Role of the magnetic anisotropy in organic spin valves 2017 Journal of Science Advanced Materials and Devices

Accepted Manuscript
Role of the magnetic anisotropy in organic spin valves
V. Kalappattil, R. Geng, S.H. Liang, D. Mukherjee, J. Devkota, A. Roy, M.H. Luong,
N.D. Lai, L.A. Hornak, T.D. Nguyen, W.B. Zhao, X.G. Li, N.H. Duc, R. Das, S.
Chandra, H. Srikanth, M.H. Phan
PII:

S2468-2179(17)30131-4

DOI:

10.1016/j.jsamd.2017.07.010

Reference:

JSAMD 114

To appear in:

Journal of Science: Advanced Materials and Devices

Received Date: 28 July 2017
Revised Date:

2468-2179 2468-2179

Accepted Date: 31 July 2017

Please cite this article as: V. Kalappattil, R. Geng, S.H. Liang, D. Mukherjee, J. Devkota, A. Roy,
M.H. Luong, N.D. Lai, L.A. Hornak, T.D. Nguyen, W.B. Zhao, X.G. Li, N.H. Duc, R. Das, S. Chandra,
H. Srikanth, M.H. Phan, Role of the magnetic anisotropy in organic spin valves, Journal of Science:

Advanced Materials and Devices (2017), doi: 10.1016/j.jsamd.2017.07.010.
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ACCEPTED MANUSCRIPT

Role of the magnetic anisotropy in organic spin valves
V. Kalappattil1,#, R. Geng2,#, S.H. Liang2, D. Mukherjee1, J. Devkota2, A. Roy3, M.H. Luong2,4,
N.D. Lai4, L.A. Hornak5, T.D. Nguyen2,*, W.B. Zhao6, X.G. Li6, N.H. Duc7, R. Das1, S.

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Department of Physics, University of South Florida, Tampa, Florida 33620, USA

Department of Physics and Astronomy, University of Georgia, Athens, GA 30602, USA
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Department of Chemistry, University of Georgia, Athens, GA 30602, USA

Laboratoire de Photonique Quantique et Moléculaire, Ecole Normale Supérieure de Cachan,

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Chandra1, H. Srikanth1, and M.H. Phan1,*

UMR 8537, CentraleSupélec, CNRS, Université Paris-Saclay, 94235 Cachan, France
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College of Engineering, University of Georgia, Athens, GA 30602, USA

Hefei National Laboratory for Physical Sciences at Microscale, Department of Physics,

University of Science and Technology of China, Hefei 230026, and Collaborative Innovation

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Center of Advanced Microstructures, Nanjing University, Nanjing 210093, China
VNU Key Laboratory for Micro-nano Technology and Faculty of Physics Engineering and

Nanotechnology, VNU University of Engineering and Technology, Vietnam National University,

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Abstract

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Hanoi, 144 Xuan Thuy Road, Cau Giay, Hanoi, Viet Nam

Magnetic anisotropy plays an important role in determining the magnetic functionality of thin
film based electronic devices. We present here, the first systematic study of the correlation
between magnetoresistance (MR) response in organic spin valves (OSVs) and magnetic
anisotropy of the bottom ferromagnetic electrode over a wide temperature range (10K – 350K).
The magnetic anisotropy of a La0.67Sr0.33MnO3 (LSMO) film epitaxially grown on a SrTiO3
(STO) substrate was manipulated by reducing film thickness from 200 nm to 20 nm. Substrate-


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induced compressive strain was shown to drastically increase the bulk in-plane magnetic
anisotropy when the LSMO became thinner. In contrast, the MR response of LSMO/OSC/Co
OSVs for many organic semiconductors (OSCs) does not depend on either the in-plane magnetic

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anisotropy of the LSMO electrodes or their bulk magnetization. All the studied OSV devices
show a similar temperature dependence of MR, indicating a similar temperature-dependent
spinterface effect irrespective of LSMO thickness, resulting from the orbital hybridization of

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carriers at the OSC/LSMO interface.

