Optimization of the multi-mem response of topotactic redox La1/2Sr1/2Mn1/2Co1/2O3- x

APL Mater. 10, 011111 (2022); https://doi.org/10.1063/5.0073490 10, 011111 © 2022 Author(s). Optimization of the multi-mem response of topotactic redox La1/2Sr1/2Mn1/2Co1/2O3−x Cite as: APL Mater. 10, 011111 (2022); https://doi.org/10.1063/5.0073490 Submitted: 30 September 2021 • Accepted: 27 December 2021 • Published Online: 13 January 2022 W. Román Acevedo, M. H. Aguirre, C. Ferreyra, et al. COLLECTIONS Paper published as part of the special topic on Materials Challenges for Nonvolatile Memory ARTICLES YOU MAY BE INTERESTED IN Recent developments and the future perspectives in magnetoelectric nanocomposites for memory applications APL Materials 10, 010901 (2022); https://doi.org/10.1063/5.0076106 Tailoring the synaptic properties of a-IGZO memristors for artificial deep neural networks APL Materials 10, 011113 (2022); https://doi.org/10.1063/5.0073056 Thickness dependence of spin–orbit torques in Pt/Co structures on epitaxial substrates APL Materials 10, 011105 (2022); https://doi.org/10.1063/5.0077074 APL Materials ARTICLE scitation.org/journal/apm Optimization of the multi-mem response of topotactic redox La1/2Sr1/2Mn1/2Co1/2O3−x Cite as: APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 Submitted: 30 September 2021 •Accepted: 27 December 2021 • Published Online: 13 January 2022 W. Román Acevedo,1,2M. H. Aguirre,3,4C. Ferreyra,1,2M. J. Sánchez,1,5M. Rengifo,1,2 C. A. M. van den Bosch,6A. Aguadero,6B. Noheda,7,8and D. Rubi1,2,a) AFFILIATIONS 1Instituto de Nanociencia y Nanotecnología (INN), CONICET-CNEA, San Martín, Argentina 2Centro Atómico Constituyentes, Av. Gral Paz 1499 (1650), San Martín, Buenos Aires, Argentina 3Instituto de Nanociencia y Materiales de Aragón (INMA-CSIC) and Dpto. de Física de la Materia Condensada, Universidad de Zaragoza, Zaragoza, Spain 4Laboratorio de Microscopíías Avanzadas, Edificio I+D, Campus Rio Ebro, C/Mariano Esquillor s/n, 50018 Zaragoza, Spain 5Centro Atómico Bariloche and Instituto Balseiro (Universidad Nacional de Cuyo), 8400 San Carlos de Bariloche, Río Negro, Argentina 6Department of Materials, Imperial College London, London SW7 2AZ, United Kingdom 7CogniGron–Groningen Cognitive Systems and Materials Center, University of Groningen (RuG), Nijenborgh 4, 9747AG Groningen, The Netherlands 8Zernike Institute for Advanced Materials, University of Groningen, Nijenborgh 4, 9747AG Groningen, The Netherlands Note: This paper is part of the Special Topic on Materials Challenges for Non-Volatile Memory. a)Author to whom correspondence should be addressed: [email protected] ABSTRACT Memristive systems emerge as strong candidates for the implementation of resistive random access memories and neuromorphic computing devices, as they can mimic the electrical analog behavior or biological synapses. In addition, complementary functionalities, such as memcapacitance, could significantly improve the performance of bio-inspired devices in key issues, such as energy consumption. However, the physics of mem systems is not fully understood so far, hampering their large-scale implementation in devices. Perovskites that undergo topotactic transitions and redox reactions show improved performance as mem systems, compared to standard perovskites. In this paper, we analyze different strategies to optimize the multi-mem behavior (memristive and memcapacitive) of topotactic redox La1/2Sr1/2Mn1/2Co1/2O3−x(LSMCO) films grown on Nb:SrTiO3. We explored devices with different crystallinities (from amorphous to epitaxial LSMCO), out-of-plane orientation [(001) and (110)], and stimulated either with voltage or current pulses. We found that an optimum memoryresponseisfoundforepitaxial(110)LSMCOstimulatedwithcurrentpulses.Undertheseconditions,thesystemefficientlyexchanges oxygen with the environment minimizing, at the same time, self-heating effects that trigger nanostructural and chemical changes that could affect the device integrity and performance. Our work contributes to pave the way for the integration of multi-mem topotactic redox oxidebased interfaces in multiple device architectures, in order to exploit their memristive