# Structural defects improve the memristive characteristics of epitaxial La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3</sub> – δ based devices

Carlos Moncasi, Gauthier Lefèvre, Quentin Villeger, Laetitia Rapenne, Hervé

Roussel, Ahmad Bsiesy, Carmen Jiménez, Mónica Burriel

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Carlos Moncasi, Gauthier Lefèvre, Quentin Villeger, Laetitia Rapenne, Hervé Roussel, et al.. Structural defects improve the memristive characteristics of epitaxial La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3</sub> – δ based devices. Advanced Materials Interfaces, 2022, 2022, pp.2200498. 10.1002/admi.202200498. hal-03764463

HAL Id: hal-03764463

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Submitted on 31 Aug 2022

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Structural Defects Improve the Memristive Characteristics of Epitaxial La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3</sub>-based Devices

Carlos Moncasi, Gauthier Lefèvre, Quentin Villeger, Laetitia Rapenne, Hervé Roussel, Ahmad Bsiesy, Carmen Jiménez, Mónica Burriel*

C. Moncasi, Q. Villeger, L. Rapenne, H. Roussel, C. Jiménez, M. Burriel

Univ. Grenoble Alpes, CNRS, Grenoble INP, LMGP, 38000 Grenoble, France

E-mail: monica.burriel@grenoble-inp.fr

G. Lefèvre, A. Bsiesy

Univ. Grenoble Alpes, CNRS, CEA/LETI Minatec, LTM, 38000 Grenoble, France

Keywords: resistive switching, manganites, structural defects, memristive devices, metal-organic chemical vapor deposition (MOCVD), epitaxial thin films, valence change memories (VCMs)

Interface-type valence change memories (VCMs) are exciting candidates for multilevel storage in Resistive Random Access Memories (ReRAM) and as artificial synapses for neuromorphic computing. Several materials have been proposed as VCM candidates and, depending of the materials and electrodes of choice, different switching mechanisms take place leading to the change in resistance. Here, we focus on La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3-δ</sub> (LSM) perovskite and, particularly, on the role of its nanostructure on the memristive device performance. The nanostructural details of the layers are modified by growing LSM epitaxial thin films on different substrates, i.e. SrTiO<sub>3</sub> (STO) and LaAlO<sub>3</sub> (LAO), by metal-organic chemical vapor deposition (MOCVD).

An interface-type memristive response is observed using Ti as active electrode and Pt as inert electrode. The modifications in the nanostructure of LSM (strain and dislocations) determine the memristive performance, leading to differences in cycle to cycle reproducibility and multilevel capabilities by the modification of the LSM's oxygen migration properties. The results we report show that nanostructure engineering is a promising approach for optimizing the performance of memristive devices, approach which can also be extended and applied to other nanoionic electrochemical devices.
# 1. Introduction

The generation of data is rapidly scaling, while current memory technologies are approaching their physical limits. With the end of Moore's law in mind, industry and academia moved the focus to new approaches and alternative materials. In this scenario, resistive random access memories (ReRAM) are considered promising candidates for the next generation of non-volatile memories, due to their fast writing speed, high storage density, high endurance, long retention times and low energy consumption.[1] By applying an electric stimulus, these devices exhibit a reversible change of resistance retained under zero-field conditions, and hence high and low resistance states (HRS and LRS, respectively) can be defined. In recent years, a great effort has been carried out to explain the origin of the effective change in resistance, and a wide range of possible mechanisms have been reported in the literature.[1,2] Some of the switching mechanisms rely on purely electronic effects, such as charge-trapping at an interface, modification of a Schottky barrier, or changes induced by electronic charge in strongly correlated electron systems of oxides, which can lead to metal-insulator transitions (MIT). On the other hand, ionic transport (drift) has also been reported as a mechanism able to induce a change of resistance in memristive devices. In electrochemical metallization memories, cations from an active electrode migrate reversibly in an insulating matrix, reaching the inert electrode and leading to a change in resistance. In valence change memories (VCMs) with electrode/oxide/electrode configuration, where the oxide can be insulating, such as $HfO_2$,[3] $TaO_x$,[4,5] $SrTiO_3$,[6] or semiconducting,[7] such as $LaMnO_3$,[8] or $La_2NiO_4$,[9–11] the resistance is modified by the drift of oxygen ions (or oxygen vacancies) and the concomitant local valence change of the cationic sublattice. Nonetheless, in many cases electronic and ionic effects interplay to give raise to the memristive response.[12,13] The switching event in VCMs can take place at different scales and the memories can be classified as: (i) filamentary, where the change of resistance and migration of oxygen takes place in a very localized region, (ii) interface-type, where the mechanism occurs at the electrode/oxide interface,[2] or (iii) volume-type, which can be considered an extension of the interface-type where eventually the whole volume of the oxide is involved in the change of resistance.[14] Although so far most of the technological effort has been devoted to filamentary-type memories, these devices suffer from strong cycle-to-cycle and cell-to-cell variability. The variability has been associated to the electroforming step, and the stochastic nature of the formation/rupture of the conductive filament, being both greatly influenced by the homogeneity of the material (interfaces and defects).[15,16] In this sense, several approaches have been proposed to control the filament formation and improve the device performance.[17] On the other side, interface-type devices have gained recent interest, as
they often do not require an electroforming step thanks to an electric-field induced gradual modification of the conducting properties of the electrode/oxide interface, and due to the possibility of defining multiple resistance states in a single device.

The mechanisms behind the memristive response can be altered by varying the nature of the oxide material, as well as its morphology and topography, in addition to the metals used as electrodes.[1,18] Among the materials proposed, perovskites such as La₁-xSrxMnO₃-δ (LSM-x, being x the strontium substitution percentage, and δ the oxygen deficiency non-stoichiometry), are of particular interest as both their electronic and ionic transport properties can be tuned by small chemical changes. The resistivity in LSM-x strongly depends on the composition and the temperature. LSM-x is insulating (or semiconducting) in a broad range of temperature (0-500 K) for strontium substitution smaller than 17.5%, while metallic conduction occurs for Sr substitution of 30-40%. In the intermediate range, i.e. x= 0.175-0.3, a MIT takes places, in which the transition temperature increases with Sr content, i.e. 283 K for x=0.175 and 369 K for x=0.3.[19] MIT has also been reported for LSM-x thin films with transition temperatures of 335 K for x=0.2 and 350 K x=0.3.[20] The carrier density, and electrical resistivity in LSM-x is determined by the Mn³⁺/Mn⁴⁺ratio and it can be modified by a change of oxygen stoichiometry, or by the cationic stoichiometry, as reported for bulk[19,21,22] and thin films.[23,24] At a given temperature, the resistivity in LSM-x (x=0-50%) decreases with Sr and oxygen content.[19–26] Compared to stoichiometric LSM-33 with δ=0, a resistivity increase of around 3 orders of magnitude was reported at room temperature for bulk LSM-33 with δ=0.07,[25] and for LSM-33 thin films with δ=0.13.[26] In LSM-x thin films, the resistivity can be further tuned by the substrate used, as seen for epitaxial LSM-33.[27] While the oxygen content determines the resistivity of LSM-x, oxygen migration is crucial to trigger the change of resistance. As has been proved in the last few years, oxygen mobility at high temperature (typically >600 °C) in La₀.₈Sr₀.₂MnO₃-δ thin films is strongly influenced by the microstructure at the nanoscale, being largely enhanced along the space-charge tubes surrounding dislocations[28,29] and along grain boundaries.[30,31] Thus, thin film engineering has emerged as a key strategy to tune functional properties to improve performance in devices such as solid oxide fuel cells[30,31] and memristive devices.[29,32]

**Table 1** summarizes selected relevant studies where different mechanisms were proposed for LSM-x (x=20-33%) memristive devices using different electrodes or by direct contact with the biased probe. Whereas several mechanisms are suggested, many of them rely on oxygen (vacancy) drift to induce the resistance change. Hence, oxygen content and mobility within the lattice are likely to play a major role in the memristive properties.[20,33] An oxygen vacancy
concentration increase was ascribed as the origin of the resistance increase in LSM-20 and
LSM-30 achieved by volumetric resistive switching using conductive-AFM (C-AFM).[20] In a
different study, the memristive response in LSM-33 perovskite films was related to the
reversible formation of the highly resistive oxygen vacancy-ordered brownmillerite phase by
*in situ* TEM in vacuum conditions.[33,34]