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Equal contributions to the work

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Keywords: LSMO; Magnetic anisotropy; Magnetoresistance; Organic Spintronics

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* Corresponding authors: (T.D. Nguyen); (M.H. Phan)


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Role of the magnetic anisotropy in organic spin valves
Abstract
Magnetic anisotropy plays an important role in determining the magnetic functionality of thin

film based electronic devices. We present here, the first systematic study of the correlation

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between magnetoresistance (MR) response in organic spin valves (OSVs) and magnetic
anisotropy of the bottom ferromagnetic electrode over a wide temperature range (10K –
350K). The magnetic anisotropy of a La0.67Sr0.33MnO3 (LSMO) film epitaxially grown on a

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SrTiO3 (STO) substrate was manipulated by reducing film thickness from 200 nm to 20 nm.
Substrate-induced compressive strain was shown to drastically increase the bulk in-plane

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magnetic anisotropy when the LSMO became thinner. In contrast, the MR response of
LSMO/OSC/Co OSVs for many organic semiconductors (OSCs) does not depend on either
the in-plane magnetic anisotropy of the LSMO electrodes or their bulk magnetization. All the
studied OSV devices show a similar temperature dependence of MR, indicating a similar

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temperature-dependent spinterface effect irrespective of LSMO thickness, resulting from the
orbital hybridization of carriers at the OSC/LSMO interface.

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1. Introduction

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Keywords: LSMO; Magnetic anisotropy; Magnetoresistance; Organic Spintronics

La1-xSrxMnO3 (LSMO, x~0.33) is a very promising candidate material for spintronic

devices applications due to its chemical stability and intriguing physical properties [1-6]. In
particular, LSMO is a half-metallic ferromagnet that acts as an excellent spin injector/detector
due to near 100% spin polarization at low temperatures [3,4]. This material also provides
manufacturing flexibility and cost-effectiveness, which are of practical importance [3].
Organic spintronics based on LSMO have attracted growing interest since Xiong et al.
reported in 2004 a giant magneto-resistance (MR) of ~40% at 11 K in a LSMO/Alq3/Co spin
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valve structure [7]. The similar effects have latter been reported in many organic
semiconductors (OSCs) [8]. An extremely large MR value of up to 300% was achieved at
10K within this type of device when an interfacial diffusion between the Co electrode and the
Alq3 organic spacer was considerably suppressed by adding a thin interlayer of arrayed Co

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nanorods [9]. This unusually large MR effect caused by a large effective Co spin polarization
has been attributed to the spinterface effect, which is generally accepted to be strong in OSVs

[8,10,11]. The accomplished MR is also associated with the long life spin of electrons in

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organic materials. As compared to inorganic semiconductor-based devices, the organic ones
are cheaper and more mechanically flexible [3]. However, it has been reported by several

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research groups that the MR of the LSMO/Alq3/Co device decreases drastically with an
increase in temperature (T < ~200 K) and reaches a relatively small value at room temperature
(0.15−9%), rendering it undesirable for practical use [3,9,12-14]. The MR temperature
dependence of the LSMO/Alq3/Co device has remained an issue of long-lasting debate

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[3,7,9,12-15]. It was attributed to the reduction of the spin relaxation rate in the Alq3 layer [7]
or/and the weakening of spin polarization of ferromagnetic electrodes at high temperature
range (T > ~200 K) [14,16]. Since the Curie temperature of Co (TC~1388 K) is much higher

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than that of LSMO (TC~360 K), a considerably reduced spin polarization of LSMO near 300

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K was naturally thought of as a possible cause for the observed small MR values [14]. While
the spin-1/2 photoluminescence detected magnetic resonance (PLDMR) study on Alq3
revealed a weak temperature dependence of spin lattice relaxation time [14], Drew et al.
employed a low energy muon spin rotation method to study the temperature dependence of
spin diffusion length, demonstrating that the spin relaxation in Alq3 dominated the MR
temperature dependence [17]. Recently, Chen et al. related the MR temperature dependence
to the spin relaxation in Alq3 for T < ~100 K, but to the surface spin polarization (or
spinterface) of LSMO for T > ~100 K [12]. This seems to be supported by Majumdar et al.
who also observed a noticeable difference in the MR ratio for T > ~100 K between the two
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LSMO/Alq3/Co devices made of LSMO films grown on SrTiO3 and MgO substrates [18].
Due to a larger lattice mismatch, a larger strain (9%) was reported in the LSMO film grown
on MgO as compared to the LSMO film grown on SrTiO3 (1%) [18]. This suggests that the
observed MR difference for T > 100 K should not be simply related to the loss of spin

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polarization carriers at the LSMO/Alq3 interface [18], but also due to the substrate-induced
strain effect [19]. On the other hand, inorganic spintronics based on LSMO has been under
investigation for a long time [1,2,4,20] and the magnetic anisotropy of LSMO has been

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reported to play a crucial role in controlling the performance of these devices [2,19,20].
Unfortunately, to our best knowledge, no work has dealt with the role of magnetic anisotropy

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of LSMO electrodes in organic spin valves.