and memcapacitive properties for data storage or neuromorphic computation. ©2022 Author(s). All article content, except where otherwise noted, is licensed under a Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/). https://doi.org/10.1063/5.0073490 I. INTRODUCTION Memristors are defined typically as metal–insulator–metal structures able to switch their resistance between different nonvolatile states1,2 and are intensively investigated, nowadays, due to their potential for the development of a new generation of nonvolatile electronic memories, coined as Resistive Random Access Memories (RRAM). In addition, many memristive systems display analog multilevel states and are, therefore, suitable to be implemented in novel neuromorphic computing devices,3as they APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-1 © Author(s) 2022 APL Materials ARTICLE scitation.org/journal/apm can mimic the electrical behavior of biological synapses.4Different neuromorphic capabilities, such as long-term synaptic potentiation/depression3or spike-time-dependant plasticity,5have been reported for memristive devices. In addition, memristor arrays with crossbar architecture were experimentally implemented in simple neural networks (perceptrons) for character recognition,6highlighting the potential of these kinds of devices for the development of bio-inspired electronics. Memristor crossbars were also shown to be suitable as in-memory computing devices,7,8 which could solve the central processing unit (CPU)-memory data bottleneck present in standard computers with von Neumann architecture. Memristive mechanisms rely on the presence and electromigration of charged defects such as oxygen vacancies (OVs),1 ubiquitously found in transition metal oxides.9The presence of these defects strongly affects the local resistivity of the material. Memristive mechanisms in oxides include the formation and disruption of OV conducting nanofilaments,10 the modulation by OVs of interfacial metal–insulator energy barriers,1or a redox reaction taking place when a reactive electrode, such as Al or Ti, is deposited on top of the insulator.11 Complex oxides with a perovskite structure are multifunctional materials displaying in many cases the memristive behavior based on OV electromigration.12 In the case of the celebrated hole-doped manganites, OVs disrupt double exchange Mn–O–Mn bonds and locally increase their resistivity.13 In this way, an oxidized (reduced) manganite displays a more metallic (insulating) electrical behavior. Standard memristive perovskites—including manganites—can stand slight amounts of OVs,leadingtodifferentreducedstatesthatareenergeticallyequivalent. Upon consecutive electrical cycling, any of these reduced states can be stabilized, and therefore, the corresponding resistance state presents a significant dispersion, affecting the device reliability. A more stable memristive response was shown for perovskites such as SrCoO3,14–17 SrFeO3,18–20 or La2/3Sr1/3MnO3,21 which are able to reversibly switch to/from a brownmillerite-like (strongly reduced) phase and are examples of the so-called topotactic redox materials. For these materials, it is possible to reversibly switch between two phases with a different structure and a large difference in oxygen content (and resistivity). As both oxidized and reduced states are linked to an energy landscape with well-defined minima, the electrical switching between both the phases is more reproducible and controlled.14,17 An important issue in topotactic redox-based memristors is the significant oxygen exchange between the device and the atmosphere—initiallyneglected in theproposed scenariostoexplain the memristive mechanisms in perovskites12,13 and more recently incorporated as an important possible mechanism22–24—related to thelargedifferenceinoxygencontentbetweenoxidized and reduced phases. We note that different reports highlight the role of moisture as a source of oxygen—and eventually protons—both for cationic25 and anionic22,24 resistive switches. In order to avoid the device structural damage upon oxygen release in topotactic redox systems and maximize the device reliability, different