Ambipolar Transport in Solution-Synthesized Graphene Nanoribbons

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Ambipolar Transport in Solution-Synthesized Graphene Nanoribbons Jia Gao, Fernando Javier Uribe-Romo, Jonathan D Saathoff, Hasan Arslan, Colin R. Crick, Sam J. Hein, Boris Itin, Paulette Clancy, William R. Dichtel, and Yueh-Lin Loo ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.6b00643 • Publication Date (Web): 05 Apr 2016 Downloaded from http://pubs.acs.org on April 6, 2016

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Ambipolar Transport in Solution-Synthesized Graphene Nanoribbons Jia Gao,1‡ Fernando Javier Uribe-Romo,2‡ Jonathan D. Saathoff,3 Hasan Arslan,2 Colin R. Crick,2 Sam J. Hein,2 Boris Itin,4 Paulette Clancy,3 William R. Dichtel,2 and Yueh-Lin Loo1* 1

Department of Chemical and Biological Engineering, Princeton University, Princeton, NJ

08544, USA 2

Department of Chemistry and Chemical Biology, Cornell University, Baker Laboratory, Ithaca,

NY 14853-1301 USA 3

School of Chemical and Biomolecular Engineering, Cornell University, Ithaca, NY 14853 USA

4

The New York Structural Biology Center. 89 Convent Ave., New York, NY 10027 USA

*E-mail: [email protected].

KEYWORDS: solution-synthesized graphene nanoribbons, ambipolar transport, field-effect devices, aerosol-assisted chemical-vapor deposition, inter-ribbon aggregation

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ABSTRACT Graphene nanoribbons (GNRs) with robust electronic band gaps are promising candidate materials for nanometer-scale electronic circuits. Realizing their full potential, however, will depend on the ability to access GNRs with prescribed widths and edge structures and an understanding of their fundamental electronic properties. We report field-effect devices exhibiting ambipolar transport in accumulation mode composed of solution-synthesized GNRs with straight armchair edges. Temperature-dependent electrical measurements specify thermally activated charge transport, which we attribute to inter-ribbon hopping. With access to structurally precise materials in practical quantities and by overcoming processing difficulties in making electrical contacts to these materials, we have demonstrated critical steps towards nanoelectric devices based on solution-synthesized GNRs.


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Despite its high carrier mobility and saturation velocity, the zero-bandgap semi-metallic electronic structure of graphene severely limits its utility for logic circuits.1,2 Reducing one of its two dimensions confers an energy gap that is inversely proportional to the width of the resulting graphene nanoribbon (GNR).3,4 Top-down approaches to lithographically define the dimensions of graphene or unzip carbon nanotubes5-8 have been demonstrated. But bottom-up approaches that embody the element of design can provide exquisite control over the structure and— ultimately—the electronic properties of GNRs.9-11 Of particular interest are solution-based syntheses of GNRs that offer a high degree of chemical and structural diversity yet specificity, and improved scalability over surface-assisted growth methods.12,13 Even with access to chemically well-defined GNRs, significant barriers remain to integrate these materials into functional electronic devices. To date, only two reports have demonstrated electrical activity in solution-synthesized GNR-based devices.14,15 Despite utilizing the same GNR, one report featured transistors operating in depletion mode, while the other highlighted hole-only devices operating in accumulation mode despite the fact that these materials are nominally electron-rich. Taken together, these reports underscore the processing difficulties and challenges associated with making electrical contact to these carbon nanostructures. We demonstrate ambipolar field-effect devices (FEDs) based on solution-synthesized GNRs with armchair edges. That we observe electron- in addition to hole transport is an indication that we are probing the intrinsic character of these materials. While the performance of these devices remains contact-limited, temperature-dependent electrical characterization reveals thermally activated charge transport.


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Results and Discussion Distinct from prior syntheses, the GNRs used in this study were prepared through the benzannulation of each alkyne moiety of a poly(p-phenylene ethynylene) (PPE), followed by cyclodehydrogenation to planarize the resulting poly(arylene) into a GNR (Figure 1a; Figures S1-S23). PPEs are among the most structurally diverse conjugated polymers, and high molecular weight samples are easily obtained. The benzannulation reaction is highly efficient16 and tolerant of steric hindrance, even along the polymer backbone, and provides poly(arylene)s that—unique to this study—are oxidized to GNRs with armchair edges (AGNRs).17-19 AGNRs with widths of thirteen carbon atoms (which translates to a width of 1.5 nm) and hexa(ethylene oxide) side chains (GNR-OHxg) were derived from PPE-OHxg. In general, AGNRs can be grouped into three families, N = 3p, N = 3p + 1 and N = 3p + 2, where N is the number of carbon atoms along the AGNR width and p is an integer.20 AGRs with N = 3p+1, including our GNR-OHxg, is predicted to exhibit larger band gaps than the other two families of AGRs for any given value of p.21 The OHxg solubilizing group was selected based on Molecular Dynamics (MD) simulations showing that the group was soluble and sufficiently long to wrap around the GNR and interfere with aggregation (Figures S24 and S25). PPE-OHxg was prepared via Sonogashira crosscoupling

polymerization

between

a

substituted

1,4-diethynylbenzene

monomer

and

1,4-diiodobenzene, and the resulting polymer was characterized spectroscopically (Figures S9 and S17) and via size exclusion chromatography/multi-angle light scattering (SEC/MALS; see Tables S1 and S2). PPE modification and GNR synthesis were characterized by monitoring changes in the

13

C NMR and Raman spectra of two PPE-OHxg samples: one with naturally

occurring isotopic substitutions and a second whose alkyne carbons were

13

C-enriched (99%).