To address this important issue, we have performed the first comparative study of the
bulk and surface magnetic properties of LSMO films with distinct thicknesses (20 nm and 200
nm), as well as the MR responses of LSMO/Alq3/Co devices using these LSMO films as

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electrodes. By reducing film thickness from 200 nm to 20 nm, the magnetic anisotropy of
LSMO was drastically increased, due to the enhanced SrTiO3 (STO) substrate-induced strain
effect, allowing for probing effects of the magnetic anisotropy of the LSMO film on the MR

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response of OSV devices. Our results indicate that instead of the in-plane magnetic anisotropy
of the LSMO electrode, the effective surface spin polarization at the OSCs/LSMO interface or

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a spinterface effect plays an important role in the spin injection and transport in
LSMO/Alq3/Co spin valve devices. Our observation of the negligible influence of STO
substrate strain on the MR response indicates that the variation in substrate strain is not of
significant concern while producing large numbers of OSV devices. For achieving a larger
MR at room temperature, however, other types of ferromagnetic half-metals needs to be

employed.
2. Experiment
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2.1. Sample Preparation
LSMO films of various thicknesses (200 nm, 50 nm, 20 nm) were grown epitaxially
on single-crystal SrTiO3 (STO) (100) substrates at 750 °C using magnetron sputtering
technique, with Ar and O2 flux in the ratio of 1:1 in a pressure 4 Pa. The films were

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subsequently annealed at 800 °C for 2 hours in flowing O2 atmosphere before slowly cooled
to room temperature. For the OSV device fabrication, the LSMO films were subsequently
patterned using standard photolithography and chemical etching techniques. The Alq3 spacer

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was thermally evaporated using an organic evaporation furnace with the evaporation rate of
0.5 Å /s at the base pressure of 2 × 10−7 torr; 15 nm cobalt (capped by 50 nm Al) top electrode

typically about 0.2 × 0.4 mm2.
2.2. Properties Characterization

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was deposited onto the Alq3 spacer using a shadow mask. The obtained active device area was

The crystallinity and crystallographic orientations in the heterostructures were

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characterized by X-ray diffraction (XRD) (Bruker AXS D8 diffractometer equipped with
high-resolution Lynx Eye position-sensitive detector using Cu Kα radiation). The in-plane
epitaxy was determined from XRD azimuthal (ϕ) scans (Philips X’pert Diffractometer). The

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surface morphologies were observed using an atomic force microscope (AFM) (Digital
Instruments III). Magnetic measurements were performed at different temperatures (10-350K)

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using a commercial Physical Property Measurement System (PPMS) (Quantum Design, Inc.)
in magnetic fields up to 5T. Temperature dependent magnetic anisotropy measurements were
measured by the radio-frequency transverse susceptibility (TS) using a RF tunnel diode
oscillator (TDO) integrated into the PPMS. The surface magnetic properties of LSMO films
were studied by regular magneto-optic Kerr effect (MOKE) and balanced magneto-optic Kerr
effect (B-MOKE). Magnetoresistance (MR) measurements on the OSV devices were
conducted using the four probe technique in the presence of an in-plane magnetic field up to 3
kOe.
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3. Results and Discussion
First we examined the structure of the grown films and performed a substrate-induced
strain analysis on them using the XRD technique. Figures 1a and 1b show the XRD θ-2θ
patterns for the 200 nm and 20 nm LSMO films, respectively. In both cases, only (00l) (l = 1,

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2, and 3) diffraction peaks of the pseudo-cubic perovskite LSMO phase (JCPDS 01-0894461) are observed along with the (00l) peaks of the STO substrates. No secondary phase
formations are present within the resolution limits of the instrument. The small lattice

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mismatch between bulk pseudo-cubic LSMO (lattice parameter, a = 0.389 nm) and cubic
STO (a = 0.3905 nm) allows for the epitaxial growth of LSMO on STO as evident from the

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close proximity of the XRD peak positions (inset to Fig. 1a). Atomic force microscopy
(AFM) images of the LSMO films are shown in insets of Fig. 1a and 1b, with a similar root
mean squared (RMS) surface roughness of ~0.6 nm. While the double exchange mechanism
alone cannot explain the magnetism of LSMO, a strong electron lattice coupling due to Jahn-