strategies, such as tuning the out-of-plane orientation of the oxide layer14,17–20 or including an additional material to behave as a by-design oxygen migration channel,26 can be implemented. In the first case, (110) or (111) orientations favor the migration of oxygen into and out of the device through planes parallel to the [001] brownmillerite axis, which displays easier anionic mobilityand facilitatestransport betweentop andbottom electrodes for those out-of-plane orientations. In the second case, the fabrication of self-assembled perovskite-Sm:CeO2composites allows an easy oxygen drift through Sm:CeO2vertical columns. Another suitable option to ease the oxidation and reduction of the memristive material is the use of a scavenging layer in contact with the active oxide.27,28 A very interesting topotactic redox perovskite is La1/2Sr1/2Mn1/2Co1/2O3−x(LSMCO), which displays an oxidized—more conducting—phase with x=0 and a rhombohedral R¯ 3c structure, and a reduced—more resistive—phase with x=0.62 and an orthorhombic Pbnm structure.29 The corresponding redox reaction is reported as La1/2Sr1/2Mn1/2Co1/2O3 ↔La1/2Sr1/2Mn1/2Co1/2O2.38 +0.31O2(g).Fromx-ray spectroscopy, it was inferred that the reduction of LSMCO triggers a change in the average oxidation state of the transition-metal cations, from +3.6 to +2.5,29 suggesting that the created OVs are mostly double charged. We have shown30 that epitaxial Nb:SrTiO3/LSMCO structures display a robust memristive behavior concomitant with a memcapacitive effect, reversible change in the capacitance between different non-volatile states.31–39 The memcapacitance found in this system, CHIGH/CLOW ≥100, was the highest reported to date by a factor of ≈10 and originates at the NSTO/LSMCO interface, where a switchable p–ndiode is formed.30 Memcapacitance presents a high potential for the development of neuromorphic applications; for instance, it has been proposed that a memcapacitor-based neural network for image recognition can perform in a similar way to the one based on memristors, but with a power consumption about 1000 times lower.40 Moreover, it has been reported that the functionality of the network increases with the CHIGH/CLOW ratio for memristive associate capacitive networks,41 designed for image recognition, evidencing the interest of developing and optimizing memcapacitive systems with a large response. Inourpreviousworkonmulti-memLSMCO,wefoundthatthe electroforming process is accompanied by a strong oxygen release and thermal effects that affect the nanostructure of the device and decouples the active zone from the rest of the device, hampering the possibility of implementing these systems in multiple device architectures.30 In this paper, we develop different alternatives to reduce the damage and preserve the device structural integrity, chemistry, andmulti-membehavior.Theseresultscouldeventuallybeextended to other topotactic systems and facilitate the integration of multiple multi-mem systems in more complex architectures, as it is required for memory or neuromorphic computing hardware. II. EXPERIMENTAL We have grown oxidized LSCMO thin films by pulsed laser deposition by using either excimer or YAG lasers. In the latter case, the deposition was assisted with high energy electron diffraction (RHEED). The films were deposited on conducting Nb:SrTiO3 (0.5 wt.%, NSTO) substrates with out-of-plane (001) and (110) orientations. The deposition temperatures ranged between RT and 850○Cinordertoobtain films with differentcrystallinities.Theoxygen pressure was in the range 0.040–0.085 mbar, while the fluence was set between 0.5 and 1 J/cm2. X-ray diffraction was measured with a Panalytical Empyrean diffractometer. Top Pt electrodes were deposited by either e-beam evaporation or sputtering and shaped APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-2 © Author(s) 2022 APL Materials ARTICLE scitation.org/journal/apm in circles with diameters between 45 and 500 μm by standard optical lithography. DC electrical characterization was performed with a Keithley 2612 source meter hooked to a probe station. AC electrical measurements were performed with a LCR BK894 impedance analyzer. This included