The 13C NMR and Raman spectra of these polymers are indicative of the expected structure with


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no evidence of butadiyne defects.22 The PPEs were subsequently benzannulated by 1, CF3COOH, and Cu(OTf)2, to provide polyarylene BPP-OHxg (Figure 1b), for which SEC/MALS (Figures S1-S5) indicated increased molecular weight corresponding to newly installed 2,3-diarylnaphthalene moieties (Table S1). BPP-OHxg also exhibited shifts in its UV/Vis and photoemission spectra (Figure 1c) relative to PPE-OHxg that were very similar to previous reports of benzannulated PPEs and other ortholinked phenylene oligomers.16,23 In particular, the absorbance of BPP-OHxg is blue-shifted significantly (λmax = 270 nm) relative to that of PPE-OHxg (λmax = 420 nm), which arises from its inability to adopt a coplanar conformation along its backbone and is a spectroscopic signature for ortho-linked phenylenes (Figure 1c). The photoemission of BPP-OHxg broadens, is redshifted, and has lower quantum yield (λF = 1.4% vs. 7.1%, see Supporting Information) compared to PPE-OHxg, which is also consistent with the photoemission observed for polyo-phenylene oligomers.23 The solution 1H NMR spectrum of BPP-OHxg is poorly resolved because the polymers adopt kinetically trapped conformations whose resonances do not coalesce below available temperatures for solution NMR experiments.16

13

C isotopic labeling studies of BPP-OHxg

indicate the formation of well-defined structures by solid-state NMR (Figure 1d). Two isotopically enriched BPP-OHxg derivatives were synthesized: BPP-OHxg-13C(2,3) (Figure 1d, red) with

13

C enrichment at the 2 and 3 positions of the newly introduced naphthalene subunits

and BPP-OHxg-13C(1) (Figure 1d, orange) with isotopic labeling at the 1-position.24 These 13C isotopic labels are found in distinct positions and bonding environments. Measurements of the peak intensity as a function of cross-polarization time (tCP) and interrupted decoupling time (tID) in the CP-MAS experiment indicated that the respective resonances correspond to quaternary


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naphthalene carbons for BPP-OHxg-13C(2,3) and

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C bound to a single hydrogen for BPP-

OHxg-13C(1) (Figures S18 and S19), each consistent with the expected structures. These spectroscopic and isotopic labeling studies indicate that the benzannulation of PPE-OHxg proceeds efficiently to produce well-defined polyarylenes amenable for oxidation to GNRs. Oxidation

of

(DDQ)/CH3SO3H

the in

BPP-OHxg CH2Cl2

yielded

using

2,3-dichloro-5,6-dicyano-1,4-benzoquinone

GNR-OHxg,

which

is

dispersible

in

N,N-

dimethylformamide (DMF) and N-methyl-2-pyrrolidone (NMP) up to concentrations of ca. 0.5 mg mL-1. The Raman spectrum of GNR-OHxg (Figure 2a) exhibits features that are characteristic of GNRs,3 including D and G bands at 1365 and 1600 cm-1, respectively, as well as peaks at 2663 (2D overtones), 2935 (2G or G' overtones) and 3164 cm-1 (D+G combination modes). The sp3hybridized C-H vibrational modes centered at 2870 cm-1 in the FTIR spectrum of GNR-OHxg (Figure 2b) indicates that the OHxg side chains are retained upon Scholl oxidation of BPPOHxg. The spectrum also shows attenuated sp2-hybridized C-H stretching modes at 3060 cm-1 relative to that of BPP-OHxg, a signature of cyclodehydrogenation.

13

C CP-MAS spectra of

GNR-OHxg (Figure S20) show two resonances at 120 and 70 ppm, consistent with the presence of carbons in graphitic and ethylene oxide bonding environments. Similarly, the

13

C CP-MAS

(Figure 1c) spectra of GNR-OHxg-13C(2,3) and GNR-OHxg-13C(1) exhibit resonances in the graphitic region of the spectrum only, consistent with the corresponding labeling position and the formation of a ribbon structure. Additional structural evidence for GNR-OHxg via X-ray photoelectron spectroscopy and absorption spectroscopy are provided in Supporting Figures S6S8 and S14-S16. Collectively, our suite of characterization reveals the absence of any structural or edge defects down to the resolution limit of these techniques. Given that the complete set of


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characterization data suggests clean cyclodehydrogenation, we estimate the average length of GNR-OHxg to be approximately 50 nm from the molecular weight of the polymer precursor, BPP-OHxg. GNR-OHxg shows no observable photoemission in the visible range of the spectrum, and instead emits in the near-infrared (NIR) with peaks at of 1409 and 1204 nm, corresponding to photon energies of 0.88 and 1.03 eV, respectively (Figure 2c). That our solution-synthesized GNR-OHxg emits well into the NIR region can be attributed to its extended conjugation. To estimate the band structure of a single GNR-OHxg ribbon, the band structure of a single GNR terminated by methoxy groups (GNR-OCH3) was calculated by Density Functional Theory (DFT) using the HSE screened hybrid functional.25,26 This functional was chosen because HSE results for graphene nanoribbons have been shown to be very similar to those derived from the more rigorous GW/Bethe-Salpeter Equation.27 These calculations suggest a band gap of 1.06 eV (Figure 2d; Figures S26-S28) for individual GNRs. For comparison, the band gap of hydrogen-terminated GNRs (GNR-H) was calculated to also be very similar, at 1.1 eV. Although GNR-OHxg is dispersible in DMF, its limited dispersibility precluded uniform deposition for device fabrication. Instead, we employed aerosol-assisted chemical-vapor deposition (AA-CVD) to deposit GNR-OHxg over large areas.28 Figure 3a shows optical micrographs of gold patterns defined by photolithography on silicon substrates onto which GNR-OHxg was deposited. The corresponding Raman map tracks the G-band intensity at 1600 cm-1 and reveals preferential deposition of GNR-OHxg on the gold patterns compared to the unpatterned regions of native silicon oxide after AA-CVD of GNR-OHxg at 250 °C. The Raman spectrum of deposited GNR-OHxg on Au is shown in Figure 3b. At 0.78, the D/G peak ratio of AA-CVD-deposited GNR-OHxg is identical to that of the GNR-OHxg prior to deposition


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(Figure 2a). This invariance in the D/G peak ratio indicates AA-CVD to be a benign procedure having negligible impact on the integrity of GNR-OHxg. The thermogravimetric data provided in Figure S23 further supports this claim. We observe similar preferential deposition of GNROHxg on other metals, including Cu and Ag, compared to on SiO2. Control experiments indicate this preference during deposition to be directly correlated with differences in the thermal conductivities of substrates. (Figures S31 and S32). We surmise that the more thermally conductive substrates must more effectively flash off the solvent as the aerosol comes into contact, leaving behind GNR-OHxg. Given this enhanced adsorption on metal surfaces, we opted to detail the electrical characteristics of FEDs that were fabricated by first depositing GNR-OHxg via AA-CVD on Cu foil, then transferring29,30 the GNR-OHxg onto dielectric surfaces (Figure 3c). FEDs were also fabricated by directly depositing GNR-OHxg on dielectric surfaces at higher temperatures; the electrical characteristics are qualitatively comparable with those of devices comprising transferred GNR-OHxg and are provided in Supporting Information for completeness. Using this approach, we can routinely transfer 12-15 nm thick GNR-OHxg onto SiO2; the surface morphology of one such film is shown as an inset in Figure 3d. Molecular Dynamics simulations estimate that, depending on the number of OHxg groups that intervene between GNRs, the thickness of a single GNR-OHxg accounting for its side chains varies between 0.3 nm and 1.3 nm. Given an experimentally determined height of 12 nm, the channels in our FEDs comprise 9-40 layers of GNR-OHxg (Figure S25). FEDs were tested under vacuum with relatively low yields (approximately 2%), which presumably reflects the challenge in having GNR-OHxg bridge the 60-nm gap between the source and drain electrodes. Of the functional FEDs, 60% exhibit electron-transporting characteristics and the remainder show ambipolar electrical characteristics. While ambipolar


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electrical characteristics have been reported in GNRs that are derived from top-down approaches,11,31 there have yet been studies on solution-synthesized GNRs that exhibit ambipolarity. Prior GNR-based devices have either operated only in depletion mode14 or solely transported holes,15 the lack of ambipolarity in these prior devices suggests either the presence of substantial edge defects or doping of GNRs, presumably during device fabrication and/or electrical characterization. In contrast, our observation that all the functional GNR-OHxg-based devices show electron transport in accumulation mode indicates that device operation is not limited by electron traps at the charge transport interface or by adsorbed oxygen and/or water that act as p-type dopants.32-34 Rather, we must be probing the intrinsic electronic properties of solution-synthesized GNR-OHxg. We present the electrical characteristics of a representative ambipolar FED in Figure 4. The chevron-type transfer curves shown in Figures 4a and b indicate that our device transports both holes and electrons in standard accumulation mode within the range of temperatures explored (78–200 K). The gate current, shown as dashed lines in Figures 4a and b, is approximately 2% of the drain current; gate leakage is thus negligible and does not impact our analysis. The formalism for analyzing the electrical characteristics of conventional field-effect transistors (FETs) specifies the minimum source-drain conductance point (Gmin) as the point at which the operation of the device converts from ambipolar to unipolar transport.35 Following this formalism, Figure 4c tracks the evolution of Gmin for electron and hole transport extracted from the transfer curves as a function of temperature. The Gmin for electron transport shifts from 6.5 V to -5.5 V as the temperature is increased from 78 K to 200 K. This shift indicates that a smaller gate bias is required to form a unipolar electron-transporting channel at higher temperatures. This dependence is expected and has been previously observed in devices having GNRs prepared


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from unzipping multi-walled carbon nanotubes.36 Interestingly, Gmin also shifts towards more negative bias with increasing temperature in the hole-transporting regime, indicating that larger gate biases are required to deplete electrons and induce hole transport. We speculate that this unusual response arises from a shift in the Fermi level of GNRs towards vacuum level with increasing temperature. As such, a more negative gate bias is required to deplete electrons from the channel. Even accounting for the presence of a Schottky barrier for charge injection, this shift in the Fermi level necessarily modifies the barrier height in the same direction, effectively favoring electron over hole injection at elevated temperatures. The temperature dependence of the minimum IDS for electron transport is shown in Supporting Information (Figure S33). Several models31, 36, 37 exist for extracting the band gap of GNRs from device characteristics. One widely accepted model31 is to fit Imin at different temperatures to the Arrhenius equation and extract the activation energy. The band gap is then estimated as twice the activation energy. While this model can adequately describe devices with near-ohmic electrical contacts, it does not account for the Schottky barrier for charge injection that is present in our devices. As such, we have chosen instead to extract the band gap from the output characteristics provided in Figure 4d using the Metal-Semiconductor-Metal (M-S-M) model.38-40 Figure 4e reveals that at a constant VDS of 10 V, the electron current (IDS) exhibits a weak dependence on VG. The extracted barrier height for electron injection decreases from 21 to 19 meV as VG increases from 0 to 40 V. Figure 4f shows that the barrier height decreases linearly with an increase in the square root of VDS. This square root dependence on VDS suggests the presence of image force at the GNR-OHxgelectrode interface that effectively lowers the Schottky barrier with increasing VDS.39 The maximum barrier height, extrapolated at VG = VDS = 0 V, is 100 and 92 meV for electron and hole injection (inset in Figure 4f), respectively. At twice the maximum barrier height,40 the band


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gap of the GNR-OHxg network estimated from the output characteristics is ca. 200 meV. While this estimated band gap for the transferred GNR-OHxg network is substantially smaller than what we had expected given an optical band gap of 1.1 eV from absorption measurements on GNR-OHxg dispersions and based on DFT calculations, it is consistent with the solid-state absorption spectrum of the transferred GNR-OHxg network that reveals broad absorbance down to < 0.4 eV (Figure S34) and is further consistent with the low on-off current ratios in our fieldeffect devices (1.8 to 3.1). In fact, the band gap for our GNR-OHxg is higher than that reported for ion-beam patterned GNRs (88 meV), the devices constructed from which also exhibited low on-off current ratio of 2 to 5.37 We believe the discrepancy in the band gaps of GNR-OHxg in dispersion and in the solid state stems from the aggregation of the ribbons. To assess the impact of aggregation, we calculated the band gaps of pairs of GNR-OCH3 and GNR-H, using a Density Functional Theory representation of their single-ribbon optimized geometries. In these calculations, the GNRs—like graphite—are Bernal stacked with two GNR faces separated by 3.35 Å and the top GNR shifted one bond length along the ribbon’s long axis relative to the bottom GNR. At 0.71 and 0.72 eV, the band gaps of pairs of GNR-OCH3 and GNR-H ribbons are smaller than those of isolated ribbons, with similar trends seen in other work.41,42 The presence of additional layers further decreases the band gap. The band gap of GNR-H is also very sensitive to the relative positions of the two GNRs (Figure S29 and S30). Laterally shifting two GNRs in-plane at a constant separation of 3.35 Å, for example, resulted in band gaps ranging from 0.3 to 1.0 eV. The variable-temperature measurements on FEDs based on solution-synthesized GNRs shed light on its mechanism of charge transport. As a first pass, we extracted temperature-dependent mobilities of GNR-OHxg from the linear regime of the output characteristics following the


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guidance of Tian et al., in which this approach was used to analyze the electrical characteristics of MoS2-based field-effect Schottky barrier transistors.43 Like GNRs that are produced by reactive-ion etching of graphene,44-46 we observe simple Arrhenius-type temperature dependence of both the electron and hole mobilities of GNR-OHxg. Quantification resulted in activation energies of 17 and 12 meV for electron and hole transport, respectively (Figures S35-S37). A similar activation energy of 13 meV for electron transport is extracted from electron-only FEDs whose operation is detailed in Figure S35. Previous studies have shown the edge structure of GNRs to critically influence electrical properties. While band-like charge transport36,

47

was

observed in graphene nanoribbons with smooth edges, thermally-activated electrical behavior was reported for graphene nanoribbons having edge defects.48,49 Given that characterization of GNR-OHxg suggests the presence of minimal edge defects, we are left to surmise that the thermally activated electrical characteristics is a reflection of inter-ribbon charge transport through a percolated network of GNR-OHxg that connects the source and drain electrodes. This activation energy reflects the charge transport barrier at GNR-GNR homojunctions. That these activation energies are comparable to those reported for charge transport across graphene grain boundaries50 and between graphene quantum dots51 indicate inter-ribbon transport to be the bottleneck in our devices.

Conclusions We have demonstratedFEDs exhibiting electron transport in accumulation mode based on solution-synthesized GNRs with well-defined edges. Electrical characterization of FEDs confirms their potential utility for electronic device applications. Variable-temperature measurements specify thermally activated transport in this class of materials and provide a


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quantitative description of their electronic properties. While there is uncertainty in the absolute magnitude of the band gap of GNR-OHxg, several approaches to estimate it from the electrical characteristics of GNR-OHxg-comprising devices consistently reveal the band gap to be substantially smaller than that of isolated ribbons. The < 200 meV band gap is consistent with stacking of GNR-OHxg, and implicates aggregation of these materials to remain a major challenge in realizing their full potential.

