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Transferable Memristive Nanoribbons Comprising SolutionProcessed Strontium Titanate Nanocubes Jiaying Wang,† Satyan Choudhary,‡ William L. Harrigan,§ Alfred J. Crosby,‡ Kevin R. Kittilstved,§ and Stephen S. Nonnenmann*,† †

Department of Mechanical and Industrial Engineering, ‡Polymer Science and Engineering Department, and §Department of Chemistry, University of Massachusetts Amherst, Amherst, Massachusetts 01003, United States S Supporting Information *

ABSTRACT: Memristors, often comprising an insulating metal oxide film between two metal electrodes (MIM), constitute a class of two-terminal devices that possesses tunable variations in resistance based on the applied bias history. Intense research remains focused on the metal− insulator interface, which serves as the crux of coupled electronic−ionic interactions and dictates the underpinning transport mechanisms at either electrode. Top-down, ultrahigh-vacuum (UVH) deposition approaches for MIM nanostructures yield highly crystalline, heteroepitaxial interfaces but limit the number of electrode configurations due to a fixed bottom electrode. Here we report on the convective selfassembly, removal, and transfer of individual nanoribbons comprising solution-processed, single-crystalline strontium titanate (STO) perovskite oxide nanocrystals to arbitrary metallized substrates. Nanoribbon transferability enables changes in transport models ranging from interfacial trap−detrap to electrochemical metallization processes. We also demonstrate the endurance of memristive behavior, including switching ratios up to 104, after nanoribbon redeposition onto poly(ethylene terephthalate) (PET) flexible substrates. The combination of ambient, aerobic prepared nanocrystals and convective self-assembly deposition herein provides a pathway for facile, scalable manufacturing of high quality, functional oxide nanostructures on arbitrary surfaces and topologies. KEYWORDS: resistive switching, interfaces, nanoparticles, strontium titanate, electronic transport

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INTRODUCTION Metal oxide sandwich structures exhibiting resistive switching phenomena demonstrate significant promise as next-generation functional memory storage devices due to simple, highly scalable structures, low power consumption, fast response times, and multistate logic potential.1−4 The low/high resistance states (LRS; HRS) across metal oxide interfaces within the conventional metal−insulator−metal structure, as reduced to true nanoscale dimensions, become greatly influenced by concomitant local electric field effects,5 Schottky barrier modulation,6−8 and defect distribution dominated transport.9 In order to harness unique size-dependent resistance phenomena, other recent studies demonstrated resistive switching character within vertically aligned carbon nanotubes,10 individual perovskite nanotubes,11 binary oxide nanowires,12 and individual TiO2 nanoparticles.13 Studies involving these quasi-zero-dimensional and one-dimensional nanostructures remain rather scarce as device fabrication involves complicated multistep lithographic processes or suffer from the lack of order or periodicity requisite of memory architectures. Here we report on the resistive switching character of individual nanoribbons comprising single-crystalline strontium © XXXX American Chemical Society

titanate nanocubes (STONCs). The convective self-assembly approach, previously utilized for preparing nanoribbons and grids of CdSe quantum dots,14 enables facile fabrication of high-quality perovskite oxide low-dimensional nanostructure arrays onto arbitrary substrates that retain their resistive switching functionality. Moreover, we demonstrate the transfer of the nanoribbons from the original substrate onto a second, arbitrary substrate possessing a different contact surface, posing significant implications for tailoring the resistive response type through substrate selection.

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RESULTS AND DISCUSSION Colloidal SrTiO3 (STO) nanocrystals with oleic acid capping ligands were prepared using a modified hydrothermal method.15 Figure 1 highlights the nanocrystalline morphology of a typical STO nanocube and convective self-assembly processes. Figure 1a shows a high-resolution transmission electron microscopy (HRTEM) image of an individual STONC that displays single-crystalline features and an average Received: January 5, 2017 Accepted: March 9, 2017 Published: March 9, 2017 A

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recently reported value of 2.7 nm for oleic acid-capped BaTiO3 nanocrystal superlattices.22 The FBD process thus yields large arrays of intrinsic STONC nanoribbons linked by oleic acid ligands (Figure 2a; inset) on a Nb:STO substrate, providing a model system to explore the resistive switching properties of nanocrystalline STO (Figure 2a).