impedance spectra, recorded for frequencies between 100 Hz and 500 kHz, and capacitance–voltage curves, recorded at 10 kHz. For the latter, the impedance analyzer was set to measure on a parallel RC circuit. In order to minimize spurious contributions in the determination of the capacitances—which could be relevant for measurements in the pF range—we have used short fixed external cables (≈0.5 m) and performed both opencircuit and short-circuit compensation before each measurement. For the latter, the shortcut of the circuit was done on the Pt top electrode of the device-under-test. To obtain an estimation of the error involved in the measurement of capacitances in the pF range, we have performed control experiments with a commercial ceramic 5.6 pF capacitor (tolerance ±5%), obtaining an error of ≈+15% in relation to the nominal value. High resolution Scanning Transmission Electron Microscopy with a High Angular Annular Dark Field Detector (STEM-HAADF) was performed using a FEI Titan G2 microscope with a probe corrector (60–300 keV). In situ chemical analysis was performed by Energy Dispersive Spectroscopy (EDS) and Electron Energy Loss Spectroscopy (EELS). Samples for TEM werepreparedby FocusedIon Beam (FIB)in aHelios650 dualbeam equipment. III. AS-GROWN SAMPLE CHARACTERIZATION In the first place, we describe the structural properties of the as-grown epitaxial LSMCO samples deposited at 800–850○C on (001) and (110) NSTO. Figures 1(a) and 1(b) display x-ray Bragg–Brentano patterns corresponding to (100) and (110) films, respectively. It is found that the LSMCO (00h) and (hh0) reflections are present next to the ones corresponding to the substrate. No indication of parasitic phases is found. The epitaxial character of the films is confirmed by in situ RHEED experiments. For instance, the inset of Fig. 1(a) shows a RHEED pattern corresponding to an as-grown (001) LSMCO film, displaying a stripy pattern FIG.1. (a)X-raydiffractionpattern correspondingto anepitaxialLSMCO thinfilm grownon(001) NSTO.Theinsetshows theRHEED patternmeasuredin situafter finishing the growth process. (b) X-ray diffraction pattern corresponding to an epitaxial LSMCO thin film grown on (110) NSTO. The inset shows an x-ray reflectivity measurement (black symbols) and fitting (red line) performed on the same sample. (c) STEM-HAADF cross section corresponding to an epitaxial (001) NSTO/LSMCO film. The inset showsazoomed-inimageoftheNSTO/LSMCOinterfacewherecationshavebeenlabeled.(d)EELSmapscorresponding toTi,O,Mn,Co,andLaelements.Homogeneous cation distributions are seen. APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-3 © Author(s) 2022 APL Materials ARTICLE scitation.org/journal/apm FIG. 2. (a) STEM-HAADF cross section corresponding to an amorphous LSMCO thinfilmon(001)NSTO.[(b)and(c)]Fast Fourier transforms corresponding to the selected areas of LSMCO and NSTO, displayed in (a), respectively. The amorphous nature of LSMCO is evidenced in the diffuse FFT without diffraction poles in comparison with the diffraction poles of the crystalline NSTO substrate. The zone axis is [100]. (d) Low magnification STEM-HAADF cross section of a NSTO/amorphousLSMCO/Ptdevice.(e) EDS line scans for Sr, Ti, Mn, La, Co, and O elements. The scan, indicated in (d) with a white line, starts at the LSMCO/Pt interface and ends in the NSTO substrate. typically found on epitaxial films with low roughness (root-meansquare roughness below 1 nm). This flat morphology is confirmed by atomic force microscopy (AFM) experiments, not shown here. The inset of Fig. 1(b) shows an x-ray reflectivity experiment performed on an LSMCO (110) sample, where it can be observed that the presence of non-damped oscillations that, after modeling the spectrum with GEN X software, indicates an LSMCO thickness of 16.3 nm and a surface roughness of 0.4 nm, in agreement with the morphologies inferred from RHEED and AFM analysis. Figure 1(c) displays a high resolution STEM-HAADF cross section corresponding to an LSMCO (001) film, evidencing the high structural quality of the film, with a well-defined and coherent NSTO/LSMCO interface. Figure 1(d) shows EELS chemical maps of the structure, which