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Teller distortions via MnO6 octahedral deformation has been demonstrated to play an

important role [21]. Although bulk manganites show small magnetic anisotropy, in thin films

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it differs substantially from the bulk because of epitaxial strain in the films [22].
The average out-of-plane (a⊥) and in-plane (a‫ )װ‬lattice parameters of the LSMO films

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with 200 nm and 20 nm thicknesses grown on STO (100) substrates were calculated from the
symmetric θ-2θ scans (performed about the LSMO (100), (200), (300) pseudo-cubic planes)
and the asymmetric 2θ-ω (or detector) scans (performed about the LSMO (110) and (211)
planes), following the same method as detailed in previous works [23,24]. A representative
detector scan is shown in the inset of Fig. 1b for the 20 nm LSMO film performed about the
LSMO (211) plane. The lattice parameters obtained from the XRD analysis were calculations
were a‫= װ‬0.381 (±0.003) nm and a⊥= 0.393 nm (±0.001) for the 20 nm film and those for the
200 nm films were a‫ = װ‬0.392 (±0.002) nm and a⊥= 0.391 nm, respectively. Difference
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between the in-plane and out-of-plane lattice parameters of each film indicates the lattice
distortion during their growing. In-plane strain was calculated by using the formula (a‫– װ‬
ao)/ao, where ao is the bulk lattice parameter of LSMO as measured from the XRD powder
diffraction [23]. A large compressive strain of – 0.020 is found in the 20 nm LSMO film,

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while a relatively small tensile strain of 0.005 is seen in the 200 nm LSMO film. Also by
comparing the a‫ װ‬and a⊥ values, we see that the 20 nm LSMO film undergoes a larger
tetragonal distortion due to lattice induced strain. This is in accordance with the thickness-

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dependent epitaxial strain study conducted by the other research group [19]. In the present
study, we aimed to understand how this strain variation would affect the surface and bulk

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magnetic properties of the LSMO films and hence the MR responses of the SV devices using
these LSMO films as ferromagnetic electrodes.

To understand how surface magnetic properties of LSMO films differ from their bulk
ones, magneto-optical Kerr effect (MOKE) measurement and vibrating sample magnetometer

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(VSM) measurements were performed over the temperature range of 10-350K. We recall that
the MOKE technique is an excellent tool for studying the surface magnetization as it is highly

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sensitive to the magnetization within the skin depth region of 10-15nm in most materials.
Figures 2a and 2b display the MOKE data at two selected temperatures (127 K and 215 K) for

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the 200 nm and 20 nm LSMO films, respectively. The coercive field (HC) of each film was
determined from the MOKE loop and the temperature dependence of HC for both films are
plotted in Fig. 2c. As compared to the 20 nm LSMO film, the values of HC (close to values of
the surface magnetic anisotropy field) are larger for the 200 nm LSMO film, especially in the
high temperature range. Figures 2d and 2e show the VSM loops at two selected temperatures
(50 K and 300 K) for the 200 nm and 20 nm LSMO films, respectively. The temperature
dependences of HC for both films are also plotted in Fig. 2f. We note that the HC value of the
200 nm LSMO film obtained from MOKE (Fig. 2c) is similar to that obtained from VSM at
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all temperatures (Fig. 2f), indicating that the surface magnetic property of the 200 nm LSMO
film was not significantly affected by the substrate-induced strain effect. By contrast, the HC
value of the 20 nm LSMO film obtained from MOKE is larger than that obtained from VSM,
and the difference tends to increase with lowering temperature (inset in Fig. 2f), indicating

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that the surface magnetic property of the 20 nm LSMO film was affected by the substrate
strain. Furthermore, VSM revealed a large difference in shape of the M-H loop (at 50 K)
when the thickness of the LSMO film was reduced from 200 nm to 20 nm (Fig. 2d and 2e).

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increased effective anisotropy of the 20 nm LSMO film.