Figure 1. (a) An HRTEM image of an individual STO nanocube (STONC), with a nominal edge length of ∼10 nm. (b) An illustration of the stop-and-go flow coating process; inset: an illustration highlighting the long, bell-shaped nature of the STONC nanoribbons. (c) An optical micrograph (scale bar: 500 μm) showing highly scalable arrays of STONC nanoribbons (length = 1.5 cm) with highly regular periodicity. (d) A plot of STONC nanoribbon height and width dimensions versus the stop time of the flow coating process.

cube length of approximately 10 nm. Grain size estimates of the Scherrer analysis of the powder XRD spectrum of the STONCs corroborates these dimension values (r = 15 ± 5 nm; Supporting Information Figure S1). The observation of lattice fringes in the HRTEM and sharp peaks within the diffraction pattern suggest phase pure STO and indicate good crystallinity quality.16 Despite the apparent high crystalline quality, the STONCs display broad emission centered at ∼465 nm resulting from self-trapped excitons17 that are also observed in bulk crystals of STO (Figure S2). Colloidal suspensions of single-crystalline STONCs in toluene served as the precursor component in fabricating highly ordered nanoribbons using a stop-and-go flow coating approach.18 This method, flexible blade deposition (FBD), as illustrated in Figure 1b, utilizes the flow of a colloidal suspension (on a rigid substrate) under a flexible poly(ethylene terephthalate) (PET) blade held at a fixed height, which significantly restricts the geometry that the STONCs assemble on the surface. The FBD approach previously yielded striped and grid patterns of multicolor CdSe quantum dots on silicon wafers and aluminum foil,14,19 gold nanoparticle ribbons on glass substrates,20 and PMMA nanoribbons on silicon.21 This technique traps the colloidal suspension between the blade and substrate by capillary forces, forcing the nanocrystals (STONCs) toward the contact line created by solvent evaporation. Here a rigid 0.7 wt % Nb-doped strontium titanate (Nb:STO) substrate is affixed to a linear translation stage and moved in a “stick-and-slip” motion to enable flow coating deposition of STO nanoribbon arrays on to the substrate, as confirmed by the optical microscopy image in Figure 1c (scale bar: 500 μm). The height and width dimensions of the nanoribbons are controlled by varying the intermittent stopping time (td) of the stage shifts (Figure 1d). Careful selection of td yielded height ranges from 100 to 425 nm and widths between 3 and 23 μm. The STONC ribbons span nearly the entire substrate (length = 1.5 cm). The average interparticle distance of STONC monolayer nanoribbon segments determined by high angle annular dark field TEM was approximately 3.2 ± 0.5 nm (Figure S3), in agreement with

Figure 2. (a) An illustration of the c-AFM measurement across an individual STONC nanoribbon (green) on a Nb (0.7 wt %): STO substrate (teal). Inset shows the oleic ligands between two STONC with a length 3.2 ± 0.5 nm. (b) An I−V plot of a STONC nanoribbon (height = 300 nm) displaying “eightwise” resistive switching character, as indicated by the arrow markers. Here “1” indicates the SET process, “2” and “3” the LRS response, and “4” the RESET process. The solid purple portion indicates the HRS while the purple-outlined region represents the LRS. (c, d) An illustration of the operating switching mechanism and the band structure across the multicube Pt/STONC interface. (c) Applying a negative voltage injects electrons from the trap sites (oxygen vacancies or ligands; magenta spheres) to the top (tip) electrode, resulting in unfilled traps that narrow the Schottky barrier and induces the HRS → LRS transition (b, “1”). (d) A positive applied voltage fills the interfacial traps with electrons (black dots in magenta spheres) and recovers the Schottky-like barrier, yielding a LRS → HRS transition (b, “4”).