together with EDS quantifications (not shown here) indicate that the stoichiometry of the LSMCO layer corresponds to the oxidized perovskite phase. Next, we describe the structure and chemistry of LSMCO films grown at low temperatures. Figure 2(a) shows a STEM-HAADF cross section corresponding to a NSTO/LSMCO film grown at 200○C. As evidenced by the Fast Fourier Transform [FFT, displayed in Figs. 2(b) and 2(c)] performed on the squared zones displayed in Fig. 2(a), corresponding to the LSCMO film and the NSTO substrate, respectively, LSMCO grows amorphously on singlecrystalline NSTO. EDS line scans shown in Fig. 2(e), going from the LSMCO/Pt interface to the NSTO substrate, as shown in Fig. 2(d), indicate that the stoichiometry of the amorphous LSMCO layer is, within the error of the technique, consistent with the one corresponding to the oxidized LSMCO phase. IV. CHARACTERIZATION OF VOLTAGE-CONTROLLED DEVICES We start by reviewing the electrical response of a 16.5 nm thick epitaxial LSMCO (001) film-based device, with an area of ≈38 ×103μm2and stimulated with voltage pulses. We refer to this sample as LSMCO1. The NSTO substrate was grounded, and the electrical stimuli were applied to the top Pt electrode. The device virgin resistance was ≈1 MΩ, and a forming process, consisting on a sequence of −7 V pulses that reduces the device resistance to ≈5×102Ω, was necessary to initialize the device. After forming, we APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-4 © Author(s) 2022 APL Materials ARTICLE scitation.org/journal/apm simultaneously recorded dynamic I–V curves and hysteresis switching loops (HSLs) that track the evolution of the device remanent resistance. For the dynamic I–V, we applied a sequence of 1 ms voltage write pulses of different amplitudes (from −VMIN to +VMAX) while measuring the current during the application of these pulses. TorecordtheHSL,weappliedasmallreadingvoltage(100mV)after each writing pulse to determine the remanent resistance state. Figures 3(a) and 3(b) display, respectively, the dynamic I–V curve and HSL corresponding to the LSMCO1 sample. Ten consecutive cycles are shown. It is seen that the dynamic I–V curves show rectifying behavior associated with the formation of a n–pdiode at the NSTO-LSMCO interface (we recall that NSTO and LSMCO are nand p-type materials, respectively30,42). It is found that the device switchesfromlow(RLOW ≈400 Ω) tohigh(RHIGH ≈3 kΩ) resistance (SET process) with positive stimuli, while the opposite transition (RESET process) is seen with negative stimuli. The HSL displayed inFig. 3(b) showsthat SET and RESET voltages are stable upon consecutive cycling, but both RLOW and RHIGH present some notable dispersion, an issue that will be addressed later. Retention and endurance experiments are shown in the supplementary material (Fig. S1). We have previously shown that the memristive behavior originates at the NSTO/LSMCO interface upon oxidation/reduction of LSMCO.30 We have successfully simulated the experimental HSL by assuming that LSMCO is in contact with an oxygen reservoir that allows the necessary oxygen release and uptake for the topotactic redox behavior, as shown in Ref. 30. This scenario was further confirmed by electrically cycling the device in vacuum, where no SET processisobserved,duetotheimpossibilityofLSCMOtotakeatmosphericoxygen to becomeoxidized (seeFig. S3in the supplementary material). A very interesting property arises from the analysis of complex impedance spectra, as shown in Figs. 3(c) and 3(d) for RLOW FIG. 3. (a) Dynamic I–V curves corresponding to an epitaxial (001) NSTO/LSMCO/Pt device stimulated with voltage. Ten consecutive cycles are shown. (b) Hysteresis switching loops corresponding to the same device and recorded simultaneously with the I–V curves shown in (a). The arrows indicate the evolution of the curves. [(c) and (d)] Impedance spectra measured on RLOW and RHIGH states, respectively. The inset in (c) shows the equivalent circuit used to fit the spectra. The fittings are shown in red lines. (e) Remanent capacitance measured on the same device after the application of write pulses between 0 and 1 V. Two clear capacitive states are