As explained earlier, the memristive response in VCMs may depend of the electrodes used,[1]
especially as a function of the work function and oxygen affinity. When the p-type
semiconductor oxide LSM-x (x=0-50%)<sup>[21]</sup> is put in contact with a large work function metal,
such as Pt (5.9 – 6.2 eV<sup>[35]</sup>), the I(V) characteristics are expected to show ohmic behavior. This
metal is generally used as counter electrode in memristive devices<sup>[36-40]</sup> as it is considered inert.
On the other side, when LSM-x is in contact with a low work function metal, such as Ti
(4.3 eV<sup>[41]</sup>), a Schottky barrier is expected to be formed with rectifying behavior. Nonetheless,
the barrier can be further modified by interface states as demonstrated by Sawa et al.<sup>[37]</sup> by
comparing two metals with similar work function, Ag and Ti, in contact with p-type
Pr<sub>0.7</sub>Ca<sub>0.3</sub>MnO<sub>3</sub> (PCMO). The larger rectifying behavior of the Ti/PCMO interface, compared
to Ag/PCMO, was attributed to the larger oxygen affinity of Ti. The Ti electrode would extract
oxygen ions from PCMO and create oxygen vacancies next to the interface, which would
change the width and/or height of the Schottky barrier. Hence, the Schottky barrier can be
modified by an interface layer, either using a metal that can be oxidized, or introducing a
defined interlayer.<sup>[1,37,42]</sup> This interlayer would also act as an oxygen reservoir accommodating
the oxygen from the film by the formation of an oxygen deficient MO<sub>x</sub> (with x < 2 for M=Ti)
oxide.

Given the variety of deposition techniques, film thicknesses, electrode-material combinations
and operating conditions reported in the literature, a generalized evaluation of the role of the
nanostructure in the LSM memristive device properties is rather difficult. To focus on this point,
in this work we selected La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3-δ</sub> (LSM) as a memristive material to build planar
interface-type resistive switching devices. We choose Ti as active electrode with high oxygen
affinity for fabricating the devices, and substrates presenting different lattice mismatch with
LSM, namely SrTiO<sub>3</sub> (STO) and LaAlO<sub>3</sub> (LAO), to provoke an alteration of the nanostructural
details of the LSM layer by the induced strain. The impact of the differences in microstructure
at the nanoscale is presented and related to the large differences in memristive characteristics,
reproducibility and volatility measured.
<table><caption>**Table 1.** Summary of relevant resistive switching studies using La₁-xSrxMnO₃ (LSM-x, being x the strontium substitution percentage), for x=20-33%, highlighting the type of switching and the mechanism proposed. The term 'probe' corresponds to electrically biased probe in direct contact with the film, whereas the term 'tip' is restrained to conductive-AFM (C-AFM) or other SPM methods in which the cantilever is biased.</caption><thead><tr><th>Thin film</th><th>Texture</th><th>Substrate</th><th>Device configuration</th><th>Device</th><th>Type of switching</th><th>Mechanism proposed</th><th>Refs.</th></tr></thead><tbody><tr><td>LSM-20</td><td>Poly.<sup>a)</sup></td><td>Ir<sup>b)</sup></td><td>Vertical</td><td>Ag/LSM-20/Ir</td><td>Interface</td><td>VCM: change in δ</td><td>[43]</td></tr><tr><td>LSM-20</td><td>Poly.<sup>a)</sup></td><td>Ir<sup>b)</sup></td><td>Vertical</td><td>Tip/LSM-20/Ir</td><td>Interface</td><td>VCM: change in δ</td><td>[43]</td></tr><tr><td>LSM-20</td><td>Epitaxial</td><td>STO</td><td>Planar</td><td>CeO<sub>2-x</sub>/LSM-20/CeO<sub>2-x</sub></td><td>Volume</td><td>VCM: MIT</td><td>[42,44]</td></tr><tr><td>LSM-20</td><td>Epitaxial</td><td>STO</td><td>Planar</td><td>Probe/LSM-20/Probe</td><td>Volume</td><td>VCM: MIT</td><td>[14,42,44]</td></tr><tr><td>LSM-20</td><td>Epitaxial</td><td>STO</td><td>Planar</td><td>Tip/LSM-20 (grounded)</td><td>Volume</td><td>VCM: MIT</td><td>[14,20]</td></tr><tr><td>LSM-30</td><td>Epitaxial</td><td>STO</td><td>Planar</td><td>Tip/LSM-30 (grounded)</td><td>Volume</td><td>VCM: MIT</td><td>[13,20]</td></tr><tr><td>LSM-30</td><td>Epitaxial</td><td>LAO</td><td>Planar</td><td>Au/LSM-30/Au</td><td>Filamentary</td><td>Conducting multifilaments along dislocations</td><td>[45]</td></tr><tr><td>LSM-30</td><td>Epitaxial</td><td>SRO<sup>c)</sup>/STO</td><td>Vertical</td><td>Ti/SCMO<sup>d)</sup>/LSM-30/SRO<sup>c)</sup></td><td>Interface</td><td>Modification of Schottky barrier</td><td>[46]</td></tr><tr><td>LSM-33</td><td>Epitaxial</td><td>STO</td><td>Planar</td><td>Ag/LSM-33/Ag</td><td>Interface</td><td>Charge (de)trapping at insulating interfacial layer</td><td>[47]</td></tr><tr><td>LSM-33</td><td>Epitaxial</td><td>Nb:STO</td><td>Vertical</td><td>Probe/LSM-33/Nb:STO</td><td>Volume</td><td>VCM: topotactic phase transition</td><td>[33]</td></tr></tbody></table>

## 2. Results

### 2.1. Structural characterization

The orientation, strain and extended structural defects of a La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3-δ</sub> (LSM) thin films can be easily modified by using different perovskite single crystal substrates with lattice cell parameters close to those of the layer. In order to compare the resistive switching response of LSM memristive devices with two different microstructures at the nanoscale, 7 nm thick LSM

<footnote><sup>a)</sup>Polycrystalline; <sup>b)</sup>Ir: Ir/TiN/SiO<sub>2</sub>/Si; <sup>c)</sup>SRO: SrRuO<sub>3</sub>; <sup>d)</sup>SCMO: Sm<sub>0.7</sub>Ca<sub>0.3</sub>MnO<sub>3</sub>,</footnote>
thin films were grown simultaneously by PiMOCVD on (001)-oriented STO and on (001)ₚc-oriented LAO (pc=pseudo cubic) substrates, hereafter named as LSM/STO and LSM/LAO, respectively. The pseudo-cubic bulk lattice parameter of LSM is $a_{LSM, pc}=3.898$ Å.[⁴⁸,⁴⁹] The LSM films grown on STO are expected to grow epitaxial as a result of the similar lattice parameter ($a_{STO}=3.905$ Å), with very small in-plane tensile strain ($\epsilon = \frac{a_{STO}-a_{LSM, pc}}{a_{LSM, pc}} = +0.18\%$). On the other side, films grown on LAO, with a smaller lattice parameter ($a_{LAO, pc}=3.79$ Å), would exhibit in-plane compressive strain ($\epsilon = \frac{a_{LAO}-a_{LSM, pc}}{a_{LSM, pc}} = -2.77\%$) and concomitant expansion out-of-plane, as reported elsewhere.[²⁸] The calibration of the thin film composition as a function of the deposition conditions was the object of a preliminary study. Several films were grown on silicon substrates using precursor solutions with different cationic concentration. These films were analyzed by wavelength dispersive spectroscopy (WDS) and Rutherford Backscattering spectrometry (RBS) to determine their exact composition (detailed information can be found in the electronic supplementary information, ESI). The precursor solution cationic concentration that yielded the targeted composition of La₀.₈Sr₀.₂MnO₃-δ (LSM) was used to grow the epitaxial films presented here.

The XRD patterns of the LSM/STO and LSM/LAO films are shown in **Figure 1a** and **b**, respectively. These show high intensity peaks assigned to the reflections of the (00l) planes of the single crystal substrates, in addition to which only the characteristic (00l)ₚc reflections of LSM can be observed next to the substrate peaks. For the LSM/STO sample the film reflections overlap with the high intensity peaks of STO, while for the LSM/LAO sample, the peaks assigned to the film appear at smaller angles than for the bulk (drop lines in Figure 1a and b), in agreement with the expected in-plane compressive strain induced by the LAO substrate. The experimentally measured out-of-plane parameter obtained for the LSM film on LAO is 3.95 Å, which corresponds to an out-of-plane elongation of $\epsilon=+1.33\%$, related to a partial relaxation of film's strain, as discussed below.