Conductive AFM (c-AFM) represents a powerful tool in the study of resistive switching. It locally applies a bias to the substrate to probe the homogeneity of the charge transport properties of the nanoribbons and nanocrystal building blocks using the conductive tip as the translational top electrode. Tedious manipulation and drift stability issues render local atomic force microscopy (AFM) of individual nanocrystals extremely difficult, thus yielding a rather limited amount of studies.5,9,13,23,24 Its recent uses also include the verification of highly conductive filaments after electroforming.25,26 For cAFM, a Pt/Ir-coated conductive tip is placed in direct contact with individual nanoribbons, thus creating a Pt/Ir(tip)/STO/ Nb:STO test structure. A relatively high niobium concentration (0.7 wt %) yields a low resistivity (0.007 Ω·cm), effectively rendering the Nb:STO substrate as a thick bottom electrode similar to those found in studies of Sr2TiO4 thin films.27 Figure 2b displays a typical current−voltage (I−V) response of an individual STONC nanoribbon, as collected using a sweeping B

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ACS Applied Materials & Interfaces voltage (±10 V, 1 Hz) applied to the Nb:STO bottom electrode and across the test structure, where the tip electrode maintained ground. Starting from zero applied bias, a negative sweep shows low current level at high resistance state (HRS) until a set voltage, where the induced low resistance state (LRS) corresponding to an increase in current (“1”). Note that the resistance switch during the set process is gradual, differing from the typical filamentary type (Figure 4c) switching that presents an abrupt change. The cutoff observed at larger applied negative voltages (and subsequent currents (I > −10 μA)) results from built-in transimpedance amplifier limitation (Asylum Research Dual-Gain ORCA Module). The resulting LRS can be maintained even when the applied voltage was removed, confirming the nonvolatile nature of the nanoribbons.1,28 As the applied bias increases toward +10 V, the current drops and a return to the HRS is observed at a reset voltage about 3 V (“4”). The (“1−2−3−4”) set−reset hysteresis observed in Figure 2b represents an “eightwise” (counterclockwise) resistive switching response (opposite to the polarity based on biasing the TE).2−4,26,27,29 This result contrasts the “counter-eightwise” response observed in 10 nm thick epitaxial intrinsic STO thin films.30 Though the nanoribbon height is approximately equivalent to the thickness of the epitaxial film previously reported, the primary difference is the apparent reduction in dislocation density31,32 within the individual STONC component nanocrystals. Extended defects such as dislocations typically enable the migration of oxygen vacancies via the filamentary conduction mechanism that dictates “countereightwise” (clockwise) switching within single crystalline STO.1−4,25−27,29,32−34 Figures 2c and 2d illustrate the possible mechanisms driving these phenomena, which closely resemble those recently determined for tip-induced resistive switching of individual TiO2 nanocrystals.13 In that study, sandwich test structures comprised individual TiO2 nanoparticles sandwiched between a Pt/Ir tip electrode and a Pt/Ir planar bottom substrate electrode. They subsequently applied a negative bias formed a depletion region at the bottom electrode as vacancies migrated toward the top electrode (tip), resulting in a LRS to HRS transition. Here we ruled out possible Schottky-like barrier modulation effects via oxygen vacancy migration due to the opposite polarity dependence observed, the down-scaling of the particle size, and the resulting reduction in vacancy conducting pathways that typically occur along extended defects (dislocations, interfaces) of single crystalline STO. In this study, the moveable top Pt/Ir tip electrode maintains ground while the bias is applied to the bottom Nb:STO electrode. Previous studies suggested a higher volume of oxygen vacancies concentrate along the extended defects due to increased reducibility compared to bulk.32 In this study, we posit vacancies accumulate at the interfaces between individual STONC. Electronic absorption studies showed that the majority of defects populate the surface of the nanocrystal.15 An oxygen-deficient surface therefore creates a space-charge depletion region that effectively hinders oxygen diffusion.2,31,32 This accordingly explains the observed lower current flow at low voltage region (Figure 3b). For these reasons, we assume the positively charged oxygen vacancies remain fixed at the Pt/ Ir(tip)/STO interface and form a trap state for the following switch. The rectifying I−V response of HRS (Figure 2b; solid purple portion between points “1” and “4”) results from the Schottky-like barrier formed along the (Pt/Ir)tip/STO interface.

Figure 3. (a) A topographic AFM image of an individual STONC nanoribbon (height = 400 nm). (b) A semilog I−V plot of the STONC nanoribbon for the 1st (red), 100th (purple), 500th (light blue), and 1000th cycle (green). (c) An endurance plot of log current vs number of cycles between the LRS (blue squares) and HRS (red diamond) measured locally over the left edge (a, ●), top (a, ▲), and right edge (a, ×) at a read voltage of −1.5 V, respectively. The average RON/ROFF ratio exceeds 103.