seen, CHIGH corresponding to RLOW and CLOW corresponding to RHIGH. APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-5 © Author(s) 2022 APL Materials ARTICLE scitation.org/journal/apm and RHIGH, respectively. The RLOW spectrum was fitted by assuming the AC equivalent circuit shown in the inset of Fig. 3(c). The set of three resistor/capacitor pairs accounts for both interfaces (R1/C1 and R3/C3) and non-interfacial LSMCO (R2/C2). As it is discussed in the supplementary material, the impedance of C2(ZC2=1/ωC2) is significantly higher than R2, so we assumed the branch that containsC2asanopencircuit.Thesamecircuitwas used tofittheRHIGH state. The values of the fitted circuit elements are shown in Table I. It is found that the memristive behavior is mainly driven by changes in R1, which changes between ≈10 and ≈500 Ω, and we associate to the resistance of the NSTO/LSMCO interface. The change in R1 is concomitant with a large change in the interface capacitance C1, whichswitchesbetween≈34 nFforRLOW and≈1.5pFforRHIGH.The existence of a large memcapacitive effect is confirmed by measuring the remanent, overall, device capacitance CREM after the application of voltage write pulses between 0 and 1 V (small enough to avoid triggering a RESET transition). These measurements are shown in Fig. 3(e) and display a CHIGH/CLOW ratio of ≈100, which is an order of magnitude higher than the figures reported so far.31–39 We also note that CREM is a function of the equivalent circuit elements (Table I) and the frequency ω, as described by the Maxwell–Wagner model.43 We assign the R3/C3elements to the LSMCO/Pt interface. We note the absence of memresistance in that interface, which is expected given the low energy interface barrier reported for p-type perovskites in contact with a high work function metal such as Pt (5.6 eV).44 On the other hand, C3displays a much smaller variation in comparison to the one observed for C1, which drives the overall device capacitance change. The described assignment of the circuit elements to both interfaces is further supported by the analysis performed in Ref. 30, where we modeled the I–V curves for the different memory states by assuming the presence of a switchable n–pdiode,located at theLSMCO interface,together with aninactive LSMCO/Pt interface. We relate the observed multi-mem behavior to the oxidation/reduction of LSMCO. An oxidized interface with NSTOresults in a RLOW–CHIGH state,while a reduced interface leads into a RHIGH–CLOW state. The large difference in oxygen content between both the states critically affects the balance between nand pcarriers at the NSTO/LSCMO interface, which controls the n–p junction depletion layer and triggers the multi-mem behavior. In Fig. 4, we analyze the structural and chemical changes produced on the device upon the application of electrical stress. Figure 4(a) shows a scanning electron microscopy top view of a TABLE I. Numeric values of the elements of the equivalent circuits used to fit the impedance spectra for all the samples. Resistances are given either in Ωor kΩ, and capacitances are given in pF or nF. Sample State R1(Ω) C1(F) R2(Ω) R3(Ω) C3(F) LSMCO1 RHIGH 500 1.5p 90 18 6n RLOW 10 34n 17 21 3n LSMCO2 RHIGH 79.9k 0.12n 70 ⋅ ⋅ ⋅ ⋅ ⋅ ⋅ RLOW 27.7k 0.14n 70 ⋅ ⋅ ⋅ ⋅ ⋅ ⋅ LSMCO3 RHIGH 8.6k 8p 55 340 2.3n RLOW 70 0.55n 15 22 2.5n LSMCO4 RHIGH 1.3k 3p 213 60 6.6n RLOW 16 49n 53 23 4.1n FIG. 4. (a) Scanning electron microscopy top view for a formed epitaxial (001) NSTO/LSMCO/Pt device stimulated with voltage. A darker ring around the landing zone of the electrical tip, related to O2release and LSMCO/Pt expelling, is observed. (b) STEM-HAADF cross section corresponding to the central zone of the image shown in (a). The original epitaxial structure recrystallizes upon electroforming in an arrangement of non-coherent perovskite nanograins. (c) EELS maps corresponding to Ti, O, Mn, Co, and La elements, for a formed (001) NSTO/LSMCO/Ptdevice.Changesinthebrightnessofthemapsindicatechemical modifications in the LSMCO layer upon forming (see the text for details). formed device, where it is evidenced the presence of a ring (darker contrast) around the landing zone of the