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This indicates that the compressive strain induced by the STO substrate gave rise to the

To quantify the effective anisotropy field value (HK) and its temperature evolution, we
measured the radio-frequency (RF) transverse susceptibility (TS) of both LSMO films.25 This
TS method has been validated by us as an excellent tool for studying the anisotropic magnetic
properties of a variety of systems from thin films [26] to single crystals [27] and nanoparticles

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[28]. Here the sample is kept inside an LC circuit which is self-resonant around 12 MHz with
a sensitivity of 1-10 Hz in 10 MHz. When a small RF field is applied perpendicular to the
sweeping DC field, change of the resonant frequency (∆fres) is directly proportional to the

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change of the magnetic susceptibility (∆χT) in the transverse direction: ∆χT/χT ∝ ∆fres/fres. As

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theoretically predicted by Aharoni et al. [29] for a Stoner-Wohlfarth particle with its magnetic
hard axis aligned parallel to the DC field, TS spectra should yield peaks at the anisotropy
fields (±HK) and switching fields (±HS) as the DC field is swept from positive to negative
saturation. TS curves taken at different temperatures for the 200 nm and 20 nm LSMO films

are displayed in Fig. 3a and 3b, respectively. For comparison of the ∆χT and HK, the TS
spectra taken at the same temperature of 20 K are plotted in Fig. 3c for both films. We note
that for the present LSMO films the switching peak is merged with the anisotropy peak, thus
causing a difference in the positive and negative HK values as well as a slight asymmetry in
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the peak height (Fig. 3c). For analysis purposes, we have consistently taken positive values of
HK for both films. As expected, HK increased with lowering temperature below the Curie
temperature of the film (Fig. 3d). It is worth mentioning that there is a big difference in HK
between the 20 nm and 200 nm LSMO films. At 20 K, the HK value of 597 Oe obtained for

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the 20 nm film is about 5 times larger than that of the 200 nm LSMO film (HK = 102 Oe).
Another important feature of note is that for the 20 nm LSMO film the values of HK obtained
from TS (Fig. 3d) are much larger than those estimated from MOKE (Fig. 2c), unlike in the

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case of the 200 nm LSMO film. For the 20 nm LSMO film, at 20 K, HK = 597 Oe and 60 Oe
were obtained from TS and MOKE measurements, respectively. Such a big difference in HK

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can be attributed to the fact that the epitaxial strain induced by the STO substrate was large in

the case of thin films (20 nm) as oppose to thick films (200 nm).

Upon understanding how the magnetic anisotropy evolved with temperature and its
difference in the 20 nm and 200 nm LSMO films, we then examined its effect on the MR

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response of the organic spin valves using these films as electrodes. Firstly, two devices of
LSMO(200nm)/Alq3(100nm)/Co(15nm) and LSMO(20nm)/Alq3(100nm)/Co(15nm) were

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fabricated. A schematic of the LSMO/Alq3/Co device is illustrated in Fig. 4a. Typical MR
curves taken at 10 K for the two devices are displayed in Fig. 4b and 4c. In order to

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reasonably compare the MR response between different devices, we have subtracted the
background MR contribution (“V shape”) such as anisotropic magnetoresistance (AMR)
caused by the magnetic electrodes [30,31] which has been labeled as blue dash lines (or MR
value taken at small applied magnetic field). There is no noticeable difference in shape of the
MR curves, except for a slightly larger MR for the 200 nm LSMO film, suggesting similarity
of their interfacial spin polarization. Figure 4d shows the temperature dependences of the
normalized balanced MOKE (B-MOKE) and the normalized bulk magnetization for the 200
nm and 50 nm LSMO films. We note that the B-MOKE signal is proportional to the
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magnitude of the surface magnetization within a certain skin depth region; therefore it’s a
very useful tool to investigate the surface magnetization of the LSMO films as a function of
temperature. We find that the surface magnetization decays faster than the bulk magnetization
with increasing temperature (Fig. 4d), which is in agreement with that reported by Park et al.

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using spin-resolved photoemission spectroscopy [32]. Furthermore, we note that for T > ~200
K the bulk spin polarization of the 200 nm LSMO film was larger than that of the 20 nm
LSMO film (see Fig. 4d). The difference in bulk spin polarization (proportional to ∆M)

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tended to increase with increasing temperature for these films, but only negligible difference
in MR was observed for these films (Fig. 4e), where the temperature dependence of

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normalized MR using various thicknesses of LSMO film was measured. It is interesting to
note that the temperature evolution of MR follows the same trend for all LSMO films (Fig.
4e), but does not resemble the temperature dependence of HC (Fig. 2f), HK (Fig. 3d), the bulk
magnetization (Fig. 4d) and even the surface magnetization (Fig. 2c and Fig. 4d) (measured

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by regular MOKE and balanced-MOKE, respectively) of these films. This important
observation points to the fact that the nature of orbital hybridization at the Alq3/LSMO
contact decides interfacial spin polarization (spinterface) and therefore the MR response. In