The initial bias sweep injects electrons into the STONC film and fills the trap sites at the tip−STONC interface. Applying a sufficiently large negative voltage to the bottom electrode releases the electrons, thus yielding unfilled traps and reduces the Schottky-like barrier width. This process enables electron tunneling, hence we observe a HRS → LRS transition (Figure 2c; “1”). Applying a large positive voltage to the bottom electrode refills the traps at the tip−STONC interface with electrons that neutralize the positive charge of the oxygen vacancies (Figure 2d; “4”). This process forms a strongly insulating interface and increases the barrier width, thus inducing the LRS → HRS transition. We note the extent that vacancy accumulation/depletion occurs at the nanoribbon/ Nb:STO interface must vary appreciably compared to heteroepitaxial STO/Nb:STO thin films. Recent studies of distorted STO/Nb:STO heterointerfaces observed rectifying behavior attributed to the modulation of interfacial vacancy depletion,35 where the electrode configuration used resulted in a symmetrical hysteresis, hallmark of two Schottky junctions. The hysteretic response in Figure 2b is identical to the observations for Au/Nb:STO/Ag36 and s-In/Nb:STO/o-In37 nanostacks possessing asymmetric Schottky/Ohmic contacts. Therefore, the local c-AFM I−V response of as-deposited STONC nanoribbons on Nb:STO is governed by conduction mechanisms at the tip−sample interface that enable tunneling via charge trapping/detrapping cycles defined by interfacial oxygen vacancy electron traps3,5,27,34,38−42 or Schottky barrier modulation.6−8,43,44 Figure 3 highlights the results of LRS/HRS switching endurance testing across STONC nanoribbons as collected locally by c-AFM. Figure 3a shows a (20 μm × 20 μm) topographic image of an individual STONC nanoribbon, as prepared by FBD on an Nb:STO substrate. The semilog I−V plot in Figure 3b exhibits relatively robust endurance properties, as demonstrated by the overlap after 1 (red), 100 (purple), 500 (light blue), and 1000 (green) cycles, respectively. The FBD process likely results in incomplete deposition of STONC toward the ribbon edges in multilayer C

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Figure 4. (a) Illustration of the STONC nanoribbon to STONC nanohelix release/transfer cycle. After the FBD process samples are submersed in water, creating a suspension of STONC nanohelices that are redeposited onto arbitrary substrate (Ag-coated PET substrate, lavender-purple shown). (b) An optical micrograph of STONC nanohelices released into water from the Nb:STO substrate (scale bar: 500 μm). Inset: an optical micrograph of an individual STONC nanohelix displaying its flat-ribbon-like geometry (scale bar: 200 μm). (c) The I−V response of an individual STONC nanoribbon post-transfer on a Ag-metallized PET substrate. (d) The STO/Ag/PET LRS branch on a log−log scale, which displays a linear slope of 1 corresponding to the Ohmic conduction model. (e) The endurance of STO/Ag/PET under voltage sweeping for 100 cycles. (f) The I−V response of an individual STONC nanoribbon post-transfer on a Pd-metallized PET substrate, exhibiting “eightwise” switching character. (g) The STO/Pd/ PET LRS branch on a log−log scale, displaying a slope of 1 at low voltages and an increased slope (∼2) at higher voltages, corresponding to the space charge limited current (SCLC) conduction model. (h) The STO/Pd/PET HRS branch of the I−(1/V) response on a log−log scale, which exhibits a negative linear slope indicative of trap-assisted tunneling (TAT) conduction.

favorably with higher quality Pt−STO interfaces created by dc sputtering and e-beam evaporation (102−104)7 or with sol−gel processed thin films with Pt electrodes (103).45 We note that STONC nanoribbons display a more triangular than rectangular geometry (see Supporting Information); hence, the stability displayed in the position-based endurance tests further confirms that a trap/detrap mechanism at the tip−(top) STONC interface defines transport, as no discernible thickness dependence is observed along either edge of the nanoribbon. A key advantage of STONC nanoribbons produced by FBD is the ability to release the nanoribbons into a liquid for redeposition onto a second, arbitrary substrate or surface topology (Figure 4a). Previous studies successfully demonstrated this process for releasing CdSe stripes and grids into an aqueous solution.14,19 Figure 4b shows an optical micrograph of