tip used to make electrical contact. This ring is a signature of oxygen release upon LSMCO reduction. In addition, LSMCO and Pt are expelled upon O2release, electrically decoupling the central zone from the rest of the device. This is confirmed by the lack of dependence of the different resistive and capacitive states with the device (virgin) area, as shown in the supplementary material (Fig. S4). Figure 4(b) shows a STEMHAADF cross section of the formed device, located in the central zone of Fig. 4(a). It is seen that the initial epitaxial LSMCO structure changes to an arrangement of non-coherent perovskite nanograins. We have previously shown that the grains in contact with the NSTO substrate correspond to oxidized LSMCO, but some of the grains in contact with the top Pt electrode present oxygen deficiency.30 In addition, other chemical changes are observed for formed LSCMO, as it is displayed in the EELS maps of Fig. 4(c). It is found that the grains in contact with the NSTO substrate retain the LSMCO cation stoichiometry, but the grains in contact with the top Pt electrode present a deficiency of Mn and Co and an excess of La. These structural and chemical changes are produced by uncontrolled power release and Joule heating upon the application of voltage pulses. Later, we will discuss different strategies to minimize these effects. We focus now on the electrical response of NSTO/amorphous LSMCO/Pt (LSCMO thickness was 36.5 nm, with a device area of ≈95 ×103μm2).WewillrefertothissampleasLSMCO2.The sample APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-6 © Author(s) 2022 APL Materials ARTICLE scitation.org/journal/apm virgin resistance was ≈100 MΩ, and for the forming process, we applied +4 V pulses that lowered the device resistance to ≈20 kΩ. Figures 5(a) and 5(b) display the dynamic I–V curves and HSL corresponding to the LSMCO2 sample. The time width of the writing pulses was around 1 ms. 100 consecutive cycles are shown with a stable behavior. Although rectifying behavior is also observed in LSMCO2 [see Fig. 5(a)], the SET (RESET) transitions take place upon the application of a positive (negative) voltage [Fig. 5(b)], contrarytothe LSMCO1sample, whichdisplayedmirrored SET/RESET polarities (see Fig. 3). This suggests a different memristive mechanism for the LSMCO2 device. RHIGH and RLOW states are ≈90 and ≈30 kΩ, respectively, with a good reproducibility upon consecutive cycling. Retention and endurance experiments are shown in the supplementary material (Fig. S2). A completely different behavior with respect to the LSMCO1 sample arises from the impedance spectra, which are shown for the LSMCO2 sample in Fig. 5(c). In this case, both spectra could be fitted by assuming a parallel resistor–capacitor combination (R1, C1) in series with another resistor (R2). The numerical values of the fitted circuit elements are shown in Table I. The striking difference with the case of LSCMO1 is the absence of memcapacitance, reflected in an unchanged value of C1≈0.1 nF between RHIGH and RLOW states. This is confirmed by the evolution of the remanent capacitance CREM vs writing voltage (range 0–1 V), displayed in the inset of Fig. 5(c), where no changes arefoundbetweenboththeresistivestates(CREM ≈2nF).Inorderto shed light on the microscopic origin of the distinct electrical behavior of the LSMCO2 sample, we performed the experiments shown in Fig. 6.Figure 6(a) displays a scanning electron microscopy top view ofthedeviceafterforming.Althoughitisseenthattheareaofthetop electrode in contact with the electrical tip used to apply the voltage was damaged, no indication of strong O2release is observed [compare with the case of the LSMCO1 sample, Fig. 4(a)], suggesting the absenceoftopotacticredoxtransitioninthecaseof LSMCO2 device. Further information can be obtained from Figs. 6(b)–6(d), displaying a high resolution STEM-HAADF cross section of the formed device [Fig. 6(b)] and EDS line scans from the film’s interface with Pt to the NSTO substrate [Figs. 6(c) and 6(d)]. Several features are observed: (i) a layer of about 12 nm in thickness of the oxide in contact with the NSTO substrate becomes crystallized after forming. (ii) A layer of about 24 nm in thickness on top of the crystallized oxide FIG.5.