The surface morphology of the films was studied by SEM and by AFM in tapping mode for LSM/STO and LSM/LAO as shown in **Figure 1c** and **d**, respectively. The as-grown films are homogeneous, dense and smooth, with root mean square roughness, $S_q$, of around 0.6 nm.
<figure><img src="image_2.png" /></figure>

<figure><img src="image_3.png" /></figure>

<figure><img src="image_4.png" /></figure>

<figure><img src="image_5.png" /></figure>

**Figure 1.** Structural characterization of the LSM films. XRD patterns of LSM films grown on a) STO and b) LAO. LSM and LAO peak are indexed in pseudo-cubic (pc) system, and STO peaks in cubic system. The films exhibit epitaxial growth. Black drop lines correspond to La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3</sub> bulk positions according to ICDD-00-053-0058. Surface imaging by SEM (left) and AFM height profile (right) of the c) LSM/STO and d) LSM/LAO films.

The crystal quality of the film was assessed by HRTEM after the fabrication of the devices. The epitaxial texture of the film is observed by the HRTEM cross-section images of LSM[001]/STO[001] and LSM[001]/LAO[001], Figure 2a and c, respectively. Figure 2a shows long-range domains of ordered atoms in the LSM film when grown on STO. However, the LSM film presents some extended structural defects when grown on LAO (Figure 2c), probably caused by the larger cell mismatch, compared to that on STO. The Fourier transform of the HRTEM images presented in Figure 2a and c (LSM/STO and LSM/LAO, respectively) is shown in the inset of Figure 2b and d, respectively, in which a mask is applied to the diffraction spots corresponding to (010)LSM and (010)LAO or STO which are superimposed. The inverse Fourier transform filtered image is presented in Figure 2b and d, where only the diffraction planes perpendicular to the interface are visible. In LSM/STO (Figure 2b), the planes are straight and extend from the bottom interface to the top, thus proving a good epitaxy with
absence of extended defects. On the contrary, the inverse Fourier transform image in LSM/LAO (Figure 2d), suggests the presence of extended defects in the LSM film appearing as discontinuity of the planes, extra planes, and deflection of planes, which may cause partial relaxation of the strain. Considering a fully relaxed LSM film the theoretical dislocation spacing can be calculated by $S = \frac{|b|}{\epsilon}$. Assuming a Burgers vector of one unit cell of the substrate, for example in the direction of [100], $b = a_{LAO}[100]$ ($|b|= 3.79$ Å), and a misfit $\epsilon = -2.77\%$, the dislocation spacing for a full relaxation is $S=13.7$ nm, which corresponds to an linear dislocation density of $7.3 \times 10^5$ cm$^{-1}$ and means one dislocation every ~ 36 unit cell of the LAO structure. In the image shown in Figure 2d a first dislocation is found at the left side of the image (marked with a "⊥" symbol, enlarged region inside a white square). In addition, a second complex defect, which could be formed by two dislocations of opposite sign cancelling each other (marked with "⊥" and "T" symbols), is found at the right side of the image, at a distance of 17-20 nm from the first one. We thus confirm that the number of dislocations observed in this region correspond well to the expected for a partially relaxed LSM/LAO film. The presence of extended defects on LSM/LAO was also reported by Navickas et al.[28], where they assigned them to threading and edge dislocations across and along the film. Other regions of LSM, both with and without top electrodes, were observed and exhibit similar film crystal quality as those presented for LSM/STO and LSM/LAO in Figure 2a and c, respectively. The average thickness of the films is 7 nm. Hence, we confirm that both films are epitaxial, with the cube-on-cube growth of the LSM perovskite on top of the substrate perovskite unit cells. However, while LSM/STO is epitaxial, presents no extended defects and a very small in-plane tensile strain, the LSM/LAO film shows indeed the presence of dislocations and in-plane compressive strain.

The presence of a $TiO_x$ interlayer between the LSM film and the Ti top electrode was observed by EDS mapping on TEM lamellae of the Pt/Ti/LSM/LAO (Figure 2e) and Pt/Ti/LSM/STO (Figure S2a) stacks using pristine devices. Figure 2f presents the chemical profile across the Pt/Ti/LSM/LAO stack. It should be noted that the slight Al concentration increase in the Pt layer most probably corresponds to a parasite signal coming from backscattered and high angle diffused electrons interacting with the aluminum washers used to fix the sample holder. Focusing first on the LSM film, we detect a slight accumulation of Mn at the bottom LSM/LAO interface, while the La, Sr and O profiles are quite homogeneous across the LSM film. Within the given resolution, no Sr segregation is observed at the top LSM surface (LSM/Ti interface), as often observed for Sr perovskites.[50–52] Oxygen clearly accumulates on the Ti/LSM interface, forming a $TiO_x$ interlayer of ~2-4 nm. On the LSM side of the Ti/LSM interface (green shaded
area in Figure 2f), the O/Ti ratio is roughly 2 and decreases to 1 after ~1 nm. Similar characteristics were found in the EDS maps and profile obtained for the Ti/LSM/STO stack (Figure S2a and b).

<figure><img src="image_6.png" /><figcaption>**Figure 2.** TEM cross-section characterization of LSM/STO and LSM/LAO films. HRTEM image of the 7 nm thick LSM film on a) STO and c) LAO. b), d) Fourier transform showing the reciprocal space of the images presented in a), b) respectively. Inverse Fourier transform image</figcaption></figure>
obtained applying a mask on the (010)LSM and (010)STO and LAO spatial frequencies on LSM/STO and LSM/LAO, highlighting the planes perpendicular to the surface. e) STEM-EDS mapping analysis of a pristine Ti/LSM/LAO stack with the presence of TiOx interlayer. f) Chemical profile across the Pt/Ti/LSM/LAO stack by integrating the signal in an area of 79x37 nm. Shaded areas correspond to the interfaces with LSM.

## 2.2. Electrical characterization

The Pt/LSM contact resistance and LSM resistivity obtained by transmission line measurements (TLM) are of the same order of magnitude for both samples, as summarized in **Table 2** (see ESI and **Figure S3** for more information).

<table><caption>**Table 2.** Results of Transmission line measurements (TLM) using Pt contacts.</caption><thead><tr><th>Sample</th><th>Contact resistance, R<sub>c</sub> [Ω]</th><th>Resistivity, ρ [Ω cm]</th></tr></thead><tbody><tr><td>LSM/STO</td><td>2.27×10<sup>4</sup></td><td>0.38</td></tr><tr><td>LSM/LAO</td><td>2.12×10<sup>4</sup></td><td>0.52</td></tr></tbody></table>

As shown in **Figure 3**, ohmic and rectifying responses were obtained in both samples for symmetric Pt/LSM\Pt and Ti/LSM\Ti planar devices, respectively. For the Pt/LSM\Pt devices ohmic behavior (a linear response) was obtained even up to ±20 V (see Figure S4), with a resistance value of 150 kΩ and 260 kΩ for LSM/STO and LSM/LAO, respectively. On the other hand, Ti/LSM\Ti devices present rectifying behavior with a resistance around 10⁶ Ω (at +0.1 V), and a hysteretic behavior in both polarities for large voltage ranges (see Figure S4). These responses are in agreement with the band alignment, confirming the passive role of the inert Pt electrode and the active role of the Ti electrode, inducing the memristive behavior. The different resistance values are coherent with the formation of a Schottky barrier, and the concomitant reduction of LSM and oxidation of Ti into an oxygen deficient TiOx layer, as observed by STEM-EDS.
<figure><img src="image_7.png" /></figure>

<figure><img src="image_8.png" /><figcaption>**Figure 3.** *I(V)* characteristics of symmetrical devices fabricated on a) LSM/STO and b) LSM/LAO. Symmetrical devices of Ti/LSM\Ti (dotted line) show rectifying behavior, and of Pt/LSM\Pt (solid line) show ohmic behavior.</figcaption></figure>

