In Situ Observation of Resistive Switching in an Asymmetric Graphene

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In Situ Observation of Resistive Switching in an Asymmetric Graphene Oxide Bilayer Structure Sungkyu Kim, Hee Joon Jung, Jong Chan Kim, Kyung-Sun Lee, Sung Soo Park, Vinayak P. Dravid, Kai He, and Hu Young Jeong ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b03806 • Publication Date (Web): 09 Jul 2018 Downloaded from http://pubs.acs.org on July 10, 2018

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In Situ Observation of Resistive Switching in an Asymmetric Graphene Oxide Bilayer Structure Sungkyu Kim,† Hee Joon Jung,†, ‡ Jong Chan Kim,∥ Kyung-Sun Lee,# Sung Soo Park,∥ Vinayak P. Dravid,*,†, ‡ Kai He,*,† Hu Young Jeong*,∥,#

†

Department of Materials Science and Engineering and NUANCE Center, Northwestern

University, Evanston, Illinois 60208, United States ‡

International Institute of Nanotechnology, Evanston, Illinois 60208, United States

∥

#

School of Materials Science and Engineering, UNIST, Ulsan 44919, Republic of Korea.

UNIST Central Research Facilities (UCRF), UNIST, Ulsan 44919, Republic of Korea.

*E-mail: [email protected]; [email protected]; [email protected]

ABSTRACT: Graphene oxide decorated with oxygen functional groups is a promising candidate as an active layer in resistive switching devices due to its controllable physicalchemical properties, high flexibility, and transparency. However, the origin of conductive channels and their growth dynamics remain a major challenge. We use in situ transmission electron microscopy techniques to demonstrate that nanoscale graphene oxide sheets bonded with oxygen dynamically change their physical and chemical structures upon an applied electric field. Artificially engineered bilayer reduced graphene oxide films with asymmetric oxygen content exhibit non-volatile write-once-read-many memory behaviours without experiencing the

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bubble destruction due to the efficient migration of oxygen ions. We clearly observe that a conductive graphitic channel with a conical shape evolves from the upper oxygen-rich region to the lower oxygen-poor region. These findings provide fundamental guidance for understanding the oxygen motions of oxygen-containing carbon materials for future carbon-based nanoelectronics.

KEYWORDS: graphene oxide, resistive switching, in situ TEM, bilayer structure, conductive filament Liquid exfoliation of two-dimensional materials has enabled various applications in nanoelectronic fields due to their solution processability and the ease of fabricating thin films and composites.1 Among these, chemically exfoliated graphene oxide (GO) is a promising material because the electrical conductivity of synthesized GO sheets bonded with various oxygen functional groups can be suitably tuned by thermal and chemical reduction processes.2-6 Because of their superior flexibility and transparency, GO-based materials have attracted much attention as insulating layers and conductive layers in wearable electronic devices. In particular, GO-based resistive random access memory (RRAM) is of great interest due to its easy fabrication, controllable film thickness, and superior performance.7-20 Many researchers have attempted to reveal

the conducting path

in resistive switching memory with a

metal/insulator/metal (MIM) structure, because the origin of the electrical path closely influences switching characteristic and device performance. The metal filaments composed by the penetration of electrode metal ions are generally believed to be the origin of the electrical conducting path in resistive switching memory with binary transition metal oxide,21-27 solid electrolytes,28-30 and carbon-based materials.31,32 In addition, direct observations of metal


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nanofilaments that clearly elucidate the resistive switching mechanism were reported using in situ transmission electron microscopy (TEM) techniques.33-48 In GO-based memories, experimental and theoretical work has been performed in an attempt to explain the electrical conduction path in GO films. The formation of metal filaments diffused by electrode metal atoms into the GO layer has been reported to explain the resistive switching mechanism.14-16 Liu et al.8 proposed that the resistive switching behaviours in all-reduced GO memory devices can be attributed to oxygen migration. The metal filament cannot form because highly reduced GO electrodes were used. Recently, Nagreddy et al.49 reported resistive switching driven by redox reactions in the interfacial region between the top electrode and GO films. An amorphous TiOx interface layer was formed by the chemical reaction between the top Ti electrode and the GO films. We also previously reported the resistive switching mechanism in GO-based memories with an interface layer.9-11 Metal electrodes with high oxygen affinity can easily cause the interface layer to react with GO films, and the migration of oxygen ions between the interface layer and GO films is key to inducing the resistive switching. Although these previous studies have shown the critical importance of metal electrodes and oxygen migration in GO films, it is still not clear the movement of oxygen ions inside the GO thin films excluding the effects of metal electrodes when an electric field is applied. Here, we report on the use of in situ TEM to create real-time imaging of the evolution of conductive graphitic channels in Pt/GO/Pt memory devices with noble metal electrode, excluding the effects of the interface layer and the metal cation diffusion from the electrodes. In addition, to clearly demonstrate the directional motion of oxygen ions without bubble destruction, we designed bilayer GO thin films with different amounts of oxygen content using the thermal reduction method. The chemical and structural changes of the GO film with respect


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to temperature are characterized through systematic ex situ and in situ heating studies. The electron diffraction (ED) patterns and electron energy loss spectroscopy (EELS) profiles of the planar stacked GO films reveal the reduction of GO films during the thermal reduction process. In real-time, we visualized the reduction of GO films involving a decrease in the thickness and an increase in the crystallinity during the set process using an in situ electrical biasing system inside TEM. Interestingly, the asymmetrically stacked structure of the pristine GO active layer, which is composed of highly reduced bottom GO films and lightly reduced top GO films, is reversed under the electric field by a drift of oxygen ions, finally leading to conical conductive channels. This work describes the possible movement of oxygen ions bonded to each graphene sheet, which is strongly related to the resistive switching phenomenon in GO-based RRAM devices.

RESULTS AND DISCUSSION Destruction of Pt/rGO/Pt devices by bubble formation. In this study, we use reduced GO (rGO) thin films as the resistive switching layers, but the reduction treatments are subject to different parameters to acquire the desired electrical conductivity. Figure 1a shows a currentvoltage (I-V) curve of a Pt/rGO/Pt device with thermally reduced GO films at 100 °C for 30 minutes. The crossbar-type Pt/rGO/Pt device has a uniform thickness of rGO films sandwiched between two electrodes, as shown in cross-sectional TEM images in of Figure 1a inset and Supporting Information Figure S1. The device was destroyed by bubbles when a bias over -2.5 V (red circle in Fig. 1a) was applied to the top electrode. Figure 1b shows the bubbles that formed along the edge of the top electrode after the negative bias sweep. The magnified SEM image (Figure 1b inset) clearly shows the morphology of bubbles in flower-like shapes, which initially


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formed at the edge of the electrode and propagated outward, indicating that O2 bubbles were formed by the movement and accumulation of oxygen ions bonded to the rGO sheets. A similar gas bubbling behaviour has been previously reported by Yang et al.50 in a Pt/TiO2/Pt device, where oxygen ions chemically bonded to Ti can drift under an electric field and form O2 gas at a certain pressure. We also found that the bubble destructions occurred during both negative and positive sweeps (Supporting Information Figure S2), because the negatively charged oxygen ions tend to escape from the rGO sheets and form O2 gas at the structurally weakest parts regardless of the applied voltage polarity. By viewing from the cross-sectional direction using TEM (Figure 1c), we observed that the partially accumulated O2 gas separated the rGO films (Figure 1c inset); as a result, the layered structure of the Pt/rGO/Pt films was destroyed by the bubbling. This migration of oxygen ions can also be confirmed by the crystallinity change in the rGO films (Figure 1d). The rGO films below the top electrode (region 2, oxygen deficient) show a darker contrast than the other region without the top electrode (region 1, oxygen rich), which is strongly related to the crystallinity of rGO films. The cross-sectional high-resolution TEM (HRTEM) images and the corresponding fast Fourier transform (FFT) patterns clearly reveal the microstructural change in the rGO films, as shown in Figs. 1e and f. The rGO sheets in region 1 exhibit an amorphous-like disordered structure, while those in region 2 represent the densely stacked layered structure, indicating the enhanced crystallinity of rGO films due to the elimination of oxygen ions.

Characterization of modified Pt/brGO/Pt devices. We confirmed that the movement of oxygen ions bonded to an rGO sheet is closely related to its electrical properties and causes a change in the crystallinity. However, a Pt/rGO/Pt device shows bubble destruction during


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electrical sweeps due to the accumulation of oxygen ions. We further modified rGO films using a thermal reduction with various optimized parameters to tune the concentration of oxygen ions contained in the active layer for an improved electrical conductivity.51-53 We synthesized a highly reduced GO (denoted as hrGO) thin films using thermal treatment at 300 °C for 30 minutes and then fabricated the Pt/hrGO/Pt device. Although a Pt/hrGO/Pt device has a lower initial resistance than a Pt/rGO/Pt device without bubble destruction, the resistive switching behaviour was not observed due to the insufficient movement of oxygen ions that hinders the formation of the conducting channels in GO films (Figure 2a). Alternatively, asymmetric bilayer structurebased memory devices can enhance resistive switching performance because the conductive filament is formed in an oxygen-rich active layer with a high resistance value during the set process, while the relatively oxygen-poor layer maintains medium resistance during switching.23,34 We engineered the bilayer-reduced GO (denoted as brGO) devices in form of Pt/brGO/Pt, consisting of a bottom layer of the highly reduced GO films in the same way as hrGO, and a top active layer formed from the lightly reduced GO (denoted as lrGO) films by thermal treatment at 50 °C for 30 minutes (details described in the experimental methods). The Pt/brGO/Pt device shows resistive switching behaviour during the negative bias sweep, as shown in Figure 2a. The distribution of oxygen ions in the brGO film is changed by the applied voltage, and the device shows the write-once-read-many (WORM) characteristics, which means that the low resistance state is continuously maintained during the reverse bias sweep. In addition, Pt/brGo/Pt devices obtained from crossbar array show uniform device-to-device switching characteristics (Supporting Information Figure S3). We conducted an in situ annealing experiment in TEM to verify the effect of the thermal treatment on GO films (Figure 2b). The GO films with randomly stacked sheets were heated up to an elevated temperature (Supporting


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Information Figure S4). We directly observed the variation in the ED patterns in the c-axisoriented GO films. The ED patterns were acquired when the temperature was decreased back to the room temperature after each heat treatment to exclude the effect of the thermally induced lattice expansion of GO films. Thus, we could compare the lattice parameter affected by only the oxygen contents. The structure of GO, close to polycrystallinity with short-range order, was maintained after baking at 50 °C due to the randomly stacked orientations. However, the ringshaped ED patterns evolved to the discrete sharp spots due to the increased crystallinity at temperatures higher than 100°C (Supporting Movie 1). In addition, the lattice spacing of the first bright ring (101ത0) and the second ring (112ത0) were larger than that of native graphite, indicating that the increased lattice spacing is attributed to the chemical bonding of oxygen functional groups (Supporting Information Figure S5).54-56 As the annealing temperature increases, the lattice spacing of GO films approaches the graphite values due to the detachment of oxygen functional groups and the restoration of long-range order crystallinity of the carbon network. To clearly characterize these behaviours, we also examined the electronic state of the C K-edge and the atomic ratio of C to O by tracking the EELS profile of GO films at each temperature during the in situ experiment (Supporting Information Figure S5). The relative ratio of oxygen to carbon decreases during the thermal reduction process, resulting in the enhancement of crystallinity of GO films correlated to the increase in the π* and σ* peaks at the C K-edge. The broad σ* peak becomes sharper and narrower when the crystallinity of GO films changes from an amorphous-like GO to the highly ordered graphitic structure.54,56 The carbonoxygen peak related to sp3 carbon bonding gradually decreases as the temperature increases, indicating that the concentration of oxygen ions bonded to GO sheets is lrGO > rGO > hrGO. This tendency is clearly verified in our intentionally designed brGO thin film by secondary ion


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mass spectroscopy (SIMS), as shown in Figure 2c. The content of O2- in brGO films is nonuniform according to the two regions of the thermally reduced GO films, while the hrGO films show a uniform profile across the entire thickness. We utilized a spherical aberration–corrected TEM with an operation voltage of 80 kV to clearly characterize the microstructural transition in brGO films. Figure 2d shows a crosssectional TEM image of the pristine state. The brGO films have a uniform thickness and contrast without the interface layer that can be formed during the fabrication process. However, the thickness and contrast change after applying a negative bias, as shown in the cross-sectional TEM image of the ON state in Figure 2f. The contrast of the GO films is not only closely related to the crystallinity, but also the atomic resolution HRTEM images clearly represent the crystalline state of GO films. The GO sheets in the cross-sectional HRTEM image in Figure 2g show a more densely and regularly stacked structure than that for the pristine state in Figure 2e, meaning that the crystallinity of the GO films increased after the set process. In addition, the top films (lrGO) show darker contrast and higher crystallinity than the bottom films (hrGO) at ON state. Figures 2h show the FFT patterns obtained from Figs. 2d and f, respectively. The reciprocal frequency (appeared as two spots) represents the interlayer spacing of the electrically reduced GO sheets. While the spot of pristine GO has round and broad divergence, the spot at ON state has a smaller angular divergence and an increased distance from the center. This difference indicates that the distribution of GO sheets became more orientated and the interlayer distance between these sheets were decreased by the movement of oxygen ions. Figure 2i shows the (0002) interlayer distance in GO films. Comparing with a perfect graphite structure, the mechanically and chemically synthesized GO has an increased value due to various oxygen functional groups; but, the interlayer distance decreased from 4.40 Å to 3.62 Å after the set


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process, because the oxygen ions bonded to GO sheets were detached and out-diffused. We believe that this resistive switching behaviour may be related to the movement of oxygen ions between lrGO and hrGO films. The negatively charged oxygen ions initially moved from the top lrGO to the bottom hrGO films when a negative bias was applied to the top electrode. The resistance of the top lrGO films decreased, and the current flowed along the electrical path between the reduced region in lrGO films and hrGO films with additional oxygen ions from the upper layer. However, the locally formed channel in lrGO could not be easily ruptured by oxygen ions detaching from hrGO when a reverse bias was applied to the top electrode, due to the lack of directionality.

In situ observations of resistive switching in a Pt/brGO/Pt device. To observe the electrical switching behaviour of brGO films in real-time, we performed an in situ biasing experiment (details described in experimental methods and Supporting Information Figure S6). Resistive switching was observed at the negative polarity region when the bias was continuously applied to the top electrode of a FIB-cleaved Pt/brGO/Pt device (Figure 3a). The microstructural changes from HRTEM were also simultaneously characterized with I-V curves, as shown in Figures 3b-e (also see Supporting Movie 2). The GO films show a random orientation and low crystallinity at their initial stage. Then the negatively charged oxygen ions diffuse from the top (lrGO layer) to the bottom regions (hrGO layer) under an electric field. The crystallinity of GO films starts to increase from their top region due to the migration of oxygen ions and the recrystallization of the carbon network. The specific area in the bottom regions was reduced in a very short time during the continuous application of bias. The massively reduced lrGO layer and the locally reduced hrGO regions are connected (marked by dashed lines in Figure 3e), and they form an electrical


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path with high conductivity in GO films. This graphitic channel with dark contrast has a wider diameter at the top electrode side. The conical-shaped nanofilaments induced by the movement of oxygen ions or the ordering of oxygen vacancies occur typically and have been reported in various oxide-based RRAMs.33,57-59 The low resistance state did not return to the original state with high resistivity during the reverse bias sweeps, which was the same as that measured in crossbar-type devices. It is notable that the graphitic region in brGO films during the set process expanded to all areas after applying the positive bias. This result can be explained by the diffusion of oxygen ions from oxygen-rich regions (non-graphitic channel regions) to the external of GO films, such as the electrodes or the vacuum environments, rather than to oxygendeficient regions (graphitic channel regions). In other words, the conducting filaments could not be ruptured by the back-diffusion of oxygen ions. We previously suggested the formation of Al nanofilaments at the interface, assisting the local formation and annihilation of conducting filaments by a strong electric field.9,10 But, in this experiment, reversible bipolar resistive switching was not observed due to the failure of local conducting graphitic channels to be annihilated, although asymmetric bilayer GO films which were artificially inserted between the two electrodes played a role as the oxygen supply and reservoir. Finally, we propose the graphitic filament growth mechanism for Pt/brGO/Pt structure, which excludes the interface layer formed by chemical reaction and the metal filaments caused by the metal diffusion from the electrode, as shown in Fig. 4. The lower hrGO films have fewer oxygen ions bonded to GO sheets than do the upper lrGO films by a different thermal reduction process. The negatively charged oxygen ions located in the lrGO films can drift to the lower hrGO films at an early stage of the set process. The interlayer distance of the GO sheets decreases when the oxygen ions detach, and the crystallinity of the GO films increases. An


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oxygen-deficient graphitic region is connected to the lower hrGO films, forming a conductive graphitic channel with a conical shape in GO films.

CONCLUSIONS We successfully demonstrated the dynamic growth of a graphitic channel in modified GO films with an asymmetric bilayer structure using in situ TEM techniques. To prevent the bubble destruction caused by accumulated oxygen gas and enhance resistive switching characteristic, asymmetric GO bilayer films with different oxygen concentrations were implemented, which showed a WORM switching behaviour. The change in the chemical state and electronic structure of the thermally reduced GO films was systematically characterized by the in situ heating experiment. As the annealing temperature increases, the resistance of the GO films decreases due to the annihilation of oxygen ions bonding. Furthermore, ED patterns and EELS spectroscopy confirmed that the crystallinity of carbon network in GO films increases at an elevated temperature and as the concentration of oxygen decreases. During in situ cross-sectional TEM measurements of Pt/brGO/Pt devices, oxygen ions in the upper lrGO films with a relatively high resistance diffused to the lower hrGO films when a negative bias was applied to the top electrode, and the graphitization of the upper lrGO films gradually formed due to the motion of oxygen. Finally, the conical shape graphitic channel was clearly visualized in the GO bilayer structure. From the in situ observation of reversible resistive switching behaviors of an Al/rGO/Pt device (Supporting Information Figure S7 and Supporting Movie 3) with the top interface layer, we also reconfirm that the metallic nanofilament at the interface region, which induces a strong local electric field, plays a critical role in the reversible bipolar resistive switching behaviours of metal/GO/metal systems. We believe that our direct observation of the


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resistive switching in GO films, based on the migration of oxygen ions bonded to GO sheets, may shed light on the understanding and improvement of future carbon-based non-volatile memory systems.

METHODS Fabrication of crossbar-type memories. Graphene oxide was synthesized from natural graphite (Graphi Kropfmühl AG, MGR 25 998 K) using the modified Hummers method. Graphite powder was sonicated for 2 h in water to achieve exfoliated GO sheets. The slurry of GO was centrifuged at 15000 rpm for 30min to remove unexfoliated platelets. Pt/GO/Pt, Pt/lrGO/Pt, and Pt/brGO/Pt devices were fabricated by the following process. Pt bottom electrode with Ti adhesion layer was deposited onto a highly doped Si and SiO2/Si substrates using a shadow mask with 60 µm width by sputter. The GO solution was spread onto the patterned bottom electrode using spin-coating method, and it was thermally annealed at each temperature. Pt top electrode was deposited on the GO films at an angle of 90º. Electrical measurements. A current-voltage (I-V) measurement was performed using a Keithley 4200 SCS semiconductor parameter analyzer. I-V curves of 5x5 memory array were characterized under ambient conditions. During in situ TEM experiment, the electrical properties were characterized by Nanofactory I-V measurement system. Chemical state characterization. XPS measurements were performed on a Sigma Probe (Thermo VG Scientific). The chemical states of carbon and oxygen in the thermally reduced GO thin films were carried out using TOF-SIMS.


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Ex situ TEM. The cross-sectional TEM samples of memory cells were prepared using a focused ion beam (FIB, FEI Helios Nano Lab 450) and were additional thinned by a low-energy Ar-ion milling system (Fischione Model 1040 Nanomill). TEM images were taken by a FEI Titan3 G2 60-300 with an imaging-forming Cs corrector at an accelerating voltage of 80 kV. In situ TEM. The heating specimen was prepared by drop casting GO solution on a lacey carbon TEM grid. The grid was mounted in situ Gatan 628 single tilt heating holder. The temperature was elevated at 30ºC/min. from the room temperature to 300 ºC. The STEM image, ED patterns, and EELS spectra during in situ heating experiment were carried out in the JEOL 2100F TEM (200kV) and the Hitachi HD2300 STEM (200kV). The in situ biasing devices were prepared using a focused ion beam (FIB, FEI Helios Nano Lab 450), and the devices-mounted TEM grid was connected a single-tilt STM-TEM holder (Nanofactory Instrument). The in situ biasing experiment was conducted in a FEI Technai F20 TEM at an accelerating voltage of 200kV.


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Figure 1. I-V behaviours and cross-sectional TEM images of a Pt/rGO/Pt device with bubble destruction. (a) I-V curve plotted on a semi-logarithmic scale and cross-sectional TEM image (inset) in a Pt/rGO/Pt memory device (scale bar: 50 nm). The bias was applied along the arrow. (b) Optical microscopy image showing the bubbles after applying the negative bias. The left inset is a SEM image of the bubbles in the marked region (scale bar: 5um). (c) The crosssectional low magnification TEM images of the bubble in the yellow rectangular region in (b). (d) The cross-sectional TEM image obtained from the edge of the bubble (marked as the red rectangular region in (c)). (d, e) Corresponding cross-sectional HRTEM images obtained from the blue rectangular regions 1 and 2 of (d), respectively. The insets of (e) and (f) are FFT patterns obtained from rGO films.


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Figure 2. Modification and direct observation of characteristics in GO bilayer structure. (a) I-V curves of Pt/brGO/Pt (black) and Pt/hrGO/Pt (blue) devices. The bias was applied along the arrow. (b) In situ observation of ED patterns in GO films obtained at elevated temperatures (scale bar: 2 nm-1). The associated ED patterns are shown in Supplementary Fig. S3. (c) SIMS profiles of hrGO and brGO films. Cross-sectional TEM images at the (d) pristine state and (f) ON state. (e), (g) Corresponding cross-sectional HRTEM images obtained from the rectangular region in (d) and (f), respectively. (h) FFT patterns obtained from (e) and (g), respectively. (i) The intensity profiles of the radial distribution function (RDF) from the center to a desired radius in FFT patterns from (h).


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Figure 3. Real-time observation of a conductive graphitic channel in a Pt/brGO/Pt memory device. (a) I-V curve of a Pt/brGO/Pt memory device during the set state. The bias was applied along the arrow. The inset shows a SEM image of a nanoscale device fabricated by FIB. (b-e) Real-time TEM images capturing the crystallization of GO films related to oxygen movement under an electric field. The arrows indicate the highly reduced GO region. The red solid line indicates the standard for the same region. (f) TEM image after applying a reverse bias (scale bar: 5 nm). The inset shows the I-V curve during the reverse bias sweep.


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Figure 4. Schematic showing the formation process of a conductive graphitic channel in brGO films. (a) Pt/brGO/Pt device in a pristine state. (b) Oxygen ions are diffused under the electric field, and the interlayer distance of GO sheets decreases. (c) The formation of the conductive channel in the brGO films induced by oxygen migration.


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ASSOCIATED CONTENT Supporting Information Supporting figures/movies. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION *E-mail: [email protected](V.P.D). *E-mail: [email protected](K.H). *E-mail: [email protected](H.Y.J). Present Addresses †If an author’s address is different than the one given in the affiliation line, this information may be included here. AUTHOR INFORMATION S.K., V.P.D., K.H., and H.Y.J. designed this work and prepared the manuscript. S.K. performed the device fabrication and electrical measurements. S.K., J.C.K., and H.Y.J. prepared the TEM sample, and S.K., H.J.J., S.S.P., and H.Y.J. performed TEM characterizations. K.L. collected the SIMS data. All authors discussed the results and commented on the manuscript at all stages. ACKNOWLEDGMENT This research was supported by the Creative Materials Discovery Program through the National Research Foundation of Korea (NRF), funded by the Ministry of Science, ICT and Future Planning (NRF-2016M3D1A1900035). This work also made use of the EPIC facility at


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Northwestern University’s NUANCE Center, which has received support from the Soft and Hybrid Nanotechnology Experimental (SHyNE) Resource (NSF ECCS-1542205) and the MRSEC program (NSF DMR-1121262) at the Materials Research Center; the International Institute for Nanotechnology (IIN); the Keck Foundation; and the State of Illinois, through the IIN.

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ACS Nano

(17) Ekiz, O. Ö.; Ürel, M.; Güner, H.; Mizrak, A. K.; Dâna, A. Reversible Electrical Reduction and Oxidation of Graphene Oxide. ACS Nano 2011, 5, 2475-2482. (18) Rani, A.; Velusamy, D. B.; Mota, F. M.; Jang, Y. H.; Kim, R. H.; Park, C.; Kim, D. H. OneStep All-sollution-Based Au-GO Core-Shell Nanosphere Active Layers in Nonvolatile ReRAM Devices. Adv. Funct. Mater. 2017, 27, 1604604. (19) Rani, A.; Kim, D. H. A Mechanistic Study on Graphene-Based Nonvolatile ReRAM Devices. J. Mater. Chem. C. 2016, 4, 11007-11031. (20) Rani, A.; Velusamy, D. B.; Kim, R. H.; Chung, K.; Mota, F. M.; Park, C.; Kim, D. H. NonVolatile ReRAM Devices Based on Self-Assembled Multilayers of Modified Graphene Oxide 2D Nanosheets. Small 2016, 12, 6167-6174. (21) Waser, R.; Dittmann, R.; Staikov, G.; Szot, K. Redox-Based Resistive Switching Memories – Nanoionic Mechanisms, Prospects, and Challenges. Adv. Mater. 2009, 21, 2632-2663. (22) Lee, M.-J.; Han, S.; Jeon, S. H.; Park, B. H.; Kang, B. S.; Ahn, S.-E.; Kim, K. H.; Lee, C. B.; Kim, C. J.; Yoo, I.-K.; Seo, D. H.; Li, X.-S.; Park, J.-B.; Lee, J.-H.; Park, Y. Electrical Manipulation of Nanofilaments in Transition-Metal Oxides for Resistance-Based Memory. Nano Lett. 2009, 9, 1476-1481. (23) Lee, M.-J.; Lee, C. B.; Lee, D.; Lee, S. R.; Chang, M.; Hur, J. H.; Kim, Y.-B.; Kim, C.-J.; Seo, D. H.; Seo, S.; Chung, U. I.; Yoo, I.-K.; Kim, K. A Fast, High-Endurance and Scalable NonVolatile Memory Device Made From Asymmetric Ta2O5−x/TaO2−x Bilayer Structures. Nat. Mater. 2011, 10, 625-630. (24) Yang, Y. C.; Pan, F.; Liu, Q.; Liu, M.; Zeng, F. Fully Room-Temperature-Fabricated Nonvolatile Resistive Memory for Ultrafast and High-Density Memory Application. Nano Lett. 2009, 9, 1636-1643.


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(25) Moors, M.; Adepalli, K. K.; Lu, Q.; Wedig, A.; Bäumer, C.; Skaja, K.; Arndt, B.; Tuller, H. L.; Dittmann, R.; Waser, R.; Yildiz, B.; Valov, I. Resistive Switching Mechanisms on TaOx and SrRuO3 Thin-Film Surfaces Probed by Scanning Tunneling Microscopy. ACS Nano 2016, 10, 1481-1492. (26) You, B. K.; Park, W. I.; Kim, J. M.; Park, K.-I.; Seo, H. K.; Lee, J. Y.; Jung, Y. S.; Lee, K. J. Reliable Control of Filament Formation in Resistive Memories by Self-Assembled Nanoinsulators Derived from a Block Copolymer. ACS Nano 2014, 8, 9492-9502. (27) You, B. K.; Kim, J. M.; Joe, D. J.; Yang, K.; Shin, Y.; Jung, Y. S.; Lee, K. J. Reliable Memristive Switching Memory Devices Enabled by Densely Packed Silver Nanocone Arrays as Electric-Field Concentrators. ACS Nano 2016, 10, 9478-9488. (28) Onofrio, N.; Guzman, D.; Strachan, A. Atomic Origin of Ultrafast Resistance Switching in Nanoscale Electrometallization Cells. Nat. Mater. 2015, 14, 440-446. (29) Raeis Hosseini, N.; Lee, J.-S. Resistive Switching Memory Based on Bioinspired Natural Solid Polymer Electrolytes. ACS Nano 2015, 9, 419-426. (30) Valov, I.; Waser, R. Comment on Real-Time Observation on Dynamic Growth/Dissolution of Conductive Filaments in Oxide-Electrolyte-Based ReRAM. Adv. Mater. 2013, 25, 162-164. (31) Rueckes, T.; Kim, K.; Joselevich, E.; Tseng, G. Y.; Cheung, C.-L.; Lieber, C. M. Carbon Nanotube-Based Nonvolatile Random Access Memory for Molecular Computing. Science 2000, 289, 94-97. (32) Zhuge, F.; Dai, W.; He, C. L.; Wang, A. Y.; Liu, Y. W.; Li, M.; Wu, Y. H.; Cui, P.; Li, R.-W. Nonvolatile Resistive Switching Memory Based on Amorphous Carbon. Appl. Phys. Lett. 2010, 96, 163505. (33) Kwon, D.-H.; Kim, K. M.; Jang, J. H.; Jeon, J. M.; Lee, M. H.; Kim, G. H.; Li, X.-S.; Park,


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1065-1072. (41) Yao, J.; Zhong, L.; Natelson, D.; Tour, J. M. In Situ Imaging of the Conducting Filament in a Silicon Oxide Resistive Switch. Sci. Rep. 2012, 2, 242. (42) Choi, S.-J.; Park, G.-S.; Kim, K.-H.; Cho, S.; Yang, W.-Y.; Li, X.-S.; Moon, J.-H.; Lee, K.J.; Kim, K. In Situ Observation of Voltage-Induced Multilevel Resistive Switching in Solid Electrolyte Memory. Adv. Mater. 2011, 23, 3272-3277. (43) Baek, K.; Park, S.; Park, J.; Kim, Y.-M.; Hwang, H.; Oh, S. H. In Situ TEM Observation on the Interface-Type Resistive Switching by Electrochemical Redox Reactions at a TiN/PCMO Interface. Nanoscale 2017, 9, 582-593. (44) Hubbard, W. A.; Kerelsky, A.; Jasmin, G.; White, E. R.; Lodico, J.; Mecklenburg, M.; Regan, B. C. Nanofilament Formation and Regeneration During Cu/Al2O3 Resistive Memory Switching. Nano Lett. 2015, 15, 3983-3987. (45) Yang, Y.; Gao, P.; Gaba, S.; Chang, T.; Pan, X.; Lu, W. Observation of Conducting Filament Growth in Nanoscale Resistive Memories. Nat. Commun. 2012, 3, 732. (46) Song, S. J.; Seok, J. Y.; Yoon, J. H.; Kim, K. M.; Kim, G. H.; Lee, M. H.; Hwang, C. S. Real-Time Identification of the Evolution of Conducting Nano-Filaments in TiO2 Thin Film ReRAM. Sci. Rep. 2013, 3, 3443. (47) Xu, Z.; Bando, Y.; Wang, W.; Bai, X.; Golberg, D. Real-Time In Situ HRTEM-Resolved Resistance Switching of Ag2S Nanoscale Ionic Conductor. ACS Nano 2010, 4, 2515-2522. (48) Peng, C.-N.; Wang, C.-W.; Chan, T.-C.; Chang, W.-Y.; Wang, Y.-C.; Tsai, H.-W.; Wu, W.W.; Chen, L.-J.; Chueh, Y.-L. Resistive Switching of Au/ZnO/Au Resistive Memory: An In Situ Observation of Conductive Bridge Formation. Nanoscale Res. Lett. 2012, 7, 559. (49) Nagareddy, V. K.; Barnes, M. D.; Zipoli, F.; Lai, K. T.; Alexeev, A. M.; Craciun, M. F.;


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Wright, C. D. Multilevel Ultrafast Flexible Nanoscale Nonvolatile Hybrid Graphene Oxide– Titanium Oxide Memories. ACS Nano 2017, 11, 3010-3021. (50) Yang, J. J.; Miao, F.; Pickett, M. D.; Ohlberg, D. A.; Stewart, D. R.; Lau, C. N.; Williams, R. S. The Mechanism of Electroforming of Metal Oxide Memristive Switches. Nanotechnology 2009, 20, 215201. (51) McAllister, M. J.; Li, J.-L.; Adamson, D. H.; Schniepp, H. C.; Abdala, A. A.; Liu, J.; Herrera-Alonso, M.; Milius, D. L.; Car, R.; Prud'homme, R. K.; Aksay, I. A. Single Sheet Functionalized Graphene by Oxidation and Thermal Expansion of Graphite. Chem. Mater. 2007, 19, 4396-4404. (52) Wei, Z.; Wang, D.; Kim, S.; Kim, S.-Y.; Hu, Y.; Yakes, M. K.; Laracuente, A. R.; Dai, Z.; Marder, S. R.; Berger, C.; King, W. P.; de Heer, W. A.; Sheehan, P. E.; Riedo, E. Nanoscale Tunable Reduction of Graphene Oxide for Graphene Electronics. Science 2010, 328, 1373-1376. (53) Jung, I.; Dikin, D. A.; Piner, R. D.; Ruoff, R. S. Tunable Electrical Conductivity of Individual Graphene Oxide Sheets Reduced at “Low” Temperatures. Nano Lett. 2008, 8, 42834287. (54) Xu, Z.; Bando, Y.; Liu, L.; Wang, W.; Bai, X.; Golberg, D. Electrical Conductivity, Chemistry, and Bonding Alternations under Graphene Oxide to Graphene Transition As Revealed by In Situ TEM. ACS Nano 2011, 5, 4401-4406. (55) Li, J.-L.; Kudin, K. N.; McAllister, M. J.; Prud’homme, R. K.; Aksay, I. A.; Car, R. OxygenDriven Unzipping of Graphitic Materials. Phys. Rev. Lett. 2006, 96, 176101. (56) Mkhoyan, K. A.; Contryman, A. W.; Silcox, J.; Stewart, D. A.; Eda, G.; Mattevi, C.; Miller, S.; Chhowalla, M. Atomic and Electronic Structure of Graphene-Oxide. Nano Lett. 2009, 9, 1058-1063.


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(57) Kim, K. M.; Hwang, C. S. The Conical Shape Filament Growth Model in Unipolar Resistance Switching of TiO2 Thin Film. Appl. Phys. Lett. 2009, 94, 122109. (58) Xue, W.; Liu, G.; Zhong, W.; Dai, Y.; Shang, J.; Liu, Y.; Yang, H.; Yi, X.; Tan, H.; Pan, L.; Gao, S.; Ding, J.; Xu, X.-H.; Li, R.-W. A 1D Vanadium Dioxide Nanochannel Constructed via Electric-Field-Induced Ion Transport and its Superior Metal-Insulator Transition. Adv. Mater. 2017, 29, 1702162. (59) Gao, S.; Liu, G.; Chen, Q.; Xue, W.; Yang, H.; Shang, J.; Chen, B.; Zeng, F.; Song, C.; Pan, F.; Li, R.-W. Improving Unipolar Resistive Switching Uniformity with Cone-Shaped Conducting Filaments and Its Logic-In-Memory Application. ACS Appl. Mater. Interfaces 2018, 10, 64536462.

Table of Contents (TOC) We report a direct visualization of graphitic filament formed by the migration of oxygen ions inside an asymmetric bilayer graphene oxide structure which effectively prevents bubble destruction during the bias sweep.


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In Situ Observation of Resistive Switching in an Asymmetric Graphene

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