(a)DynamicI–VcurvescorrespondingtoanamorphousNSTO/LSMCO/Ptdevicestimulatedwithvoltage.100consecutivecyclesareshown.(b)Hysteresisswitching loops corresponding to the same device and recorded simultaneously with the I–V curves shown in (a). The arrows indicate the evolution of the curves. (c) Impedance spectra measured on RLOW and RHIGH states, respectively. The fittings of the spectra are shown as red lines. The inset shows the remanent capacitance measured on the same device after the application of writing pulses between 0 and 1 V. No memcapacitance is found. (d) Simulated hysteresis switching loop with (solid line) and without (dashed line) the inclusion of the non-linear circuit element, emulating the Poole–Frenkel conduction mechanism. The experimental HSL is shown with blue squared symbols. APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-7 © Author(s) 2022 APL Materials ARTICLE scitation.org/journal/apm FIG. 6. (a) Scanning electron microscopy top view of a formed amorphous NSTO/LSMCO/Pt device stimulated with voltage. No evidence of strong O2release is found. (b) High resolution STEM-HAADF cross section corresponding to the same device. The lamella was prepared close to the border of the fused Pt zone shown in (a). [(c) and (d)] EDS line scans corresponding to Sr, Ti, Mn, La, Co, and O elements for the same device. As indicated in (c), the scan starts at the interface between the film and the Pt top electrode and ends in the NSTO substrate. The formation of a bottom crystalline LCMO layer and an amorphous STO top one is observed (see the text for details). layer remains amorphous and displays a different STEM-HAADF Z-contrast, suggesting that it presents a different chemistry in relation to the bottom layer. The latter is confirmed by the EDS line scans displayed in Figs. 6(c) and 6(d), which indicate the presence of an oxide bilayer, consisting of a top layer of SrTiO3(STO) and a bottom layer with a stoichiometry close to the double perovskite La2CoMnO6(LCMO). The presence of LCMO, which does not present a topotactic transition and was reported to increase its conductivitywiththe OV content,45 atthe interface withNSTO explains the absence of memcapacitance for the LSMCO2 device. We note that the stack observed for the formed LSMCO2 device, consisting in a crystalline LCMO/amorphous STO bilayer, resembles the case of memristive crystalline TiO2/a-Si system reported in Ref. 28. Regardingthememristivebehavior,wehavesuccessfullyreproduced in Fig. 5(d) the experimental remanent resistance loop by modeling the OV dynamics between STO and LCMO layers using the Voltage-Enhanced-Oxygen Vacancy (VEOV) model.13,46 Both STO and LCMO behave as n-type semiconductors in which OVs act as electronic dopants reducing their resistivities.45,47 We assume the LCMO/NSTO interface as ohmic and, therefore, not contributing to the memristive effect. On the other hand, the Pt/STO interface is of Schottky-type48 and favors the generation of strong electric fields upon the application of electrical stress, promoting OV electromigration between STO and LCMO. The simulation assumes a 1D chainofLCMOandSTOnanodomains,abletoaccommodatedifferent OV contents. Each nanodomain is characterized by a resistivity thatisrelatedtothelocalOVdensity.Foragivenvalueoftheapplied electricalstimulus,ineachsimulationstep,theOVprofileisupdated through a set of balance equations for the OV transition rates, and the resistance of the sample is computed.13 In the present case, the chain of nanodomains is divided into three regions that define two interfaces: the Schottky Pt/STO interface (LI) and the STO/LCMO interface (RI). Upon the application of positive electrical stimuli to the Pt top electrode, OVs electromigrate from the LI to the RI, with the concomitant reduction of the total resistance. For a negative polarity of the applied stress, the opposite process takes places and the resistance increases. The simulations produce a squared remanent resistance loop, shown by the dashed line in Fig. 5(d), with APL Mater. 10, 011111 (2022); doi: 10.1063/5.0073490 10, 011111-8 © Author(s) 2022