The memristive behavior was studied on asymmetric planar Pt/Ti/LSM\Pt devices, in which the voltage bias was applied to the Pt/Ti active electrode while the Pt electrode was grounded. The initial resistance state (IRS) of the pristine devices was measured (at +0.1 V) in two sets of devices: (i) devices with electrodes of different area, *A*, separated by 50 µm and, (ii) devices with different separation between electrodes, *L*, and with a fixed electrode surface *A* of 200×200 µm². The values obtained for the two sets of devices for LSM/STO and LSM/LAO are presented in **Figure 4**. Figure 4a and b show that the IRS is independent of the substrate and of the distance between contacts, *L*=50, 100 and 200 µm, showing a negligible contribution of the LSM sheet resistance to the total IRS measured. On the other hand, Figure 4c and d show the dependence of IRS with contact area. As a first approach the total resistance can be calculated as the sum of series resistances of each material and interface, i.e. $R_{Pt} + R_{C1} + R_{LSM} + R_{C2} + R_{Ti}$, being $R_{LSM} = R_S \frac{L}{W}$, where $R_S$ is the sheet resistance of the LSM film, *L* is the length between the electrodes, and *W* is the width of the electrodes. Using the sheet
resistance of LSM and the contact resistance between Pt/LSM ($R_{C1}$) obtained by TLM, it can be inferred that the $R_{C2}$ component governs the total resistance, being responsible of 61-76% and 71-97% of the total resistance for LSM/STO and LSM/LAO, respectively. The large contribution of the $R_{C2}$ component to the IRS is in good agreement with the formation of a Schottky barrier and a $TiO_x$ layer, as shown by STEM-EDS (Figure 2e) as well as the depletion of oxygen in the LSM film under Ti. Both these phenomena are expected to increase the $R_{C2}$ resistance contribution. The larger IRS distribution measured for LSM/LAO compared to LSM/STO could come from a non-homogeneous structural defect distribution within the LSM film, leading to a certain inhomogeneity of the $TiO_x$ layer in terms of oxidation level and thickness, as well as different gradient in the concentration of oxygen vacancies inside LSM.

<figure><img src="image_9.png" /></figure>

**Figure 4.** Initial resistance state (IRS) of the Ti/LSM|Pt pristine memristive devices. a), b) Distance and c), d) area dependence study on LSM/STO and LSM/LAO, respectively. The area of the electrode (A) and distance between electrodes (L) in the distance study is $A=200^2 \ \mu m^2$ and $L=50$, 100 and 200 $\mu m^2$, while for the area study is $A=100^2$, $200^2$, $300^2$, and $400^2 \ \mu m^2$ and $L=50 \ \mu m$. The insets in a) and c) show a schematic top-view of the electrode arrangement as function of L and A, respectively. Box range: 25th to 75th percentile. Whiskers: to 10th (lower
limit) and to 90th (upper limit) percentile. Median value: horizontal line in box. Mean: open data point. Outliers: solid data points. For plots c) and d), the dashed gray lines correspond to slope -1. In LSM/STO (LSM/LAO), the data of the distance dependence study was obtained from 15, 15, 30 (16, 16, 32) devices for L=50, 100 and 200 µm², respectively; and for the area dependence study it was obtained from 9 (12) devices for each size.

Memristive behavior could be obtained for the Pt/Ti/LSM\Pt planar devices of all sizes. Here we focus on devices with the same length between electrodes. All the devices exhibit clockwise log *I*(*V*) (or counter-eightwise *I*(*V*)) characteristics where the HRS RESET and LRS SET are programmed at positive and negative voltage, respectively. An initialization procedure was carried out to evaluate the operating voltage amplitude range of the LSM-based devices. The procedure consists of symmetric bipolar writing voltage sweeps, 0 V → +*V*<sub>write</sub> → 0 V → +*V*<sub>write</sub> → 0 V (defined as cycle), where *V*<sub>write</sub> was increased every five cycles in steps of 2 V up to 20 V, and then decreasing back to 2 V. Figure S5 presents the *I*(*V*) characteristics of the initialization process for LSM/STO and LSM/LAO for standard devices (A=200x200 µm² and L=50 µm). For these standard devices the highest applied voltage of 20 V corresponds to an electric field of 4 10⁵ V/m. However, as previously explained, the largest contribution to the device's total resistance comes from the *R*<sub>C2</sub> component, which corresponds to the depleted LSM (under the electrode) and the LSM/Ti (TiO<sub>x</sub>) interface, implying that the electric field drop will not be uniformly distributed between the electrodes, but concentrated at the region close to the Ti electrode. The resistance state reached is measured at +0.1 V after each half cycle. All the measured devices exhibit hysteretic behavior at large applied voltage, but not always at small voltage. This memristive behavior is observed without the need of a previous high-voltage electroforming step. Here we have considered that the devices either do not present memristive behavior or that their memristive response is unreliable when the HRS/LRS ratio at small voltage (+0.1 V) is smaller than 1.3 and/or the ratio is very variable, with no voltage and/or cycles dependence, usually due to noisy measurements. For LSM/STO, more than 75% of the devices of all sizes (100² to 400² µm²) showed memristive behavior, while in the case of LSM/LAO it was found to be more likely for the small devices (with >75% success for 100² and 200² µm² devices). On the other hand, reliable switching in large devices was only obtained for 20% (300² µm²) and 50% (400² µm²) of the cases.

As is clearly shown in **Figure 5** (standard devices, largest voltage range), the *R*(*V*) characteristics of the LSM-based devices fabricated on STO and LAO are different. Contrary to the LSM/LAO sample, the *R*-*I*(*V*) curves in LSM/STO are asymmetric and similar to those
reported for La0.8Sr0.2MnO3/STO by C-AFM.[20] **Figure 6a** and **b** presents for LSM/STO and LSM/LAO, respectively, the evolution of HRS and LRS obtained after each half-cycle as a function of the total number of cycles (bottom axis) and the writing voltage amplitude (top axis).
In LSM/STO devices, during the initialization process, represented by open symbols in Figure 6a, the HRS and LRS are clearly distinguishable from $V_{write}=12$ V, after an abrupt resistance increase occurring during the RESET process. At this voltage, the HRS/LRS ratio is 1.9. Then, the evolution of the resistance states with increasing writing voltage is mainly dominated by a HRS increase until a $V_{write}$ of 20 V, where the HRS/LRS ratio is around 3. Next, as the voltage amplitude is decreased (cycles 51 to 100 in Figure 6a), the HRS shows a staggered decrease, as expected for a multilevel capacity memory, but the LRS remains practically constant at $V_{write}$ smaller than 16 V, with HRS/LRS ratio of 2.2. The same procedure was carried out a second time (solid data points in Figure 6a). The resistance evolution is similar, however the final resistance states reached are higher than during the first complete sequence.
A similar increasing trend was observed in a third and fourth repetition of the complete cycling procedure (not shown). It can be thus concluded that the Pt/Ti/LSM\Pt planar devices on LSM/STO do not show reproducible behavior at high operating voltages (e.g. ±20 V), as the resistance states constantly increase by cycling. However, as will be later shown through pulsed measurements, these devices are able to reproducibly work at lower amplitude voltages (e.g. ±10 V).
<figure><img src="image_10.png" /></figure>

<figure><img src="image_11.png" /><figcaption>**Figure 5.** Resistance-voltage (-electric field) characteristics of initialized Ti/LSM\Pt standard devices fabricated on the a) LSM/STO and b) LSM/ LAO films showing gradual SET and RESET transitions.</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>**Figure 6.** Evolution of the HRS and LRS resistance states achieved after each cycle (bottom axis) and as a function of writing voltage (top axis) for standard devices on a) LSM/STO and b) LSM/LAO. The resistance was read at +0.1V after each half cycle. The devices are initialized during the first experiment at which the writing voltage amplitude is increased from 2 V to 20 V and decreased back (open data points). The experiment was carried out a second time (solid data points) to assess reproducibility.</figcaption></figure>

On the other hand, the R-I(V) characteristics of the Pt/Ti/LSM\Pt planar devices on LSM/LAO show a similar shape in both polarities (Figure 5b). Similar curves have been reported for GdBaCo₂O₅/LaNiO₃ devices grown on STO using Ag and Pt electrodes,[12] and for La₀.67Sr₀.33MnO₃/LAO using two Ag electrodes,[47] where the formation of an interfacial layer due to surface defects acts as trapping/detrapping center modifying the Schottky barrier and the resistance. For LSM/LAO devices, during the initialization the value of the HRS drastically increases at low voltages, with two well differentiated states (HRS/LRS=1.6) at 10 V, (Figure 6b, open symbols). The HRS/LRS ratio increases with voltage amplitude reaching a value of 5 at 20 V, for the device here presented. The increase in HRS/LRS ratio is due to a simultaneous increase in HRS and a decrease in LRS with increasing voltage amplitude. More interestingly, both states show a staggered evolution towards an intermediate resistance value
when the voltage amplitude is decreased. Repeating the same voltage sequence procedure, the
resistances states (solid symbols in Figure 6b) are perfectly reproducible. The devices on
LSM/LAO also show reproducible performance for large number of cycles and under alternate
varying voltage amplitude, i.e. 10 cycles at ±20 V, followed by 35 cycles at ±10 V and 10 cycles
at ±20 V, as shown in Figure S6. Hence, it can be concluded that devices on LSM/LAO show
multilevel capacity and reproducible performance both at high and low amplitude voltages.

