, reciprocal space maps of the 3/3/3 trilayer sample were measured at room temperature the 33ID beamline at the Advanced Photon Source with a Pilatus 100 K pixel detector.35 Specular diffraction scans of the film and substrate Bragg peaks for the [2LSCO/4LSMO]6 and [2LSCO/6LSMO]4 superlattices were measured using a Rigaku Smartlab diffractometer equipped with a Ge(220) double bounce monochromator.

Temperature and magnetic field-dependent measurements of the magnetization of the samples were performed by superconducting quantum interference device (SQUID) magnetometry using a Quantum Design MPMS 3 system. The temperature-dependent magnetization curves were measured on warming up the sample with an applied 0.1 T magnetic field after field cooling in an 0.2 T magnetic field. The magnetization measurements were performed with the applied magnetic field oriented either parallel to the film [100] axis (in-plane) or to the film [001] axis (out-of-plane).

The evolution of the RHEED pattern during the growth of the first LSCO and LSMO layers is shown in Fig. 1. A transition in the RHEED spectra from the diffraction pattern for the MgO substrate surface in Fig. 1(a) to 3D spots [Fig. 1(b)] after deposition of 2 uc of LSCO is indicative of initial island growth and surface roughening. Additionally, an increase in the spacing of the RHEED streaks indicates a decrease in the in-plane lattice constant from 4.21 Å for the MgO surface to 3.87±0.02 Å for the LSCO adlayer. On deposition of 6 uc of LSMO, the RHEED pattern transitions from the 3D pattern to 2D streaks shown in Fig. 1(c) indicative of a smoothening of the film surface. The initial roughening is attributed to a 3D Volmer–Weber island growth to relax the large film-substrate lattice mismatch.36 

FIG. 1.

Reflection high energy diffraction (RHEED) pattern recorded along high symmetry zone axes during the growth of a 2 uc LSCO/6 uc LSMO bilayer for (a) the initial MgO substrate, (b) the first 2 uc of LSCO indicating 3D growth, and (c) 6 uc of LSMO indicating a transition to 2D growth. The shift of the (2 0) diffraction pattern of the film is indicative of a contraction of the in-plane lattice constant from 4.21 Å for the MgO substrate to 3.87 Å for the LSCO and LSMO layers. The vertical dashed lines indicate the position of the bulk MgO reflections.

FIG. 1.

Reflection high energy diffraction (RHEED) pattern recorded along high symmetry zone axes during the growth of a 2 uc LSCO/6 uc LSMO bilayer for (a) the initial MgO substrate, (b) the first 2 uc of LSCO indicating 3D growth, and (c) 6 uc of LSMO indicating a transition to 2D growth. The shift of the (2 0) diffraction pattern of the film is indicative of a contraction of the in-plane lattice constant from 4.21 Å for the MgO substrate to 3.87 Å for the LSCO and LSMO layers. The vertical dashed lines indicate the position of the bulk MgO reflections.

Close modal

To confirm the relaxation of strain in the LSCO/LSMO layers, reciprocal space maps were measured at the 33ID beamline at the Advanced Photon Source with an incident photon energy of 15.5 keV. Figure 2 shows the reciprocal space map around the LSCO/LSMO film (111) and the MgO (111) Bragg peak for a trilayer 3 uc LSCO/3uc LSMO/3 uc LSCO sample on MgO. The location of the film Bragg peak corresponds to a relaxed pseudo-cubic lattice constant of a=b=c=3.87±0.01 Å.

FIG. 2.

(a) Schematic of 3 uc LaSrCrO3/3 uc LaSrMnO3/3 uc LaSrCrO3 heterostructure grown on (001)-oriented MgO by molecular beam epitaxy. (b) The reciprocal space map around the film and substrate (111) Bragg peaks indicate film strain relaxation. The yellow and black vertical dashed lines indicate the in-plane reciprocal vector corresponding to bulk MgO and LSMO, respectively.

FIG. 2.

(a) Schematic of 3 uc LaSrCrO3/3 uc LaSrMnO3/3 uc LaSrCrO3 heterostructure grown on (001)-oriented MgO by molecular beam epitaxy. (b) The reciprocal space map around the film and substrate (111) Bragg peaks indicate film strain relaxation. The yellow and black vertical dashed lines indicate the in-plane reciprocal vector corresponding to bulk MgO and LSMO, respectively.

Close modal

The out-of-plane lattice constants of the superlattices are determined from specular diffraction scans around the film and substrate (002) Bragg peaks. The (002) Bragg peaks for [2ucLSCO/4ucLSMO]6 and [2ucLSCO/6ucLSMO]4 superlattices on MgO substrates are shown in Fig. 3. The fits to the measured data are obtained using the GenX x-ray analysis program.37 The average out-of-plane lattice constant determined from the fits for the LSMO and LSCO layers are 3.868±0.005 and 3.862±0.005 Å, respectively. The calculated lattice parameters obtained from the fit are very close to the bulk LSMO lattice constant, which further supports the lattice relaxation in LSCO/LSMO heterostructures on MgO.38 

FIG. 3.