ribbons (Figure S4), yielding a possible thickness dependence. We therefore also performed local endurance tests at different tip positions along the nanoribbon from left−right at the locations marked by the ●, ▲, and × marks on the topographic image found in Figure 3a, respectively. Endurance profiles of LRS/HRS switching current collected at a read voltage of −1.5 V are shown in Figure 3c, where blue (square marker) lines indicate the LRS and red (diamond marker) lines indicate the HRS. The results shown in Figure 3c demonstrate robust and thickness-independent LRS/HRS values for the nanoribbons. The average RON/ROFF ratio exceeds 103 over a span of 1000 switching cycles with no apparent position dependence. While the ratio does not quite approach the reported value (106) of UHV pulsed laser deposition-prepared 10 nm thick epitaxial thin films,30 the value does compare D

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ACS Applied Materials & Interfaces STONC nanoribbons floating on the water surface upon removal by dissolution of a poly(acrylic acid) sacrificial layer and of nanoribbons forming helical mesostructures when released away from the liquid−air interface (inset). Previous results on CdSe and PMMA nanoribbons have demonstrated that careful selection of nanoribbon dimensions directly influences the equilibrium helical shape and geometry-defined mechanical properties.21,46 To demonstrate the versatile nature of FBD-deposited nanoribbons, we released the structures into deionized water and then transferred and redeposited the ribbons onto a flexible PET substrate with Ag and Pd metal coating for examination via c-AFM. Virtually all studies involving resistive switching of thin films remain limited by the top-down lithographic processes to evaporate or sputter various top electrode materials on a fixed film/metal/substrate bottom electrode configuration.45,47,48 The FBD transfer process provides the unique advantage of selecting the bottom electrode material to induce or potentially enhance a desired transport mechanism within functional oxide nanostructures. The I−V response of the Ir/STO/Ag/PET structure shown in Figure 4c exhibits switching characteristics opposite of the eightwise switching displayed by the same nanoribbons on the Nb:STO substrate in Figure 2b. The I−V response displays highly linear Ohmic behavior under the LRS and sharp, pronounced LRS ↔ HRS transitions, indicative of formation and dissolution processes of filament induced by cation migration with the chemical active electrode being Ag.2,3 This curve fitting of the log−log response confirmed the LRS exhibits Ohmic character, with a slope of 1, which confirms the formation of metallic filaments (Figure 4d). Applying a positive bias to the Ag bottom electrode creates Ag cations (Ag → Ag+ + e−) that drift through the STO nanoribbons toward the top (tip) counter electrode, where they subsequently reduce to Ag atoms and form the highly conductive filaments responsible for the steep current increase. Conversely, an applied negative voltage induces Joule heating and oxidizes the Ag atoms at the counter electrode, which causes the conductive filament to rupture and starts the LRS → HRS transition (Figure 4e). Moreover, the transferred nanoribbons possessed an increased RON/ROFF ratio (104) and a reasonable endurance lifetime of ∼100 cycles (see Figure S6). The slight variation in LRS current during endurance testing is attributed to the randomly branched growth of filaments near the inert Pt/Ir electrode.49 Improving the endurance stability within transferred nanoribbons warrants further studies involving optimization of the release, redeposition, and secondary substrate/contact selection. STONC nanoribbons were also transferred onto Pdmetalized flexible PET substrates to create a Pt/Ir/STO/Pd/ PET test structure, which displayed entirely different switching characteristics, as observed in Figure 4f−h. In this case, the nanoribbons possessed counter-eightwise switching character but showed no rectifying behavior at the HRS. This is due to the relative symmetry of the sample structure. Figure 4g shows the LRS response in a log−log scale, exhibiting an increasing slope from 1 (I ∝ V) to 2 (I ∝ V2) with increasing applied voltage, thus suggesting transport based on the space charge limited current model (SCLC). The HRS response (Figure 4h) displays a linear response that trends with ln I ∝ 1/V, fitting the trap-assisted tunneling model (TAT)42 and emphasizing the critical role that traps play during switching at the Ir-tip/STO interface. To induce an HRS → LRS transition, an applied