In addition, pulsed measurements were also carried out on LSM/STO and LSM/LAO standard
memristive devices. The writing/read procedure consisted of alternative Vwrite=±10 V pulses of
60 ms duration, each followed by a read measurement. This voltage amplitude was selected
from the previous I(V) cycling measurements (see Figure 6), so that both the LSM/STO and
LSM/LAO devices would operate reversibly and in a reproducible manner. In addition, a study
of the influence of the pulse duration (60 vs 600 ms) was carried out by performing a sequence
of alternative Vwrite=±10 V pulses on LSM/STO and LSM/LAO initialized devices grown on
similar epitaxial LSM thin films (8 nm thick in this case). Figure 7a and b show the resistance
states read for LSM/STO and LSM/LAO devices, respectively, for 100 bipolar consecutive
pulses. Both devices show a stable HRS with cycling, whereas the LRS increases during the
first cycles and remains quite stable after around 10 repetitions. The HRS/LRS ratio for
LSM/STO and LSM/LAO at the end of the 100 consecutive pulses (Vwrite=±10 V) is around 1.3
and 3, respectively. These ratios are almost two times smaller than the values obtained by
performing I(V) voltage sweeps at the same voltage amplitude (i.e. 0 V → +10 V → 0 V → -
10 V → 0 V). Figure S7 a and b show the resistance states read for LSM/STO and LSM/LAO
devices (8 nm thick LSM), respectively, for 128 bipolar consecutive pulses. For LSM/LAO
stable HRS and LRS values are already obtained after a single pulse of 60 (or 600 ms) duration,
while for the LSM/STO device about 20 pulses of 60 ms (or 8 pulses of 600 ms) duration are
needed to reach the stable values. In addition, for the LSM/STO device an increase of the
resistance states with pulse width is observed, while for the LSM/LAO device the HRS does
not vary its value for longer pulse duration, while the LRS does seem to slightly decrease.
<figure><img src="image_13.png" /><figcaption>**Figure 7.** Pulse operation of the Pt/Ti/LSM\Pt memristive devices working using ±10 V/60 ms writing voltage pulses. Evolution of HRS and LRS at +0.1 V over 100 consecutive bipolar pulses for devices on a) LSM/STO and b) LSM/LAO.</figcaption></figure>

Next, the stability of the resistance states over time, i.e. their retention, was evaluated for two standard pristine devices on LSM/STO and on LSM/LAO. The devices were first initialized, and then cycled 5 times at ±20 V before programming the desired resistance state by a voltage sweep. The *R(V)* characteristics and HRS/LRS values are representative of each sample. The HRS/LRS ratio prior to the HRS and LRS retention measurements is 3.1 and 2.4, respectively, for LSM/STO devices, and 23 and 18, respectively, for LSM/LAO devices. Next, the evolution of the resistance with time, measured at +0.1 V for five hours, is presented in **Figure 8**. For both devices, the resistance states present certain volatility before stabilizing. The resistance state relaxation curves were fitted to a two-phase exponential decay. For the LRS relaxation, similar time constants were obtained for both samples, i.e. τ₁ ~ 10 min and τ₂ ~ 3 h, which suggest that the same phenomena could be taking place in both devices. However, the HRS resistance decay in the LSM/STO device is somewhat faster than in the LSM/LAO device, with time constants of τ₁ ~ 8 min, τ₂ ~ 1 h for LSM/STO devices, and τ₁ ~ 20 min, τ₂ ~ 6 h for LSM/LAO devices. Furthermore, as shown in Figure S8, the stability of four different multi-
states has been confirmed for the LSM/LAO optimized devices, corresponding to two HRS states (programmed at +10 V and at +20V) and to two LRS (programmed at -10 V and at -20V).

<figure><img src="image_14.png" /></figure>

**Figure 8.** HRS and LRS resistance state relaxation with time for Ti/LSM\Pt memristive devices programmed at ±20 V on LSM/STO (solid line) and LSM/LAO (dashed line). The resistance is measured at +0.1 V.

Finally, in order to evaluate the area dependence of the switching of the multilevel and reproducible Ti/LSM\Pt devices in LSM/LAO, the HRS and the LRS was measured for devices with different contact area ranging from 100$^2$ to 400$^2$ µm$^2$. Both states were programmed at different writing voltages, as presented in **Figure 9**, showing a clear resistance decrease by increasing the pad size, as expected for interface-type switching. Whereas the HRS shows a linear dependence with area and a slope of -1, the LRS shows a weaker dependence, which might be related to a non-homogeneous tunneling effect in the case of the larger size devices. The electroforming-free nature of the devices, together with the HRS and LRS dependence with device area, and the planar geometry of the device strongly support the existence of interface-type resistive switching (non-filamentary).[18]
<figure><img src="image_15.png" /><figcaption>**Figure 9.** Area dependence resistance of the HRS (blue shade, top) and LRS (grape shade, bottom) at different writing voltages for Ti/LSM\Pt devices in LSM/LAO. Only one device is presented for each area.</figcaption></figure>

**3. Discussion and oxygen switching model**

The Pt/Ti/LSM\Pt memristive devices fabricated on LSM/STO and LSM/LAO showed clockwise log *I*(*V*) characteristics with a gradual change in resistance, showing that the HRS and LRS can be tuned by the writing voltage amplitude. As has been widely reported for PCMO devices, the switching response could be triggered by the Ti contact acting as oxygen reservoir via the formation of an insulating oxygen deficient TiOx interlayer (see Figure 2e-f) capable of modifying the conducting interface.[46] This interlayer would facilitate oxygen exchange from (to) the LSM films with the concomitant formation (annihilation) of oxygen vacancies. This hypothesis is in agreement with the drift of negatively charged oxygen ions towards (away from) the Ti electrode under positive (negative) voltage during HRS RESET (LRS SET) applied to the Ti electrode (see sketch in **Figure 10**). For the initialized LSM/LAO sample, the resistance states are very reproducible and pivot around an intermediate resistance value (Figure 6b) with HRS/LRS ratios between 3 and 23. Therefore, the multilevel resistances could be related with the gradual oxidation (reduction) of TiOx and increasing (decreasing) oxygen vacancy concentration in LSM for the HRS (LRS). This mechanism is fully reversible under symmetric operating voltage. However, it can be observed that the LSM/STO devices only present resistive switching behavior after a relatively high RESET process (at + 12V). We
suggest that the large increase in resistance occurring at this stage is related to the formation of oxygen vacancies in LSM, where the oxygen ions would be incorporated in the TiOₓ interlayer (with increase of thickness or further oxidation). The presence of oxygen vacancies in LSM would allow for the resistive switching behavior with two well differentiated HRS and LRS states, which occurs thereafter. This observation highlights the importance of the initialization step in creating oxygen vacancies to trigger the switching. This hypothesis is supported by the work of Luo et al.[⁴⁰], where resistive switching in Pt/LSM20\Pt devices was only obtained for reduced films (with presence of oxygen vacancies), while no switching and ohmic behavior was obtained for oxygen stoichiometric LSM20-based devices. In addition, smaller ratios are measured for the LSM/STO devices, where the resistance states are only reproducible at intermediate voltage ranges (up to 10 V), while at higher voltages they both states increase with operation (Figure 6a) due to the continuous increase of the LRS. This suggests that for the LSM/STO devices the HRS reached at large operating voltage cannot be fully reversed by applying the same voltage of opposite polarity to obtain the LRS. The differences observed between both samples can be explained by differences in oxygen mobility. As a first approximation we compare the high-temperature oxygen diffusion coefficients reported in the literature for LSM, although for our devices the voltage (field) would be the driving force, resulting in an oxygen drift instead of thermal diffusion. Similar bulk oxygen diffusion coefficients have been mesured at 600 °C for LSM films around 130 nm thick grown on LAO and STO.[²⁸,²⁹] Nonetheless, the presence of space-charge tubes around dislocations has been reported to be the cause of the enhanced oxygen diffusion along these defects,[²⁹] as experimenally measured by other groups along epitaxial thin films[²⁸] and along grain boudaries.[³⁰,³¹] In terms of oxygen migration, a device in LRS would RESET to HRS by the drift of oxygen towards the positively biased Ti electrode as a result of the electrical bias.[³³] This operation is favored in the case of LSM/LAO due to the presence of many extended defects, such as dislocations, with a high concentration of oxygen vacancies, that would act as fast oxygen diffusion paths (see Figure 10d-f), leading to a larger HRS/LRS window. This hypothesis is supported by the differences in resistive switching pulse-width dependence observed for devices grown on the two substrates (see Figure S7). On the other hand, during LRS SET process, the re-incorporation of oxygen and drift within LSM might be limited for the LSM/STO devices due to the high quality of the epitaxial LSM film, which does not present many extended defects (Figure 10b-c).[²⁸,²⁹,³¹] Eventually, the concentration of oxygen in TiOₓ and oxygen vacancies in LSM/STO will increase with cycles raising the total resistance of the device.
<figure><img src="image_16.png" /><figcaption>**Figure 10.** a) Sketch of a Pt/Ti/LSM\Pt memristive device for LSM/STO films, and d) sketch of the half-loops dislocations present in LSM/LAO films, in agreement to reference.[53] Cross-section scheme of the oxygen drift at the Ti/LSM interface in the memristive devices during operation at large voltage (i.e. ±20 V). Scheme of oxygen drift during HRS and LRS programming b), c) in LSM/STO- and d), e) LSM/LAO-based devices. In all cases the bias is applied on the Pt/Ti electrode while the Pt electrode is grounded. We note that the extended defects may present different lengths and orientations, and do not necessarily always connect both electrodes as it may be derived from Figure 10d.</figcaption></figure>

As for the relaxation process, it was measured at +0.1 V and thus an oxygen drift is not expected. However, a short-scale transient fickian diffusion and/or a redox reaction at the interface may occur as a result of the chemical gradient induced during operation. The increase of resistance of the LRS can be considered as either a small reincorporation of oxygen from the LSM back to the Ti via a redox reaction, or a gradient re-equilibration within the layers. These phenomena would be similar for both types of devices, LSM/STO and LSM/LAO, as similar time constants were found. Similarly, the relaxation of the HRS shows a decay of resistance, in this case faster in LSM/STO devices, that could be related to the back-exchange of oxygen from $TiO_x$ into the LSM film or a re-equilibration within the layers. The differences between substrates in this case might be related to local differences in stoichiometry at close proximity of the interface after the application of the RESET voltage.

**4. Conclusion**

Epitaxial thin films of $La_{0.8}Sr_{0.2}MnO_{3-δ}$ with different nanostructural details have been grown by pulsed injection metal-organic chemical vapor deposition on $SrTiO_3$ and $LaAlO_3$ perovskite
substrates. The negligible strain present in the LSM/STO films and the substrate-induced strain retained in LSM/LAO, leading to an out-of-plane cell enlargement of +1.33% was corroborated by X-ray diffraction and HRTEM. LSM/STO showed perfect epitaxial growth with long-range atomic order and practically no defects, while LSM/LAO showed the presence of extended defects, such as dislocations. By microscopy and spectroscopy techniques, the spontaneous formation of TiOₓ at the Ti/LSM interface was measured in pristine devices, likely resulting from oxygen-scavenging of the Ti electrode and the concomitant oxygen vacancy formation in the LSM film, which introduces a significant contribution to the initial resistance value.

The gradual clockwise log I(V) characteristics, and the HRS and LRS electrode size dependence support the existence of interface-type switching, where the tunable resistance states programmable at different writing voltages would be related to different levels of oxidation and reduction of LSM and Ti in these VCM memristive devices. Reproducible performance is achieved in both types of devices by operating at intermediate voltages, i.e. ±10 V, as shown by pulsed measurements over at least 100 cycles. HRS RESET occurs when a positive voltage is applied to the Ti electrode, and hence oxygen drifts from LSM toward the Ti electrode, which is oxidized. LRS is obtained under negative polarization, where oxygen drifts back into “bulk” LSM. While the films resistivity, topography and morphology were similar in both samples, the Ti/LSM\Pt planar devices showed substrate-dependent memristive behavior, particularly for high voltage operation ranges (± 20 V). Whereas LSM/STO shows an increase of the LRS with continuous cycling under bipolar symmetrical operation, LSM/LAO exhibits reproducible performance and larger operating windows. The presence of extended defects, such as dislocations, observed by HRTEM in LSM/LAO, would allow faster oxygen drift through LSM. Instead, the absence of extended defects in LSM/STO would hinder the drift of oxygen into the LSM, resulting in an incomplete SET to intermediate resistance states, leading to a gradual increase of the LRS. As for the retention, both states experience a small relaxation that might be related to an oxygen back-diffusion process at the Ti/LSM interface as a consequence of the chemical gradient induced during operation. Here, we have highlighted the importance of the nanostructure in epitaxial LSM-based memristive devices with gradual change in resistance and multilevel states. Based on these results, we propose their suitability as memristive devices and neuromorphic computing elements to be evaluated by future electrical measurements, including endurance studies, on designed vertical LSM-based memristive devices with controlled level of structural defects. The implications of these results could also be extended to other perovskite systems and technologies, such as micro solid oxide fuel cells and electrolysisers. Hence, engineering the nanostructure by introducing extended structural defects which enhance the
oxygen mobility can result in an effective way of improving the performance of different
nanoionic electrochemical devices.

**5. Experimental Section/Methods**

*Material Deposition:* Epitaxial 7 nm thick La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3-δ</sub> (LSM) films were grown by pulsed injection metal-organic chemical vapor deposition (Pi-MOCVD) in aJIPELEC reactor described elsewhere.[54,55] A homogeneous solution of metal-organic precursors was prepared using La(thd)<sub>3</sub>, Sr(thd)<sub>2</sub>-tetraglyme, and Mn(thd)<sub>3</sub>, [thd=2,2,6,6,-tetramethyl-3,5-heptanedionato] from Strem Chemicals dissolved in m-xylene from Alfa Aesar. The cationic concentration in the solution was optimized to grow films with 20% strontium substitution and [(La+Sr)/Mn] ratio close to 1. The total concentration of the solution was 22.5 mM, with [(La+Sr)/Mn] precursor ratio of 2 and [Sr/(La+Sr)] precursor ratio of 0.55.

The films were deposited on 1x1 cm² single crystals of LaAlO₃ (100)pc (LAO) and SrTiO₃ (100) (STO) substrates purchased from Crystec. For this the solution was evaporated at 250 °C and the reaction chamber was heated up at 675°C. The deposition pressure was 5 Torr and oxygen was added into the chamber with a partial pressure of 2 Torr. The gas flow rates were kept constant during deposition at 0.3 slm for O₂ and 0.4 slm for Ar. The liquid precursor solution was injected at 1 Hz with an opening time of 2 ms for 3600 pulses at a feeding rate of approximately 0.11 mL/min using argon as carrier gas.

*Characterization:* The orientation of the film and presence of crystalline phases were studied by XRD in Bragg-Brentano θ-2θ configuration in a Bruker D8 Advance serie II diffractometer with Cu Kα1 radiation (λ=1.5406 Å) equipped with rotating holder. The instrumental shift and displacement was corrected by aligning the STO and LAO substrate peaks to their reference, ICDD-00-035-0734 and ICDD-00-031-0022, respectively. SEM was performed in a FEG ZEISS GeminiSEM 300 to study surface morphology. AFM topography scans were performed in a Bruker Icon AFM in tapping mode to analyze the grain size and roughness of the film. The nanostructural details of the film were analyzed by TEM-EDS on electron transparent lamellas obtained by Focused Ion Beam milling using a FEI dual-beam Helios 450S. A platinum layer was sputtered to prevent damage from the ion beam (rough and fine milling at 30 kV and 2-8 kV operating voltage, respectively). The lamellas were observed at 200 kV in a Tecnai Osiris FEI microscope equipped with a Super-X detector for Energy Dispersive X-ray (EDS) spectroscopy (for the LSM/LAO sample). Complementary, cross-sectioned TEM samples were prepared by automated polishing, the latter using a MultiPrepTM system (Allied High Tech Products Inc.) The final polishing was performe using a felt-covered disc impregnated with a
silica solution until the appearance of the first extinction fringe among those of equal thickness. Ar-milling was then used to reduce the total thickness. These samples were observed at 200 kV in a JEOL JEM 2010 LaB₆ TEM (with a resolution of ~0.19 nm). Further STEM EDS measurements were collected by STEM with a JEOL 2100F FEG microscope operating at 200 kV with a 0.2 nm resolution in scanning mode (for the LSM/STO sample). The cationic composition of the film was analyzed by Wavelength Dispersive Spectroscopy (WDS) and by Rutherford Backscattering Spectrometry (RBS) performed at the Centro de Micro-Análisis de Materiales – Universidad Autónoma de Madrid (Spain). WDS was performed using a CAMECA SX50 spectrometer in different areas of each sample at different energies in the range 10–16 KeV. The data was analyzed using the software STRATAGem. RBS measurements were performed using the 5 MeV tandetron accelerator. The cations were analyzed with a 2 MeV He⁺ beam and silicon surface barrier detectors placed at a scattering angle of 170°. The analysis of O and C was carried out by 4He⁺ ion beam at oxygen and carbon resonant frequency, 3.035 MeV and 4 MeV, respectively. A charge of 12 µC was used to obtain each spectra. SIMNRA code was used to simulate the spectra and obtain the elemental composition.