Measured (circles) and simulated (solid lines) specular diffraction of the (002) Bragg peaks for the [2LSCO/2LSMO]10 superlattice and the [2LSCO/6LSMO]4 superlattice on the MgO substrate. The diffraction spectra are offset for clarity. L is the reciprocal lattice unit of the MgO substrate, where 1r.l.u.=(1/4.212)Å1.

FIG. 3.

Measured (circles) and simulated (solid lines) specular diffraction of the (002) Bragg peaks for the [2LSCO/2LSMO]10 superlattice and the [2LSCO/6LSMO]4 superlattice on the MgO substrate. The diffraction spectra are offset for clarity. L is the reciprocal lattice unit of the MgO substrate, where 1r.l.u.=(1/4.212)Å1.

Close modal

To determine the effect of the lattice relaxation on the magnetic properties of the heterostructures, magnetization as a function of the applied magnetic field and the temperature was measured by SQUID magnetometry. The temperature-dependent curves were measured with a 1000 Oe magnetic field applied either in-plane or out-of-plane on warming after field cooling in a 0.2 T applied magnetic field. The magnetization as a function of magnetic field and temperature for the [2LSCO/2LSMO]10 and [2LSCO/4LSMO]6 superlattices on MgO is shown in Figs. 4(a) and 4(b), respectively, with the magnetic field oriented either in-plane or out-of plane as shown in the inset in Fig. 4(a). For both samples, the magnetization is significantly reduced in the out-of-plane direction compared to the in-plane magnetization indicative of an in-plane magnetic easy axis. For the [2LSCO/2LSMO]10 sample, the out-of-plane magnetization does not reach the saturation value measured for the in-plane orientation.

FIG. 4.

(a) Temperature-dependent in-plane and out-of-plane magnetization measurements with 1000 Oe applied field on warming after cooling in a 0.2 T field. (b) Field-dependent magnetization curves at 10 K for [2LSCO/2LSMO]10 and [2LSCO/4LSMO]6 superlattices grown on (001)-oriented MgO by molecular beam epitaxy.

FIG. 4.

(a) Temperature-dependent in-plane and out-of-plane magnetization measurements with 1000 Oe applied field on warming after cooling in a 0.2 T field. (b) Field-dependent magnetization curves at 10 K for [2LSCO/2LSMO]10 and [2LSCO/4LSMO]6 superlattices grown on (001)-oriented MgO by molecular beam epitaxy.

Close modal

Furthermore, the observation of ferromagnetism in the heterostructure with a LSMO thickness of 2 uc is attributed to the removal of magnetic dead layers when using LSCO as spacer due to the interfacial structural coupling and the anti-ferromagnetic coupling between Cr and Mn across the LSCO/LSMO interface.6 The measured magnetic moments for the relaxed heterostructures on LSMO are expected to be close or identical to analogous heterostructures grown on LSAT where the lattice mismatch is small (0.2%).29 The magnetic moments for the [2LSCO/2LSMO]10 and [2LSCO/4LSMO]6 heterostructures on LSAT at 10 K are 1.5 μB/Mn and 2.8 μB/Mn, respectively. The corresponding moments for the [2LSCO/2LSMO]10 and [2LSCO/4LSMO]6 on MgO are 1.4 and 2.4 μB/Mn, respectively. The increase in the magnetization per Mn with increasing LSMO thickness is attributed to the contribution to the total magnetization of the Cr spins aligned anti-parallel to the applied magnetic field.6,29

To quantify the magnetic anisotropy (MA) between the magnetic easy and hard axis, we calculate the magnitude of the effective anisotropy constant, Keff, by the area enclosed between the in-plane, Mip, and out-of-plane, Moop, magnetization curves shown in Fig. 4(b) for the [2LSCO/4LSMO]6 heterostructure.39,40Keff for the [2LSCO/4LSMO]6 configuration is approximately 2×105ergs/cm3, which is comparable to that of strained LSMO.8 The measured anisotropy for the relaxed heterostructure is unexpected since the magnetocrystalline anisotropy for bulk LSMO is 1 order of magnitude lower.41 

For thin films, the effective anisotropy energy Keff can be expressed in terms of a bulk magnetocrystalline contribution KV and a surface/interface contribution KS as39 