Materials and Methods Materials. All starting materials and solvents, unless otherwise specified, were obtained from commercial sources and used without further purification. All reactions were performed at ambient laboratory conditions, and no precautions were taken to exclude oxygen or atmospheric moisture unless otherwise specified. Anhydrous solvents: tetrahydrofuran (THF), methylene chloride, acetonitrile (MeCN), N,N-dimethylformamide (DMF), toluene were purified using a custom-built, alumina-column based solvent purification system. N-methyl-2-pyrrolidone (NMP) was distilled under vacuum (10 mtorr) from molecular sieves. Diisopropylamine and 1,2dichloroethane (DCE) were distilled from CaH2 under N2 atmosphere. Trimethylsilylacetylene13

C2 (99% atom 13C) was obtained from Aldrich Chemical Co. Synthesis of compounds S1-S6,

monomer S7 and isotopically enriched materials are provided in the Supporting Information. Compounds S5, S5-13C4, and 1 were prepared according to published methods.15 General Synthesis of PPE-OHxg. Monomer S7 (41 mg, 0.124 mmol) 1,4-diiodobenzene (41 mg, 0.124 mmol) and 4-tert-butyl-iodobenzene (See Supporting Table S1 for amounts), Pd(PPh3)4 (7 mg, 0.006 mmol) and CuI (2 mg, 0.010 mmol) were loaded in a flame dried 50 mL Schlenk flask equipped with a magnetic stirrer. The flask was evacuated under dynamic vacuum


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to 150 mtorr and refilled with N2 three times. Under N2 flow, anhydrous MeCN (12.4 mL) and anhydrous iPr2NH (4.2 mL) were added via cannula. The mixture was stirred at 40 °C for 4 days covered from light, after which was transferred in air to a round bottom flask and the mixture was concentrated to dryness in a rotary evaporation. The obtained dark green oil was suspended in CH2Cl2 (1 mL) and added dropwise into Et2O (500 mL) forming a brown pale precipitate in a yellow solution and the suspension was allowed to rest for 6 h. The brown precipitate was separated by vacuum filtration using a 2 µm nylon filter membrane, rinsed with water, HCl 1M (aq), methanol and Et2O until rinse solvent was colorless. The obtained solid was dried under dynamic vacuum (10-2 torr) for 12 h to afford PPE-OHxg (118 mg, 92% yield) as a dark brown waxy solid. 1H NMR (500 MHz, dmso-d6, 25 °C) δ 8.20 (br), 7.84 (br), 7.70 (br), 7.39 (br), 7.28 (br), 7.23 (br), 7.11 (br), 4.27 (br), 3.85 (br), 3.65-3.39 (br), 3.18 (br).

13

C NMR (126 MHz,

dmso-d6, 25 °C) δ 156.79, 137.23, 133.56, 133.28, 132.39, 131.07, 130.91, 129.37, 128.25, 127.99, 127.50, 126.99, 126.06, 118.81, 106.89, 71.00, 69.79, 69.72, 69.60, 69.53, 69.30, 68.72, 68.69, 67.29, 57.64. FTIR (ATR, cm-1) 3058, 2867, 2205, 1627, 1603, 1493, 1389, 1349, 1261, 1232, 1196, 1098, 1005, 972, 936, 888, 852, 814, 748, 701, 669. General synthesis of BPP-OHxg. PPE-OHxg (80 mg of, 0.080 mmol, based on repeating unit) and Cu(CF3SO3)2 (10 mg, 0.027 mmol) were loaded in a 5 mL round bottom flask equipped with a magnetic stirrer. C2H2Cl4 (2 mL) was added and the mixture was degassed by three cycles of freeze-pump-thaw, backfilled with N2 and heated to 90 °C with stirring. 2-(2-phenyl-ethynyl)benzaldehyde 6 (206 mg, 1.152 mmol) and CF3COOH (30 µL, 44 mg, 0.046 mmol) were added at 90 °C with stirring in portions of 80 mg of 6 and 10 µL of acid each 3 hours. After 9 h, the mixture was cooled to room temperature, quenched with saturated NaHCO3 (aq). After effervescence stopped, the mixture was extracted with CH2Cl2 (3 × 50 mL), dried over


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anhydrous Na2SO4, filtered through celite and concentrated to dryness. The obtained brown oil was dissolved in C2H2Cl4 (1 mL) and precipitated into Et2O (250 mL), forming a brown solid in a pale yellow suspension. The mixture was allowed to rest 8 h after which the solid was isolated by vacuum filtration using a 2 µm nylon filter paper, rinsed with HCl 2 M (aq), water, MeOH and Et2O until rinse liquors were colorless. The obtained dark brown solid was dried under dynamic vacuum (10-2 torr) for 12 h affording 118 mg (91% yield). 1H NMR (500 MHz, dmsod6, 25 °C) δ 7.95 (br), 7.52-7.06 (br), 4.15 (br), 3.43 (br), 3.16 (br). 13C NMR (126 MHz, dmsod6, 25 °C) δ 138.79, 135.74, 132.68, 128.34, 71.80, 70.33, 70.10, 69.39, 67.96. FTIR (ATR, cm1

) 3055, 2867, 1627, 1602, 1487, 1447, 1394, 1348, 1244, 1197, 1100, 937, 888, 851, 751, 698.

General synthesis of GNR-OHxg. BPP-OHxg (21 mg, 0.017 mmol based on repeat unit) was loaded in a 500 mL Schlenk flask equipped with a magnetic stirrer. The flask was evacuated under dynamic vaccum to 150 mtorr and the filled with N2 three times. Anhydrous CH2Cl2 (80 mL) and CH3SO3H (10 mL) were added via cannula, and the solution was stirred and cooled to 0 °C in an ice/water bath. 2,3-dichloro-5,6-dicyano-1,4-benzoquinone, (DDQ, 0.114 g, 0.506 mmol) was dissolved in 2.5 mL of anhydrous CH2Cl2 and the mixture was bubbled with N2 for 5 min. The DDQ solution was added drop wise via syringe to the CH2Cl2 solution under N2, and the mixture was stirred at 0 °C for 1 h. The mixture was warmed up to room temperature and stirred for 48 h to form a dark solution. The mixture was quenched with saturated NaHCO3 (aq) yielding a black precipitate suspended in the organic phase. After effervescence stopped, the mixture was extracted with CH2Cl2 (3 × 100 mL). The black solid (present in both aqueous and organic phases) was separated by filtration from each phase and combined. The obtained black solid was rinsed with water, MeOH, EtO2 and CH2Cl2, and dried under dynamic vacuum (10-2


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torr) for 8 h affording GNR-OHxg (18 mg 70% yield). FTIR (ATR, cm-1) 3064, 2865, 1717, 1447, 1348, 1240, 1198, 1095, 948, 880, 829, 756, 701. Characterization. High-resolution 1H,

13

C and

11

B nuclear magnetic resonance (NMR) spectra

were recorded using Varian INOVA-400, 500 and 600 MHz spectrometers. Solid-state NMR was recorded at ambient pressure in a Varian INOVA-400 spectrometer using a magic anglespinning (MAS) probe with 7 mm (outside diameter) silicon nitride rotors. The magic angle was adjusted by maximizing the number and amplitudes of the signals of the rotational echoes observed in the

79

Br MAS FID signal from KBr. Cross-polarization with MAS (CP-MAS) was

used to acquire 13C data at 100.7 MHz. The CP contact time varied from 100 to 7000 µs. High power two-pulse phase modulation (TPPM) 1H decoupling was applied during data acquisition. The MAS sample-spinning rate was 8.2 kHz. Recycle delays between scans varied between 3 and 10 s, depending upon the compound as determined by observing no apparent loss in the 13C signal from one scan to the next. High-field, solid-state NMR was measured at The New York Structural Biology Center using a Bruker Avance I spectrometer operating at a 1H frequency of 750 MHz and equipped with 4-mm HX probe. Spinning frequency was 15 kHz. All spectra were acquired at temperature of –5 °C. RF fields were 100 kHz for 1H and 50 kHz for

13

C, using 3 ms of spin locking time and 50%

linear ramp for 1H-13C cross polarization and TPPM composite pulse decoupling. Recycle delay was 3 s, number of transients varied between 1000 and 20,000. The 13C chemical shifts are given relative to tetramethylsilane as zero ppm, calibrated using the methylene carbon signal of adamantane, assigned to 37.77 ppm, as secondary reference.


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Matrix-assisted laser desorption-ionization time-of-flight mass spectrometry (MALDI-TOF MS) was performed in a Waters MALDI Micro MX, 7,7,8,8-tetracyano-quinodimethane (TCNQ) was used as matrix and the mass ranges were calibrated using poly(ethylene oxide) standards. Fourier-transform infrared (FTIR) spectra were recorded using a Thermo Nicolet iS10 with a diamond attenuated total reflectance (ATR) attachment. A total of 64 transients were collected for each sample with a resolution of 0.05 cm-1. Data presented in absorbance mode was processed as follows: raw data was converted to absorbance, smoothed (0.783 cm-1 level), and the baseline was subtracted with a zero-set 3200 and 2650 cm-1 with a linear baseline in between these frequencies. Data presented in % transmittance was uncorrected. Ultraviolet/visible/near infrared (UV/vis/NIR) absorbance spectra were recorded on a Cary 5000 spectrophotometer with a mercury lamp. Quartz cuvettes rated for transparency in the near infrared region containing the pure solvent of interest for the measurement was used for baseline. Photoemission and excitation spectra were recorded on a Horiba Jobin Yvon Fluorolog-3 fluorescence spectrophotometer equipped with a 450 W Xe lamp, double excitation and double emission monochromators, a digital photon-counting photomultiplier and a secondary InGaAs detector for the NIR range. Correction for variations in lamp intensity over time and wavelength was achieved with a solid-state silicon photodiode as the reference. The spectra were further corrected for variations in photomultiplier response over wavelength and for the path difference between the sample and the reference by multiplication with emission correction curves generated on the instrument. The fluorescence quantum yield, ΦF, was determined using anthracene in cyclohexane as reference; its ΦF = 36% ± 0.04%.52 The Raman spectra shown in Figure 2 were recorded on a Renishaw InVia confocal Raman microscope with excitation wavelengths of 488 nm and 785 nm. Dispersions were drop-cast on


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borosilicate glass slides or silicon wafers, solid samples were compressed between borosilicate glass slides. Backgrounds were fitted with a polynomial function and subtracted from the sample spectra. Raman mapping of the G-band intensity, shown in Figure 3a, and the Raman spectrum of AA-CVD-deposited GNR-OHxg, shown in Figure 3b, were performed using a Horiba ARAMIS Raman spectrometer with an excitation wavelength of 532 nm. Thermogravimetric analysis from 20-600 °C was carried out on a TA Instruments Q500 Thermogravimetric Analyzer under a N2 flow using a 5 K min-1 scan rate. Size exclusion chromatography with multi-angle laser light scattering (SEC-MALLS) with THF as mobile phase was performed on two 7.5-µm columns (PolyPore, Varian, Inc.) connected in series. THF was the mobile phase at 1.0 mL min-1 flow from a Shimadzu LC-20AD isocratic pump. The detectors consisted of a miniDawn S2 three angle, light-scattering system, followed downstream by an Optilab Rex differential refractometer from Wyatt Technologies. Samples were prepared in 1 mg mL-1 concentration in BHT-stabilized THF and filtered through a 2 µm inorganic membrane filter prior to injection. Average molecular weights were determined using one-point on-line dn/dC. X-ray photoelectron spectroscopy (XPS) was measured in a Surface Science Instruments XProbe SSX-100 spectrometer with operating pressure < 10-9 torr using monochromatic Al K Xrays at 1486.6 eV. Photoelectrons were collected at an angle of 55o from surface normal in a hemispherical analyzer with pass energy of 50 V and detected with a SSI Position Sensitive, resistive anode, 40 mm × 40 mm detector, electronically defined as 128 active channels with maximum count rate of 1,000,000. Samples were prepared by drop-casting from dispersions in CHCl3, DMF or NMP onto gold-coated silicon wafers. The wafers were cleaned with acid piranha solution previous to use. Data analysis and deconvolutions were performed in CasaXPS:


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Processing software for XPS (Casa Software Ltd.). Peak deconvolutions were performed with Shirley backgrounds, Gaussian-Lorentzian-Product functions with 30% Lorentzian mixing (GL(30) function). Peak position, full-width-at-half-maximum and concentration were fitted and extracted. The AFM image shown as inset in Figure 3d was collected using a Veeco Dimension NanoMan AFM. Aerosol-assisted chemical vapor deposition. Solution-synthesized graphene nanoribbons with GNR-OHxg were deposited by AA-CVD to generate sufficient coverage for device fabrication. An aerosol containing GNRs was generated using an ultrasonic humidifier. The aerosol was transferred into a horizontal-flow hot-walled chemical-vapor deposition reactor, constructed from a quartz tube furnace (inside diameter 36 mm), using Argon as the carrier gas at a flow rate of 500 mL min-1. 30 mL of GNR-OHxg dispersion, at a concentration 50 µg L-1 in chloroform, was used for each deposition. The substrates were positioned at the center of the heating zone and tilted 15° relative to the bottom of the furnace to facilitate the deposition of GNR-OHxg during AA-CVD. The substrates were typically 16 mm × 16 mm. The set-point temperature of furnace was varied between 250 and 350 °C, with the upper limited dictated by chemical degradation of GNR-OHxg. Field-effect device fabrication and electrical characterization. Bottom-gate, top-contact fieldeffect devices comprising GNR-OHxg were fabricated on heavily doped Si substrates with 90nm thick thermally grown silicon dioxide as the gate dielectric. Contact pads having 5 nm of Ti and 30 nm of Au were defined by standard photolithography, electron-beam evaporation and liftoff.53 After the deposition of GNR-OHxg, E-beam lithography was performed to define source and drain electrodes comprising 10 nm of Ti and 40 nm of Au that electrically connect the


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contact pads. The device channel width and length were kept constant at 2 mm and 60 nm, respectively. A final photolithography step was performed to pattern the active channel region. Oxygen plasma treatment removed GNR-OHxg outside the channel regions and electrically isolated neighboring devices. Three-terminal devices were placed in a Lakeshore probe station (Lake Shore Cryotronics, Inc., Westerville, USA) and cooled with liquid nitrogen. The electrical characteristics were measured in vacuum (< 5 × 10-5 torr) using an Agilent 4155C semiconductor parameter analyzer as a function of temperature from 78 K to 200 K. Mobilities were calculated from transfer curves collected in the linear regime using equation (1) below: ௅

ଵ

ଵ

೚ೣ

ವೄ

u = ݃௠ × ௐ × ஼ × ௏

(1)

where gm is the transconductance, L and W are the channel length and width, respectively, of the device defined by the geometry of source and drain electrodes. VDS is source-drain bias, and Cox is the gate capacitance. Here, we estimated the capacitance by assuming a dielectric constant of 38.3 nF cm-2 for the 90-nm thick SiO2 gate dielectric. Molecular Dynamics Calculations. All Molecular Dynamics (MD) simulations were run using LAMMPS.54 In all simulations, the OPLS-AA force field was used to describe the bonded and non-bonded interactions of the system.55 Partial charges for NMP and graphene van der Waals parameters were altered to match those used by Shih et al., which had been validated against experiment.56 To generate Figures S24 and S25, we used a 10 Å cut-off for long-range interactions with a van der Wals tail correction. In addition, long-range Coulombic interactions were estimated using a particle-particle, particle-mesh approximation.57 This was done in a similar fashion to our previous MD work on GNRs.58 To calculate GNR-OHxg thicknesses, the


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simulations were run in vacuum with a 15 Å long-range cut-off without any additional longrange corrections. Additional information can be found in the Supporting Information. Density functional theory calculations. All Density Functional Theory (DFT) calculations were made using Gaussian 09.59 Periodic boundary conditions were used along a vector parallel to the GNRs. To reduce calculation costs, OHxg side chains were replaced with OCH3 side chains or the GNR edges were terminated with hydrogen atoms. The band structures of GNR-OCH3 and GNR-H were calculated using the HSE functional.25,26 For single GNRs, the geometry was optimized using the 6-31G(d) basis set, and the bands were calculated using the 6-311(d,p) basis set. For bilayer GNRs, the 6-31(d) basis set was applied to calculate the band structure. This basis set was found to reproduce the 6-311(d,p) band structure very closely and was significantly less computationally expensive. Additional information can be found in the Supporting Information.