*Device Fabrication and Characterization:* Device electrical characterization was carried out on planar devices fabricated in the PTA cleanroom facilities (Grenoble). A specific mask was designed to include: (i) facing Ti and Pt square electrodes of different sizes (100², 200², 300² and 400² µm²) separated 50 µm; (ii) 200×200 µm² facing Ti and Pt electrodes separated at different distances (50, 100 and 200 µm); and (iii) 2000x200 µm² Pt electrodes at different distances for transmission line measurements (TLM). The pattern was transferred to the samples by laser-assisted lithography process (µPG 101 from HEIDELBERG Instruments) and electron-beam induced metal evaporator (MEB550 from PLASSYS). Prior to the metal evaporation the exposed patterned surface of the sample was treated with a mild Ar⁺ etch to enhance metal adhesion and remove potential residues of the photoresist. Pt (approx. 100 nm thick) was used as counter electrode, and Ti (approx. 10 nm) capped with Pt (approx. 90 nm) to prevent atmospheric oxidation was selected as active electrode. The electrodes were evaporated in high vacuum conditions on top of each LSM thin film sample, forming planar devices. Devices formed by 200×200 µm² electrodes and separated by 50 µm are referred to as standard devices.

The current-voltage (I(V)) characteristics in voltage-control mode were analyzed under ambient conditions using a Keithley 4200A-SCS semiconductor characterization system equipped with pre-amplifiers modules. Two tungsten probes (tip radius 10 µm) were used to contact the
devices. All measurements were performed applying voltage on Pt/Ti electrode and grounding
the Pt counter electrode. The bipolar sweeps followed the sequence: 0 V → +V<sub>amp</sub> → 0 V → -
V<sub>amp</sub> → 0 V. After each half-sweep the resistance state values were measured at V<sub>read</sub>=+0.1 V
in a sweep from 0 V to +0.1 V (~10 mV/s).

**Supporting Information**

Supporting Information is available from the Wiley Online Library or from the author.

**Acknowledgements**

This work has been supported by the ANR funded project "Mangaswitch" (ANR-17-CE24-0038) and has benefited from characterization equipment of the Grenoble INP - CMTC platform supported by the Centre of Excellence of Multifunctional Architecture Materials "CEMAM" n°ANR-10-LABX-44-01 funded by the "Investments for the Future" Program, and from the "Plateforme Technologique Amont" de Grenoble, with the financial support of the "Nanosciences aux limites de la Nanoélectronique" Fundation" and CNRS Renatech network. In addition, part of this work, carried out on the Platform for Nanocharacterisation (PFNC), was supported by the "Recherches Technologiques de Base" program of the French National Research Agency (ANR). The authors thank M. Manso-Silván and Vicente Torres-Costa for their help with the RBS measurements, and Eirini Sarigiannidou for fruitful discussions on the TEM imaging and dislocation quantification.

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Published online: ((will be filled in by the editorial staff))

**References**

[1] R. Waser, R. Dittmann, C. Staikov, K. Szot, *Adv. Mater.* **2009**, *21*, 2632.

[2] S. Bagdzevicius, K. Maas, M. Boudard, M. Burriel, *J. Elec* **2017**, *157*.

[3] B. Govoreanu, G. S. Kar, Y. Y. Chen, V. Paraschiv, S. Kubicek, A. Fantini, I. P. Radu, L. Goux, S. Clima, R. Degraeve, N. Jossart, O. Richard, T. Vandeweyer, K. Seo, P. Hendrickx, G. Pourtois, H. Bender, L. Altimime, D. J. Wouters, J. A. Kittl, M. Jurczak, *Tech. Dig. - Int. Electron Devices Meet.* *IEDM 2011*, 729.

[4] M. J. Lee, C. B. Lee, D. Lee, S. R. Lee, M. Chang, J. H. Hur, Y. B. Kim, C. J. Kim, D. H. Seo, S. Seo, U. I. Chung, I. K. Yoo, K. Kim, *Nat. Mater.* **2011**, *10*, 625.
[5] F. Miao, J. P. Strachan, J. J. Yang, M. X. Zhang, I. Goldfarb, A. C. Torrezan, P. Eschbach, R. D. Kelley, G. Medeiros-Ribeiro, R. S. Williams, *Adv. Mater.* **2011**, *23*, 5633.

[6] K. Szot, W. Speier, G. Bihlmayer, R. Waser, *Nat. Mater.* **2006**, *5*, 312.

[7] R. Dittmann, J. P. Strachan, *APL Mater.* **2019**, *7*, DOI 10.1063/1.5129101.

[8] B. Meunier, E. Martinez, R. Rodriguez-Lamas, D. Pla, M. Burriel, C. Jimenez, Y. Yamashita, O. Renault, *ACS Appl. Electron. Mater.* **2021**, *3*, 5555.

[9] K. Maas, E. Villepreux, D. Cooper, C. Jiménez, H. Roussel, L. Rapenne, X. Mescot, Q. Rafhay, M. Boudard, M. Burriel, *J. Mater. Chem. C* **2020**, *8*, 464.

[10] K. Maas, E. Villepreux, D. Cooper, E. Salas-Colera, J. Rubio-Zuazo, G. R. Castro, O. Renault, C. Jimenez, H. Roussel, X. Mescot, Q. Rafhay, M. Boudard, M. Burriel, *Adv. Funct. Mater.* **2020**, *30*, 1.

[11] K. Maas, C. Wulles, J. M. Caicedo, B. Ballesteros, V. Lafarge, J. Santiso, M. Burriel, *J. Mater. Chem. A* **2022**, DOI 10.1039/d1ta10296f.

[12] S. Bagdzevicius, M. Boudard, J. M. Caicedo, L. Rapenne, X. Mescot, R. Rodríguez-Lamas, F. Robaut, J. Santiso, M. Burriel, *J. Mater. Chem. C* **2019**, *7*, 7580.

[13] C. Moreno, C. Munuera, S. Valencia, F. Kronast, X. Obradors, C. Ocal, *Nano Lett.* **2010**, *10*, 3828.

[14] J. C. Gonzalez-Rosillo, R. Ortega-Hernandez, J. Jareño-Cerulla, E. Miranda, J. Suñe, X. Granados, X. Obradors, A. Palau, T. Puig, *J. Electroceramics* **2017**, *39*, 185.

[15] D. J. Wouters, Y. Y. Chen, A. Fantani, N. Raghavan, in *Resist. Switch. From Fundam. Nanoionic Redox Process. to Memristive Device Appl.* (Eds: D. Ielmini, R. Waser), **2016**, pp. 597–622.

[16] M. Lanza, H. S. P. Wong, E. Pop, D. Ielmini, D. Strukov, B. C. Regan, L. Larcher, M. A. Villena, J. J. Yang, L. Goux, A. Belmonte, Y. Yang, F. M. Puglisi, J. Kang, B. Magyari-Köpe, E. Yalon, A. Kenyon, M. Buckwell, A. Mehonic, A. Shluger, H. Li, T. H. Hou, B. Hudec, D. Akinwande, R. Ge, S. Ambrogio, J. B. Roldan, E. Miranda, J. Suñe, K. L. Pey, X. Wu, N. Raghavan, E. Wu, W. D. Lu, G. Navarro, W. Zhang, H. Wu, R. Li, A. Holleitner, U. Wurstbauer, M. C. Lemme, M. Liu, S. Long, Q. Liu, H. Lv, A. Padovani, P. Pavan, I. Valov, X. Jing, T. Han, K. Zhu, S. Chen, F. Hui, Y. Shi, *Adv. Electron. Mater.* **2019**, *5*, 1.