Keff=μ00HSMoopdHμ00HSMipdH=KV+2πMS+KStFM,
(1)

where KV is the intrinsic bulk anisotropy and KS/tFM is the interface anisotropy due to the interfacial exchange bias and/or surface and interface induced structural and orbital reconstructions and tFM is the thickness of the ferromagnetic LSMO layer.15,23,42 The 2πMS term is the shape magnetic anisotropy. A positive Keff corresponds to a perpendicular magnetic anisotropy and Keff<0 corresponds to an in-plane easy axis.39 Thus, the negative Keff observed for the LSCO/LSMO heterostructures suggests that the interfacial contribution KS/tFM, which depends on the film thickness, is negative and its contribution to the total magnetic anisotropy will decrease with increasing tFM. For the [2 LSCO/2 LSMO] heterostructure, the out-of-plane magnetization does not the reach saturation magnetization observed for the measurements with the magnetic field oriented in-plane suggesting that Keff is dominated by KS. We note that the net magnetization measured by SQUID for the LSCO/LSMO heterostructures depends on the magnetization in both the Cr and Mn layers; thus, the effect of magnetic anisotropy on both the LSMO and LSCO layers determines the total magnetic response.6,29 Further studies are required to determine how the orientation of the external field affects the interfacial magnetic exchange at the LSCO/LSMO interface.

To confirm that the easy in-plane axis observed for the relaxed LSCO/LSMO heterostructures on MgO is independent of strain, element-selective x-ray magnetic circular dichroism (XMCD) measurements are carried out as a function of an applied magnetic field for a [2LSMO/2LSCO]10 superlattice under coherent 2% compressive strain on a LaAlO3 substrate.29 The XMCD measurements were performed at the 4.0.2 beamline at the Advanced Light Source using total electron yield with an incident beam angle at 30° to the sample surface. The x-ray absorption spectra (XAS) at the Mn L-edge measured for right [I(σ+)] and left [I(σ)] circular polarized incident photons with a 0.5 T magnetic field applied in the plane of the sample are compared in Fig. 5(a). The XMCD signal is determined from the difference between I(σ+) and I(σ). Furthermore, XMCD magnetic hysteresis loops are obtained by fixing the photon energy at the Mn L3 edge and measuring the XMCD signal as the applied magnetic field is swept between 0.5T and 0.5 T. We compare the XMCD hysteresis loops measured at the Mn L3 absorption edge at 150 K for a magnetic field applied in-plane and out-of-plane as shown in Fig. 5(b). The XMCD results indicate an in-plane easy axis for the compressively strained LSCO/LSMO heterostructure; hence, the observed MA is attributed to interactions at the LSCO/LSMO interface, which dominate the magnetocrystalline and shape anisotropy for ultra-thin LSMO layers.

FIG. 5.

(a) X-ray absorption spectra measured at the Mn L-edge for right [I(σ+)] and left [I(σ)] circular photon polarizations and the XMCD signal, I(σ+)-I(σ) for a [2LSCO/2LSMO]10 heterostructure on LaAlO3. The measurements are performed at 150 K with a 0.5 T magnetic field applied in the plane of the sample. (b) Magnetic field dependence of the XMCD signal measured at the Mn L3-edge and 150 K for a [2LSCO/2LSMO]10/LaAlO3 sample with the field applied either parallel or perpendicular to the film surface.

FIG. 5.

(a) X-ray absorption spectra measured at the Mn L-edge for right [I(σ+)] and left [I(σ)] circular photon polarizations and the XMCD signal, I(σ+)-I(σ) for a [2LSCO/2LSMO]10 heterostructure on LaAlO3. The measurements are performed at 150 K with a 0.5 T magnetic field applied in the plane of the sample. (b) Magnetic field dependence of the XMCD signal measured at the Mn L3-edge and 150 K for a [2LSCO/2LSMO]10/LaAlO3 sample with the field applied either parallel or perpendicular to the film surface.

Close modal

In conclusion, we have investigated the effect of LSCO spacer layers on the magnetic properties of thin LSMO layers grown on (001)-MgO substrates. The LSCO/LSMO heterostructures are found to be relaxed as a result of the large lattice mismatch with the MgO substrate with strain relaxation occurring within 2–3 unit cells. The magnetic easy axis of the LSCO/LSMO heterostructures is found to lie in the plane of the film due to a dominant contribution of the LSMO interface magnetic anisotropy to the total magnetic anisotropy. This work demonstrates the critical role interfacial interactions play in modulating the magnetic states of transition metal oxide heterostructures.

S.K., Y.O., and D.P.K. acknowledge financial support from the U.S. National Science Foundation under Grant No. NSF DMR-1751455. This work was performed in part at the Analytical Instrumentation Facility at North Carolina State University, which is supported by the State of North Carolina and the National Science Foundation (Award No. ECCS-1542015). This work made use of instrumentation at AIF acquired with support from the National Science Foundation (NSF) (DMR-1726294). The AIF is a member of the North Carolina Research Triangle Nanotechnology Network (RTNN), a site in the National Nanotechnology Coordinated Infrastructure (NNCI). The authors acknowledge use of the SQUID and PPMS facility in the Department of Materials Science and Engineering at North Carolina State University. Use of the Advanced Photon Source was supported by the U.S. Department of Energy (DOE), Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. This research used resources of the Advanced Light Source, which is a DOE Office of Science User Facility under Contract No. DE-AC02-05CH11231.

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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