negative voltage to the Pd bottom electrode injects electrons into the STONC ribbons, which are then captured by the trap sites along the Ir-tip/STO interface and subsequently emitted to the tip. This two-step tunneling process is likely assisted by oxygen vacancies and the oleic acid ligands between nanocubes. To verify the role of ligand traps during switching, STO nanoribbons were also transferred onto a Pd-metalized, silicon substrate and annealed at 500 °C for 2 min under ambient conditions to remove all ligands. The fitting results are presented in Figure S8. The LRS character remains identical for the flexible Pd/PET substrate and the hard Pd/Si substrate, as both follow a linear relation at low voltages and a quadratic behavior at higher voltages. However, the HRS character for nanoribbons on Pd/Si substrates deviates significantly from the TAT model and more closely resembles the behavior of Fowler−Nordheim tunneling (Figure S8d), which suggests a much weaker trap-assisted effect. Since the annealing step removed the presence of ligands, the change in conductive mechanism confirms that the ligands play a critical role as traps for assisting electron tunneling.50 Note that the ligands exert certain effect on the reduction of current, thus demonstrating their potential role in reducing power consumption. The c-AFM I−V response of transferred STONC nanoribbons on flexible, Pd-metallized PET substrates was also collected under a fixed curvature, as shown in the inset of Figure 5. STO/Pd/PET assemblies were affixed to a stainless

Figure 5. I−V response of an individual STONC nanoribbon on a Pd/ PET substrate under strain. Inset: the as-transferred nanoribbons are affixed to a steel rod of known diameter, inducing a substrate strain of ∼3%. The local I−V response collected via c-AFM retains the identical (repeatable) switching behavior and same curve fitting (see the Supporting Information).

steel nail shaft, inducing a strain of ∼3% (see Supporting Information). Curve fitting of the I−V response in Figure 5 (see Figure S9a,b) confirmed the nanoribbons retained both the SCLC-based LRS and TAT-based HRS character displayed by the as-transferred STO ribbons on the Pd/PET substrate (Figure 4f). The stability exhibited by the nanoribbons in a bending state demonstrates enormous promise as a candidate for flexible electronic device platforms where electrical and mechanical reliability still present significant challenges.51

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CONCLUSION In summary, we introduced a facile, inexpensive, scalable fabrication method that yields memristive nanoribbon structures comprising solution-processed, single-crystalline complex oxide nanoparticles possessing complete transferability to arbitrary substrates. As-fabricated SrTiO3 nanoribbons display stable eightwise switching, in contrast to the countereightwise polarity exhibited by single-crystal STO thin films E

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coating with a frequency of 70 kHz and a spring constant of 2 N/m (ASYELEC-01) was used. Each bottom electrode was attached to the sample holder via silver paint (TED PELLA, Inc.) to form a electrical path between the sample bias and the sample surface. During the measurement, the sample was biased between ±10 V at the bottom electrode using while the tip is held at 0 V potential. Current was collected through the tip and subsequently passed through a dual gain (ORCATM) transimpedance amplifier. Prior to I−V measurements, a topography of the sample surface was taken in AC mode. To do a local I−V plot, the tip was moved to a different position of the ribbon and engaged to it with a set point of 0.02 V. In the initial electroforming, a negative voltage of −10 V was applied to the bottom electrode to switch the system into LRS.

with appreciable dislocation densities. The ability to transfer nanoribbons to arbitrary substrates creates a rich parameter space by which any desired polarity response, improved RON/ ROFF ratio, and switching type is potentially tailored through bottom electrode metallization selection (or change). The transferability and durability of memristive properties posttransfer or in a strained state boast promise as flexible and printed electronic elements. We expect this study will spur work that explores metal (electrode) oxide interface-induced functionality within solution-processed nanostructure constituents in metal−insulator−metal architectures, including barrier modulation in ferroelectric52 and ferromagnetic53 tunnel junctions, screening lengths in multiferroics,54 oxygen electrocatalytic activity,55 and lattice−electronic interactions within oxide heterostructures.56

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b00220. Additional experimental details and analysis, highresolution TEM and HAADF images, I−V curves, powder X-ray diffraction, electronic absorption and emission of colloidal STO NCs, and X-ray diffraction of STO NC powders (PDF)