[17] T. Li, H. Yu, S. H. Y. Chen, Y. Zhou, S. T. Han, *J. Mater. Chem. C* **2020**, *8*, 16295.

[18] A. Sawa, *Mater. Today* **2008**, *11*, 28.

[19] A. Urushibara, Y. Moritomo, T. Arima, A. Asamitsu, *Phys. Rev. B* **1995**, *51*, 103.
[20] J. C. Gonzalez-Rosillo, S. Catalano, I. Maggio-Aprile, M. Gibert, X. Obradors, A. Palau, T. Puig, *Small* **2020**, *16*, 1.

[21] J. Mizusaki, Y. Yonemura, H. Kamata, K. Ohyama, N. Mori, H. Takai, H. Tagawa, M. Dokiya, K. Naraya, T. Sasamoto, H. Inaba, T. Hashimoto, *Solid State Ionics* **2000**, *132*, 167.

[22] J. Mitchell, D. Argyriou, C. Potter, D. Hinks, J. Jorgensen, S. Bader, *Phys. Rev. B - Condens. Matter Mater. Phys.* **1996**, *54*, 6172.

[23] Y. Shen, C. Wang, L. Li, Q. Shen, L. Zhang, *Key Eng. Mater.* **2014**, 602–603, 902.

[24] S. Jiang, G. Wang, X. Ma, *Vacuum* **2015**, *120*, 32.

[25] Y. Baikov, E. I. Nikulin, B. T. Melekh, V. M. Egorov, *Phys. Solid State* **2004**, *46*, 2086.

[26] J. R. Sun, H. W. Yeung, H. K. Wong, T. Zhu, B. G. Shen, *Eur. Phys. J. B* **2003**, *35*, 481.

[27] S. K. Chaluvadi, F. Ajejas, P. Orgiani, S. Lebargy, A. Minj, S. Flament, J. Camarero, P. Perna, L. Méchin, *J. Phys. D. Appl. Phys.* **2020**, *53*, DOI 10.1088/1361-6463/ab8e7b.

[28] E. Navickas, Y. Chen, Q. Lu, W. Wallisch, T. M. Huber, J. Bernardi, M. Stöger-Pollach, G. Friedbacher, H. Hutter, B. Yildiz, J. Fleig, *ACS Nano* **2017**, *11*, 11475.

[29] J. M. Börgers, J. Kler, K. Ran, E. Larenz, T. E. Weirich, R. Dittmann, R. A. De Souza, *Adv. Funct. Mater.* **2021**, 2105647, 1.

[30] E. Navickas, T. M. Huber, Y. Chen, W. Hetaba, G. Holzlechner, G. Rupp, M. Stöger-Pollach, G. Friedbacher, H. Hutter, B. Yildiz, J. Fleig, *Phys. Chem. Chem. Phys.* **2015**, *17*, 7659.

[31] A. M. Saranya, D. Pla, A. Morata, A. Cavallaro, J. Canales-Vázquez, J. A. Kilner, M. Burriel, A. Tarancón, *Adv. Energy Mater.* **2015**, *5*, 1.

[32] D. Pla, C. Jimenez, M. Burriel, *Adv. Mater. Interfaces* **2017**, *4*, 1600974.

[33] L. Yao, S. Inkinen, S. Van Dijken, *Nat. Commun.* **2017**, *8*, 1.

[34] L. Yao, S. Inkinen, H. P. Komsa, S. van Dijken, *Small* **2021**, *17*, 1.

[35] G. N. Derry, Z. Ji-zhong, *Phys. Rev. B* **1989**, *39*, 1940.

[36] B. Meunier, E. Martinez, R. Rodriguez-Lamas, D. Pla, M. Burriel, M. Boudard, C. Jiménez, J. P. Rueff, O. Renault, *J. Appl. Phys.* **2019**, *126*, DOI 10.1063/1.5125420.

[37] A. Sawa, T. Fujii, M. Kawasaki, Y. Tokura, *Appl. Phys. Lett.* **2004**, *85*, 4073.

[38] T. Liu, M. Verma, Y. Kang, M. Orlowski, *Appl. Phys. Lett.* **2014**, *073510*, 1.

[39] L. Goux, P. Czarnecki, Y. Y. Chen, L. Pantisano, X. P. Wang, R. Degraeve, B.
Govoreanu, M. Jurczak, D. J. Wouters, L. Altimime, *Appl. Phys. Lett.* **2010**, *97*, DOI 10.1063/1.3527086.

[40] X. D. Luo, R. L. Gao, C. L. Fu, W. Cai, G. Chen, X. L. Deng, H. R. Zhang, J. R. Sun, *Phys. B Condens. Matter* **2016**, *483*, 99.

[41] F. Borgatti, C. Park, A. Herpers, F. Offi, R. Egoavil, Y. Yamashita, A. Yang, M. Kobata, K. Kobayashi, J. Verbeeck, G. Panaccione, R. Dittmann, *Nanoscale* **2013**, *5*, 3954.

[42] R. Ortega-Hernandez, M. Coll, J. Gonzalez-Rosillo, A. Palau, X. Obradors, E. Miranda, T. Puig, J. Suñe, *Microelectron. Eng.* **2015**, *147*, 37.

[43] X. Chen, N. Wu, J. Strozier, A. Ignatiev, *Appl. Phys. Lett.* **2006**, *89*, 2004.

[44] J. C. Gonzalez-Rosillo, R. Ortega-Hernandez, B. Arndt, M. Coll, R. Dittmann, X. Obradors, A. Palau, J. Suñe, T. Puig, *Adv. Electron. Mater.* **2019**, *5*, 1800629.

[45] H. K. Lau, C. W. Leung, *J. Appl. Phys.* **2008**, *104*, DOI 10.1063/1.3043801.

[46] A. Sawa, T. Fujii, M. Kawasaki, Y. Tokura, *Appl. Phys. Lett.* **2006**, *88*, 86.

[47] Y. W. Xie, J. R. Sun, D. J. Wang, S. Liang, B. G. Shen, *J. Appl. Phys.* **2006**, *100*, 1.

[48] T. Grande, J. R. Tolchard, S. M. Selbach, **2012**, DOI 10.1021/cm2030608.

[49] M. Gaudon, C. Laberty-Robert, F. Ansart, P. Stevens, A. Rousset, *Solid State Sci.* **2002**, *4*, 125.

[50] B. Koo, K. Kim, J. K. Kim, H. Kwon, J. W. Han, W. C. Jung, *Joule* **2018**, *2*, 1476.

[51] H. Jalili, J. W. Han, Y. Kuru, Z. Cai, B. Yildiz, *J. Phys. Chem. Lett.* **2011**, *2*, 801.

[52] H. S. Lee, C. S. Park, H. H. Park, *Appl. Phys. Lett.* **2014**, *104*, 0.

[53] E. Navickas, Y. Chen, Q. Lu, W. Wallisch, T. M. Huber, J. Bernardi, M. Stöger-Pollach, G. Friedbacher, H. Hutter, B. Yildiz, J. Fleig, *ACS Nano* **2017**, *11*, 11475.

[54] J. P. Sénateur, F. Felten, S. Pignard, F. Weiss, A. Abrutis, V. Bigelyte, A. Teiserskis, Z. Saltyte, B. Vengalis, *J. Alloys Compd.* **1997**, *251*, 288.

[55] R. Rodriguez-Lamas, D. Pla, O. Chaix-Pluchery, B. Meunier, L. Rapenne, H. Roussel, X. Mescot, F. Wilhelm, A. Rogalev, M. Boudard, Q. Rafhay, C. Jiménez, M. Burriel, *Beilstein J. Nanotechnol.* **2019**, *10*, 389.
Carlos Moncasi, Gauthier Lefèvre, Laetitia Rapenne, Hervé Roussel, Ahmad Bsiesy, Carmen Jiménez, Mónica Burriel*

Structural Defects Improve the Memristive Characteristics of Epitaxial La<sub>0.8</sub>Sr<sub>0.2</sub>MnO<sub>3</sub>-based Devices

Asymmetric planar Pt/Ti/LSM\Pt memristive devices show clockwise log *I*(*V*) characteristics with a gradual changes in resistance, and good retention and reproducibility, suitable for neuromorphic computing applications. Compared to defect-free epitaxial LSM/STO devices, the extended defects present in LSM/LAO boost the oxygen drift in the films, improving the LSM/LAO devices memristive properties.

<figure><img src="image_17.png" /></figure>