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FIGURES

Figure 1. (a) Schematic depiction and (b) chemical reactions to transform PPEs (purple) into GNRs (green) featuring benzannulation of alkyne subunits on PPE-OHxg followed by Scholl oxidation. (c) UV-visible absorption and emission (inset) spectra of PPE-OHxg and BPP-OHxg (in 1,2-dichloroethane) (d) Partial OHxg-13C(1) (orange). Partial 13

13

13

C CP-MAS spectra of BPP-OHxg-13C(2,3) (red) and BPP-

C MAS NMR spectra of isotopically enriched GNR-OHxg-

C(2,3) (purple) and GNR-OHxg-13C(1) (black).


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Figure 2. (a) Raman spectrum (λex = 488 nm) of GNR-OHxg displaying the D, G, and the corresponding 2D and 2G overtones as well as D+G combination vibrational features. (b) FT-IR spectra of BPP-OHxg (red) and GNR-OHxg (green) indicate a decrease in intensity associated with the aromatic C-H stretching mode relative to the OHxg C-H stretching mode upon cyclodehydrogenation. (c) Near-infrared (NIR) photoemission spectrum (DMF, λex = 450 nm) of GNR-OHxg. (d) HSE-calculated band structure of single- and double-stacked GNR-OCH3 near the Fermi level (only 10 bands shown), indicating bandgaps of 1.06 and 0.71 eV, respectively.


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Figure 3. (a) Optical microscopy images of Au patterns on SiO2 (left column), and Raman mapping of the G band of the same substrates after they were subjected to AA-CVD of GNROHxg at 250 °C (right column), scale bar is 5 um. (b) Raman spectrum of AA-CVD-deposited GNR-OHxg on Au excited at 532 nm. (c) Scheme for transferring GNRs onto arbitrary substrates. (d) Height profile of GNR-OHxg on SiO2 after transfer from copper foil, inset shows the corresponding Atomic Force Microscope (AFM) image.


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Figure 4. (a, b) Transfer characteristics (IDS vs. VG) of a GNR-OHxg-based FED in its hole- and electron-transporting regimes at 78 K (black), 140 K (blue) and 180 K (red), the dot lines show the gate leakage current at those temperatures. (c) The evolution of the minimum conduction point (Gmin) for unipolar hole (red) and electron (blue) transport with temperature. (d) Output characteristics (IDS vs. VDS) of the same device at 180 K. (e) Arrhenius plot of IDS at VDS = 10 V with VG increasing from 0 to 40 V as a function of temperature, symbols are experimental data and the solid lines are fits to the data. (f) Plot of the effective barrier height for electrons (red


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square) and holes (blue square) as a function of the square root of VDS at VG= 0 V. Solid lines are linear fits to the data. Inset shows the extrapolation at VDS = 0.

ASSOCIATED CONTENT Supporting Information Syntheses and characterization of materials, fabrication and electrical characterization of GNROHxg-based device and computational calculations. This material is available free of charge via the Internet at http://pubs.acs.org AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]. Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ‡These authors contributed equally. Conflict of Interest The authors declare no competing ﬁnancial interest. ACKNOWLEDGMENT We gratefully acknowledge funding from the National Science Foundation through its Nanoelectronics Beyond 2020 Initiative (CHE-1124754) and co-sponsorship by the Semiconductor Research Corporation through its Nanoelectronics Research Initiative (2011-NE2205GB). W.R.D. acknowledges the Arnold and Mabel Beckman Foundation for a Beckman Young Investigator Award. J.G. also acknowledges the Netherlands Organisation for Scientific


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Research for a Rubicon post-doctoral fellowship (680-50-1202). This work made use of the Cornell Center for Materials Research Shared Facilities, which are supported through the NSF MRSEC program (DMR-1120296). The data collected at the New York Structural Biology Center (NYSBC) was made possible by a grant from NYSTAR. REFERENCES 1. Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Katsnelson, M. I.; Grigorieva, I. V.; Dubonos, S. V.; Firsov, A. A. Two-Dimensional Gas of Massless Dirac Fermions in Graphene. Nature 2005, 438, 197-200. 2. Avouris, P. Graphene: Electronic and Photonic Properties and Devices. Nano Lett. 2010, 10, 4285-4294. 3. Barone, V.; Hod, O.; Scuseria, G. E. Electronic Structure and Stability of Semiconducting Graphene Nanoribbons. Nano Lett. 2006, 6, 2748-2754. 4. Son, Y.-W.; Cohen, M. L.; Louie, S. G. Energy Gaps in Graphene Nanoribbons. Phys. Rev. Lett. 2006, 97, 216803. 5. Bai, J. W.; Duan, X. F.; Huang, Y. Rational Fabrication of Graphene Nanoribbons Using a Nanowire Etch Mask. Nano Lett. 2009, 9, 2083-2087. 6. Kosynkin, D. V.; Higginbotham, A. L.; Sinitskii, A.; Lomeda, J. R.; Dimiev, A.; Price, B. K.; Tour, J. M. Longitudinal Unzipping of Carbon Nanotubes to Form Graphene Nanoribbons. Nature 2009, 458, 872–876. 7. Campos-Delgado, J.; Romo-Herrera, J. M.; Jia, X.; Cullen, D. A.; Muramatsu, H.; Kim, Y. A.; Hayashi, T.; Ren, Z.; Smith, D. J.; Okuno, Y.; et al. Bulk Production of a New Form of sp2 Carbon: Crystalline Graphene Nanoribbons. Nano Lett. 2008, 8, 2773-2778.


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Graphical Table of Contents


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Ambipolar Transport in Solution-Synthesized Graphene Nanoribbons

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