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order to rule out the possibility that the spin relaxation dominates the MR temperature

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dependence in our Alq3-based devices [7,12,17] we further investigated the time dependent
MR of various OSCs as the interlayers in the same device structure LSMO
(50nm)/OSC/Co(15nm)/Al. Figure 4f shows the time dependence of normalized MR for
different OSCs, and it is worth noting that they all follow a similar trend in terms of
temperature regardless of their different spin diffusion lengths [8,33-35]. If their spin
diffusion length is temperature dependent, one should see different trends of temperaturedependent MR, which are not observed in our experiments. Combined MOKE, VSM, TS and
MR data provide solid evidence that the in-plane magnetic anisotropy does not play a decisive
role in determining the MR of LSMO/Alq3/Co devices, which is in stark contrast to what is
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observed in LSMO-based inorganic devices [1,2,4]. Our findings suggest that the spinterface
of organic/LSMO is quite universal. For achieving larger MR effects at room temperature,
either other half-metal ferromagnets such as Heusler alloys [36] need to be used or the
interface between LSMO and organics needs to be engineered [37-39], such as by

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chemisorption [33].
4. Conclusion

In summary, we have studied the effects of SrTiO3 substrate-induced strain on the

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magnetic properties of 20 nm and 200 nm thick LSMO films and the MR responses of
LSMO/Alq3/Co spin valve devices using these films as ferromagnetic electrodes. We have

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observed that by reducing the LSMO film thickness from 200 nm to 20 nm, the compressive
strain was induced which drastically increased the bulk in-plane magnetic anisotropy but not
the surface in-plane magnetic anisotropy. Combined MOKE, VSM, TS and MR experiments
provide solid evidence that the in-plane magnetic anisotropy plays no decisive role in

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LSMO/Alq3/Co spin valve structures. Our study supports the hypothesis that the effective
surface spin polarization at the OSCs/LSMO interface dominates the MR temperature

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dependence, and that the spin diffusion length of most OSCs is temperature independent.
Acknowledgements

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Research at the University of South Florida was supported by the U.S. Department of Energy,
Office of Basic Energy Sciences, Division of Materials Sciences and Engineering under
Award No. DE-FG02-07ER46438 (structural, VSM and TS studies). Work at the University
of Georgia was supported by the Start-up Fund and the Faculty Research Grant (MOKE and
MR studies). Work at University of Science and Technology of China was supported by
National Natural Science Foundation of China and NBRPC (2015CB921201) (growth of thin
films).

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Figure captions

Figure 1. XRD patterns of (a) 200 nm and (b) 20 nm LSMO films. Inset of Fig. 1a show an
enlarged XRD peak portion for 200nm LSMO film, and inset of Fig. 1b shows a typical

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detector scan for the 20 nm LSMO film. Insets show the AFM images (3x3µm) of both films.
Figure 2. MOKE loops taken at 127 K and 215 K for the (a) 200 nm and (b) 20 nm LSMO
films. (c) Temperature dependence of coercive field (HC) derived from the MOKE loops for

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both samples. VSM loops taken at 50 K and 300 K for (d) 200 nm and (e) 20 nm LSMO films.
(f) Temperature dependence of coercive field (HC) derived from the VSM loops for both

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samples. The inset of Figure 2f shows the temperature dependence of HC for the 20 nm
LMSO film derived from MOKE and VSM, respectively.

Figure 3. Bipolar TS scans taken at different temperatures for the (a) 200 nm and (b) 20 nm
LSMO films. (c)TS scans taken at 20 K for both samples for comparison. (d) The temperature

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dependences of effective anisotropy field (HK).

Figure 4. (a) Schematic of the LSMO/Alq3/Co spin valve device used in this study. Typical
MR curves taken at 10 K of the spin valve device using (b) 200 nm LSMO film and (c) 20nm

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LSMO film. The blue dash line indicates the MR signal after deducing the AMR component

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orientated from the magnetic anisotropy of LSMO film itself. (d) Temperature dependence of
normalized balanced MOKE of 200 nm and 50 nm LSMO films, and temperature dependence
of normalized magnetization of 200 nm and 20 nm LSMO films. (e) Temperature dependence
of normalized MR of the OSV devices using different LSMO thicknesses. (f) Temperature
dependence of normalized MR of the devices using various organic materials as the spacers.
(Nguyen et al. [35]; Liang et al. [34]; Geng et al. [33])

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