METHODS

Colloidal STO Nanocrystal Synthesis: Colloidal undoped SrTiO3 nanocrystals were prepared using a modified hydrothermal method without hydrazine as reported previously. Nanocrystals were characterized using absorption/emission, powder X-ray diffraction, and transmission electron microscopy (TEM). The nanocrystals display a broad emission peak centered at ∼465 nm that has is assigned to self-trapped excitonic emission associated with oxygen vacancies.17 The powder X-ray diffraction pattern of these nanocrystals display only the cubic perovskite phase of strontium titanate. Highresolution TEM images (JEM 2200FS-JEOL) of colloidal SrTiO3 NCs deposited on a 3 nm carbon film supported on 400 mesh copper TEM grids (Electron Microscopy Sciences). TEM analysis reveal that the nanocrystals are cuboidal with average edge dimensions of ∼10 ± 2 nm. The nanocrystal concentration was estimated from analysis of the Ti concentration obtained by inductively coupled plasma optical emission spectroscopy (ICP-OES, PerkinElmer 4300 DV) and the average particle size from TEM (see the Supporting Information). STONC Nanoribbon Convective Self-Assembly, Removal, and Transfer: The “stop-and-go” flow coating method described previously14 was used to create STO nanoribbons from colloidal STONCs building blocks. The custom-built instrument uses a poly(ethylene terephthalate) film (82 ± 7 μm in thickness, supplier McMaster-Carr) as a flexible blade, with a “scored” region (2 mm) at the edge. This blade is attached to an adapter with pitch and roll variability, allowing the scored region of the blade to be aligned parallel to a substrate and brought into contact. The substrate is attached to a nanopositioner (Burleigh Inchworm controller 8200), which is programmed to perform a series of stop-and-go steps, with stop time (td), velocity (v), and spacing (d) as user-defined variables. STO colloidal solution (8 μL) is injected between the substrate and scored region of the blade. The capillary forces confine the injected solution to the scored region and meniscus height is dictated by the capillary interactions. STO nanoribbons are flow coated on commercially available Nb-doped STO (0.7 wt % Nb, MTI Corp.) and silicon (P(100) 0−100 ohm·cm; single-side polished; UniversityWafer Inc.) substrates. In order to release the STO nanoribbons from the substrate, poly(acrylic acid) (PAA) is used as a sacrificial layer. An aqueous solution of PAA (20 mg/mL) is spun-coat (3000 rpm, 30 s) on a prewashed silicon wafer. Nanoribbons are released by dissolving the PAA sacrificial layer (height ∼50 nm) with water. Once the nanoribbons were released, transfer was completed by placing the second, arbitrary substrate in direct contact with the surface of the water meniscus. All commercial materials were used as received. The Ag and Pd electrode deposition (thickness = 100 nm) was performed by RF- and DC-sputtering of metal targets (Kurt Lesker, 99.99%) at room temperature and a base pressure of 10−6 Torr. The metallized-PET films and substrates were rinsed with isopropyl alcohol and toluene and dried with a filtered stream of compressed air. Local Conductive AFM (c-AFM) Measurements: The electrical measurement was conducted at room temperature on an ORCAConductive AFM (Cypher ES, Asylum). A silicon probe with Pt/Ir tip

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (S.S.N.). ORCID

Stephen S. Nonnenmann: 0000-0002-5369-9279 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research was partially supported by University of Massachusetts Amherst start-up funding and the UMass Center for Hierarchical Manufacturing (CHM), a NSF Nanoscale Science and Engineering Center (CMMI-1025020). A.J.C. and S.C. gratefully acknowledge financial support from the Army Research Office (W911NF-14-1-0185). The authors also acknowledge use of the facilities at the CHM Conte Nanotechnology Cleanroom for thin film deposition and lithography processes. We also thank Alexander Ribbe for collecting the SEM and HAADF-TEM images of the nanocrystals/nanoribbons at the W. M. Keck Electron Microscopy Facility at UMass. J.W. performed all substrate/ electrode preparation, c-AFM measurements, and transport analysis and cowrote the manuscript. S.C. performed all nanoribbon fabrication and transfer to arbitrary substrates. W.L.H. performed hydrothermal synthesis of STONC. A.J.C. cowrote the manuscript and supervised all nanoribbon formation. K.R.K. cowrote the manuscript and supervised all solution processing of the nanocube precursors. S.S.N. led the overall project, supervised general fabrication, local probe, and data analysis studies, and wrote the manuscript.

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REFERENCES

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DOI: 10.1021/acsami.7b00220 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX