Article pubs.acs.org/JPCC
Facile Determination of Bulk Charge Carrier Concentration in Organic Semiconductors: Out-of-Plane Orientation Hopping Conduction Characteristics in Semicrystalline Polythiophene Chang-Yong Nam* Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, New York 11973, United States S Supporting Information *
ABSTRACT: In this report, we demonstrate a straightforward two-terminal-device-based electrical measurement and analysis scheme that can simultaneously determine the bulk free charge carrier concentration and out-of-plane charge mobility in organic semiconductors by understanding the transition behavior of device current−voltage characteristics from ohmic to space charge limited conduction. As a model system, we characterize the properties of a semicrystalline poly(3-hexylthiophene) (P3HT) conjugated polymer film in which free carrier concentration is systematically controlled by adjusting oxygen doping level. The observed dependence of out-of-plane charge mobility on the carrier concentration is analyzed in the context of percolative variable range hopping conduction, and we identify the rate-limiting charge hopping process in P3HT and correlate it with the role of disordered polymer regions in mediating the charge transport between neighboring crystalline polymer lamellar domains.
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INTRODUCTION Organic semiconductors attract significant attention for their applications in electronic and energy conversion devices that potentially offer lightweight, innovative form factors, and low material and manufacturing costs. Particularly, organic conjugated polymers have shown important applications in energy conversion devices such as solar cells. Fundamental electronic material parameters such as bulk free charge carrier concentration and mobility affect the charge transport characteristics in conjugated polymers and associated device performances. For instance, the free carrier concentration influences device current on/off ratio and turn-on voltage of polymer organic field effect transistors (OFETs),1 as well as the contact barrier potential,2 charge recombination rate,3 and open circuit voltage4 in organic photovoltaic (OPV) devices. Similarly, the charge mobility is closely related with the switching speed of OFETs5 along with charge carrier collection and recombination and energy conversion efficiency in OPVs.6,7 Fundamentally, the free carrier concentration and mobility are interrelated properties with their correlation reflecting the intrinsic characteristics of hopping-type charge transport mechanism in organic semiconductors.8 Facile and accurate measurements of charge carrier concentration and © 2012 American Chemical Society
mobility are naturally important not only in determining new organic semiconductor systems suitable for high-performance electronic and energy conversion devices but also in understanding the basic charge transport processes in these materials. Several experimental schemes are available for measuring the charge carrier concentration and mobility in organic semiconductors. Conventionally, Hall measurement can quantify both carrier concentration and mobility, but its application to organic materials is nontrivial due to their low charge mobility and high resistivity.9 It also measures the in-plane mobility and can be less relevant in the organic energy conversion devices having an out-of-plane charge transport orientation, such as OPVs, due to the highly anisotropic charge transport characteristics in organic semiconductors.10,11 Meanwhile, the charge carrier concentration (conversely, doping concentration) in organic semiconductors has been determined by metal−insulator−semiconductor (MIS) junction devices, including OFETs1 and MIS diodes,12,13 where respective device turn-on gate voltage and capacitance−voltage characteristics Received: August 16, 2012 Revised: October 17, 2012 Published: October 24, 2012 23951
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Figure 1. (a) Scanning electron microscopy micrograph of glass/AuPd/P3HT/AuPd device cross-section after 3 h of dedoping vacuum annealing at 400 K (scale bar = 200 nm). (b) Double-logarithmic I−V plot of the dedoped P3HT device at Poxygen = 0.1 mTorr. The ohmic (slope = 1) and SCLC regions (slope = 2) are indicated by dashed and dash-dotted lines, respectively. The slope crossover occurs at VTh ≈ 0.14 V. (c) A representative current vs time plot at V = 0.1 V describing the oxygen doping procedure in the P3HT device. The dedoped sample in panel b was first exposed to the elevated Poxygen (20 mTorr) at time = 0 s in dark condition. Note that there is negligible change in current over time compared to the dedoped state (∼0.1 μA at V = 0.1 V in panel b), suggesting an insignificant level of oxygen doping under dark condition. When ∼1.8 mW/cm2 intensity light (from a halogen lamp) is illuminated on to the sample (at time = 200 s), there is an appreciable rise in photocurrent. Once the light is turned off, the sample stays at a higher conducting state even after a gradual decay of photocurrent, indicating that the light activates oxygen doping. The procedure can be successively repeated with increasing Poxygen to systematically increase the oxygen doping level in P3HT.
electrical contacts increases quadratically with increasing applied voltage (V) and, importantly, in proportion to the bulk charge mobility (μ) with the following relationship:22 JSCLC = (9/8) εoεrμ (V2/d3), where εo is the vacuum permittivity, εr is the relative dielectric constant of the semiconductor layer, and d is the device active layer thickness. Here, the magnitude of SCLC relies on the amount of the excess charge carriers injected from the contact to the semiconductor (nSCLC), which in turn increases with applied voltage by the following expression derived by equating JSCLC to a drift current density J = neμ(V/d), with charge concentration n and elementary charge e:
could be related to the carrier concentration. Planar OFETs are also widely used to characterize the charge mobility,14,15 but the in-plane charge transport confined near the semiconductor− dielectric interface is not compatible with the out-of-plane bulk charge transport in vertical geometry devices, similarly to Hall measurement. Alternatively, transient photocurrent techniques, including time-of-flight16,17 and photo-CELIV (charge extraction with a linearly increasing voltage)18,19 measurements, can determine the mobility along the out-of-plane orientation, but they may require relatively thicker organic semiconductor layers, compared to OPVs, and/or somewhat complicated model-dependent analyses.9 In this article, we demonstrate a straightforward twoterminal-device-based electrical measurement and analysis scheme that can simultaneously determine the bulk free charge carrier concentration and out-of-plane charge mobility in organic semiconductors by understanding the transition behavior of device current−voltage characteristics from ohmic to space charge limited conduction. As a model system, we characterize the properties of a semicrystalline poly(3hexylthiophene) (P3HT) conjugated polymer film in which free carrier concentration is systematically controlled by adjusting oxygen doping level. The observed dependence of out-of-plane charge mobility on the carrier concentration is analyzed in the context of percolative variable range hopping conduction, and by combining the X-ray scattering analysis of the internal polymer structure, we indentify the rate-limiting charge hopping process in P3HT and correlate it with the role of disordered polymer regions in mediating the charge transport between neighboring crystalline polymer lamellar domains.
nSCLC =
9 εoεr V 8 e d2
(1)
These injected excess charge carriers compete with the preexisting bulk free carriers, whose concentration (nohmic) is invariant under increasing bias and responsible for the ohmic current (Johmic = nohmiceμ(V/d)) in a low voltage bias regime. Since nSCLC increases with applied bias, naturally there exists a threshold voltage VTh at which nSCLC = nohmic =
9 εoεr VTh 8 e d2
(2)
For V > VTh, the charge conduction mode now transitions from ohmic to spaced charge limited regime as nSCLC > nohmic, with overall J−V characteristics following SCLC behavior. This competing behavior between ohmic and space charge limited conduction, originally described by Lampert in 1950s,23 enables the determination of nohmic without knowing μ by experimentally measuring VTh at which a lower voltage J ∝ V relationship changes to J ∝ V2 in a higher bias range (i.e., the change of slope from one to two with increasing voltage in a double-logarithmic current−voltage (I−V) plot). We applied this method to characterize the nohmic in a P3HT film having systematically varied oxygen doping levels and studied the correlation between nohmic and out-of-plane mobility (μ⊥) obtained from both ohmic and space charge limited conduction characteristics using a vertical geometry device structure. Widely utilized in OFETs and OPVs, P3HT is a
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RESULTS AND DISCUSSION Being relatively easy to implement and analyze, the space charge limited current (SCLC) measurement has been widely used to determine the out-of-plane bulk charge mobility in organic semiconductors.20,21 Specifically, if the semiconductor is charge-trap-free or having energetically shallow or completely filled charge traps, the current density limited by space charge in the semiconductor layer (JSCLC) between two ohmic 23952
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higher values (Figure 2a), indicating that nohmic increased with Poxygen. The SCLC in a higher voltage range (V > ∼3 V),
solution-processable semicrystalline conjugated polymer that has anisotropic charge transport properties due to the lamellar stacking of polymer chains perpendicular to a substrate surface, and the concomitant in-plane alignment of π−π stacking of polymer backbones, where the latter orientation supports charge transport more efficiently than the vertically stacked lamella that have insulating alkyl chain interlayers.10 In P3HT, a molecular oxygen acts as a p-type dopant and increases both nohmic and μ by forming a charge transfer complex with a thiophene ring.1,24 This doping process is enhanced by light12 and can be reversed by heating the polymer in vacuum.25 The vertical P3HT device structure in our study comprises a ∼300 nm thick P3HT film (spin-cast in air) sandwiched between 30 nm thick bottom and top AuPd contacts with a crossbar geometry (device area A = 1 mm2), as the device cross-section is shown in Figure 1a. In the experiment, in order to drive out the oxygen inadvertently incorporated into the P3HT active layer during the film processing in air, we first thermally annealed the P3HT device in a dark condition at 400 K for 3 h under 0.1 mTorr vacuum (i.e., dedoping). In a double-logarithmic I−V plot, the dedoped P3HT exhibits a clear transition from ohmic (slope of one) to space charge limited conduction (slope of two) at VTh ≈ 0.14 V (Figure 1b), which yields nohmic = 2.9 × 1014 cm−3 from eq 3, and εr ≈ 3 for P3HT. This carrier concentration is at least 1 order of magnitude lower compared to what has been reported in unannealed P3HT,1 supporting an effective dedoping by the vacuum anneal. We then quantified the μ⊥ corresponding to the obtained nohmic from the ohmic region (V < VTh) using the relationship μ⊥= σ/(nohmice) with ohmic conductivity σ extracted from a linear I−V plot. We find μ⊥ = 5.2 × 10−5 cm2 V−1 s−1, which reasonably matches the value (6.0 × 10−5 cm2 V−1 s−1) extracted independently from nohmic in the space charge limited regime (VTh > V) using the slope of a linear portion of √I−V data ((9Aεoεrμ⊥/(8d3))1/2), confirming the negligible influence of electric field and validity of the measurement scheme. We use these procedures to obtain the μ⊥ vs nohmic data by systematically increasing the oxygen doping level of the active P3HT layer and discuss their implication on the charge transport in the semicrystalline polymer film. It is noted that the measured high-bias I−V characteristics do not display any sign of deviation from the perfect space charge limited conduction behavior (slope of two) throughout the different doping conditions, reconfirming the negligible influence of electric field on the measured mobility in our probed bias range. We achieved the controlled oxygen doping in the initially dedoped P3HT by increasing the sample chamber oxygen pressure (Poxygen) and illuminating ∼1.8 mW/cm2 intensity light (halogen lamp) on to the device for ∼20−800 s (Figure 1c). Here, the light is diffuse-scattered from the vacuum chamber wall and transmitting through the thin semitransparent top metal contact, reaching the sandwiched P3HT layer. Consistent with previous reports,12,25 combining increased Poxygen with light illumination was essential for achieving an appreciable oxygen doping, and increasing Poxygen alone negligibly affected the P3HT conductance. We find that VTh and corresponding nohmic of P3HT increase with elevating oxygen doping levels. When the doping procedure was repeated sequentially with increasing Poxygen (up to 750 Torr), the corresponding I−V characteristics of the doped P3HT displayed a substantial increase (∼100×) of ohmic current in a low voltage region (V < ∼3 V), with the VTh extending to
Figure 2. (a) Double-logarithmic I−V plots of the P3HT device after successive oxygen doping procedure with increasing Poxygen from 0.1 mTorr to 750 Torr (each curve labeled with corresponding Poxygen). The dashed lines indicate ohmic regions (slope = 1). Note that the voltage at which I−V curve departs from the slope of one region (i.e., VTh) increases with elevating Poxygen. Along with the increased magnitude of ohmic current, this shows the increased charge carrier concentration. (b) A double-axis graph displaying measured VTh (open circle) in panel a and corresponding nohmic (closed circle) calculated by eq 2 for the given Poxygen applied during oxygen doping. (c) Simulated VTh vs nohmic plots constructed using eq 2 for various device active layer thicknesses d (labeled). For a certain nohmic of interest, d can be chosen rationally such that the corresponding VTh stays in a reasonable voltage range practically accessible during the experiment.
however, showed relatively insignificant change in magnitude as the I−V curves for different Poxygen collapse to nearly one curve (Figure 2a). Since μ⊥ is the only parameter that can affect the magnitude of SCLC in this case, it implicates a weak dependence of μ⊥ on the given ranges of oxygen doping and associated nohmic as we discuss further later. We calculate nohmic for each Poxygen using eq 2 and find that nohmic overall increases ∼20× from 2.9 × 1014 cm−3 to 7 × 1015 cm−3 in response to the 23953
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Figure 3. (a) Measured double-logarithmic μ⊥ vs nohmic data plot. The μ⊥ was obtained from both ohmic (open square) and SCLC (closed square) regions in Figure 2a. The dashed line indicates a linear fit of the data set to eq 3 (slope ≈ 0.2). (b) Schematics describing the hopping conduction model under exponential distribution of localized states. Top left: Energy E vs distance x plot with the short horizontal lines and open circle indicating localized states and a hopping charge carrier, respectively. R is the interspacing between two neighboring hopping sites. Top right: E vs density of states N(E) plot with conduction level energy Ev. Above Ev, the density of localized states follows the exponential distribution, N(E) ∝ exp((Ev − E)/(kTo)), with the characteristic energy width kTo. Bottom: Two neighboring localized hopping sites represented by square energy wells with spacing R. Corresponding electron wave probability function |ψ|2 is shown for each site with α−1 denoting the wave function decay length outside the localized state.
7-decade increase in Poxygen to 750 Torr, as the corresponding VTh reaches ∼3 V (Figure 2b). The relatively small change in nohmic for the given Poxygen increase may suggest that only 0.1% of the solubilized oxygen in the P3HT film (∼6 × 1018 cm−3 at Poxygen = 750 Torr, assuming solubility ∼0.2 (v/v)24) provided free charge carriers, but it is more likely originating from the sandwich device geometry that causes a slow lateral diffusion of oxygen as well as an insufficient light-induced doping activation in the P3HT layer. Though a moderate range of nohmic was probed in our case, the method can characterize a wider range of nohmic by rationally choosing the device active layer thickness. Figure 2c, plotted using eq 2, provides a guideline for a suitable device thickness for measuring certain nohmic using a reasonable bias voltage (Figure 2c). Using the measurement scheme and the controlled doping procedure described above, we obtained the μ⊥ vs nohmic data and found a weak power law dependence of μ⊥ on increasing nohmic, which is consistent with a hopping conduction of charge carriers among localized states under the influence of percolation. In detail, the measured μ⊥ (from both ohmic and space charge limited regions) is in 10−5 cm2 V−1 s−1 range and increases by ∼70% in response to over one-decade increase in nohmic (from 3 × 1014 cm−3 to 7 × 1015 cm−3) as the doublelogarithmic μ⊥−nohmic plot (Figure 3a) displays a linear increase with slope ∼0.2. This is similar to what has been reported by Tanase et al. in P3HT having a comparable range of carrier concentration.26 The observed power law dependence suggests that the carrier transport occurs via percolative variable range hopping (Figure 3b) in which μ⊥ and nohmic have the following relationship:26 μ⊥ =
To / T σo ⎛ (To/T )4 sin(πT /To) ⎞ ⎟ ⎜ × nohmicTo/ T − 1 e⎝ 8Bc α 3 ⎠
localized state (conversely, overlap parameter between localized states). The eq 3, analytically derived by Vissenberg and Matters originally,8 captures the key physics of charge localization, hopping conduction, and percolation, which are highly relevant to the charge transport processes in disordered organic semiconductors, and fitting our μ⊥−nohmic data to the equation yields To ≈ 361 K and α−1 ≈ 0.16 nm. The extracted To value is relatively smaller than the previously reported value in unannealed P3HT (425 K),26 corresponding to a slight reduction of the exponential energy distribution width of the conduction levels by ∼6 meV. Such a reduction in the energy level distribution width has been observed in thermally annealed P3HT and can be explained by an improved polymer chain conformation.28 Unlike To, the obtained α−1 is nearly identical to the previously reported value,26 and this provides insights on the rate-limiting charge hopping process during the out-of-plane charge transport in P3HT. On the basis of the general hopping conduction theory,29 the hopping overlap parameter α−1 describes the decay length of electron wave function outside the potential well of a localized state (Figure 3b). In order for an electron hopping to be realized, the two neighboring states must be close in space and energy, providing the nexus of variable range hopping theory. Using the typical Miller− Abrahams hopping probability expression30 and the exponential distribution that typically describes the localized states in organic semiconductors,8,26 one can readily derive the analytic expression correlating α−1 and the average interspacing R of two neighboring hopping states as31(see the Supporting Information for detailed derivation)
R= (3)
3To −1 α 2T
(4)
Considering that T = 300 K and To ≈ 361 K in our experiment, we then find that R ≈ 1.8α−1 (approximately 2α−1), which indicates that the rate-limiting charge hopping in this case occurs between two nearest neighboring localized states. In fact, the exact R ≈ 2α−1 relationship could be derived previously if the role of percolation during the transport was considered.32 What these relationsships suggest in our examined P3HT is
where σo is the a conductivity prefactor (∼106 S/m for P3HT),26 To is the effective temperature representing energy width of exponential distribution of localized states, Bc is the critical number of bonds per site for onset of percolation (∼2.8 for three-dimensional disordered system),27 and α−1 is a characteristic decay length of the electron wave function on a 23954
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Figure 4. Two-dimensional GIWAXS intensity data (a) and corresponding line plots along qxy (lateral in-plane orientation) and qz (vertical out-ofplane orientation) (b) measured from a Si/AuPd/P3HT/AuPd sample after 3 h of vacuum annealing at 400 K. The inset in panel a shows a schematic of a unit crystalline P3HT lamellar domain, with the line and sheet representing an individual polymer chain and a lamellar sheet, respectively. From peak positions in panel b, we find that the π−π stacked lamellar sheets (in-plane) have a lattice spacing of 0.37 nm, and the lamellar-stacked polymer chain (out-of-plane) 1.6 nm as indicated in the inset of panel a. (c) A schematic describing the charge transport model along the out-of-plane orientation in the semicrystalline P3HT film. The blue lines represent amorphous polymer chains located between crystalline lamellar domains. When an out-of-plane charge transport is forced by external electrical field, the π−π overlapped polymer chains in the amorphous region (indicated by red dashed circles) mediate the charge hopping between vertically apart crystalline domains.
then that R is at least 0.32 nm and compares well with the π−π stacking distance of P3HT chains, ∼0.37 nm, which we quantified from the glazing-incidence wide-angle X-ray scattering (GIWAXS) (Figure 4a,b) and is consistent with previously reported values.33 This strongly implicates that the rate-limiting charge hopping process is most likely to occur between two neighboring P3HT chains through overlapped π electron states (π−π stacking). The notion, however, appears counterintuitive given the lateral in-plane π−π stacking orientation of P3HT (i.e., perpendicular P3HT lamellar stacking) observed from the GIWAXS data (Figure 4a,b) and the vertical out-of-plane charge transport orientation in our device. This can be reconciled by considering the semicrystalline nature of P3HT that has crystalline lamellar domains surrounded by a disordered amorphous region. It is first noted that the strong interchain π−π stacking delocalizes charge carriers within a P3HT lamellar domain as previously confirmed by the charge-induced optical measurement in P3HT FETs10,34 and the photogeneration in P3HT thin films;35 therefore, the rate-limiting charge hopping across the π−π stacked polymer chains is less likely to occur within a lamellae itself but rather expected to happen in certain parts of a disordered region between vertically situated neighboring crystalline P3HT lamellar domains (Figure 4c). This scenario is in line with the model that Kline et al. proposed to explain the enhanced field effect (i.e., in-plane) mobility in high molecular weight P3HT, where overlapped long polymer chains in the disordered boundary region bridge neighboring crystalline P3HT lamellar domains.36 Given the delocalization of charge carriers within the lamellar domains, the identification of π−π stacked polymer chains as the rate-limiting hopping site supports the suggested role of disordered region in mediating the charge transport between crystalline domains and, more importantly, implicates that the same disordered region facilitates the out-of-plane orientation charge transport as well. This discussion potentially leads to the significant notion of having a properly structured disordered region in semicrystalline conjugated polymers for achieving the optimal charge transport and associated perform-
ances in the devices requiring the out-of-plane charge transport orientation (e.g., OPVs).
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CONCLUSIONS In summary, we measured the carrier-concentration-dependent charge mobility along the out-of-plane orientation in the semicrystalline P3HT by using the systematically controlled oxygen doping levels in P3HT and the transition characteristics of charge transport from ohmic to space charge limited conduction. By correlating the analyzed hopping conduction characteristics with the polymer internal molecular packing examined by X-ray scattering, we uncovered the role of the disordered polymer region in mediating the charge hopping between crystalline lamellar nanodomains along the out-ofplane charge transport orientation. The results demonstrate the implication of having suitable polymer structure and molecular weight for optimized disordered polymer region that can lead to the enhanced charge transport and performance in semicrystalline conjugated polymer energy conversion devices.
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METHODS All processing was performed in ambient air unless indicated otherwise. Device and GIWAXS Samples Preparation. First, P3HT powder (American Dye Source) was dissolved in monochlorobenzene and filtered using a 0.45 μm pore-size syringe filter to obtain 3 wt % solution. For a device sample, 30 nm thick AuPd (60/40) was deposited through a shadow mask by thermal evaporation (Lesker PVD 75) at base pressure ∼8 × 10−7 Torr to pattern 1 mm × 7 mm bottom contacts on a 0.7 mm thick glass substrate. For a GIWAXS sample, 30 nm thick AuPd was blanket-deposited on a 0.5 mm thick polished Si wafer. The P3HT solution was then spin-cast at 600 rpm for 45 s (with 5 s of 1500 rpm final drying sequence) on both glass and Si substrates to form a ∼300 nm thick P3HT layer. Finally, the device and the GIWAXS samples were completed by depositing a 30 nm thick top AuPd layer (1 mm × 7 mm pattern for the device sample to form a cross-bar geometry with 1 mm2 active device area and no pattern for the GIWAXS sample). The 23955
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device sample was cleaved after all measurements were completed, and its cross-section was observed in a scanning electron microscope (Hitachi S4800) Device Electrical Measurement. The fabricated device was loaded in a vacuum probe station (Newport) equipped with a sample heater and an oxygen input. The device dedoping was performed following the annealing procedure described in the main text. Device I−V characteristics were obtained using a semiconductor parameter analyzer (Agilent 4156C) at systematically elevated oxygen pressure and doping level as described in the main text. GIWAXS Measurement. The GIWAXS sample prepared above was first dedoped (i.e., annealed) using the identical procedure applied to the device sample. The GIWAXS data were obtained at the X9 undulator-based beamline of the National Synchrotron Light Source at Brookhaven National Laboratory. The slit-collimated incident X-ray beam (λ = 0.0886 nm) was focused onto the P3HT sample, located inside the vacuum chamber (pressure ≈ 0.3 Torr), through the top AuPd layer, using Kirkpatrick−Biaz mirrors (200 μm × 80 μm beam spot size) with corresponding beam footprint on the sample spreading out in proportion to the inverse incident angle. The two-dimensional charged-coupled device (CCD) detector, positioned ∼270 mm away from the center of the sample stage, collected the GIWAXS intensity image. Finally, the data conversion to q-space was accomplished using the reference GIWAXS data collected from silver behenate powder, and the diffraction peaks in qxy and qz line plots were fitted to a Lorentzian profile, Δq/(4(q − qo)2 + Δq2), to determine the lattice constant (2π/qo).
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ASSOCIATED CONTENT
S Supporting Information *
Analytic derivation of relationship between electron wave function decay length and average interspacing of neighboring localized states. This material is available free of charge via the Internet at http://pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This research was carried out at the Center for Functional Nanomaterials, Brookhaven National Laboratory (BNL), which is supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-AC02-98CH10886. I thank Htay Hlaing and Benjamin Ocko at the Condensed Matter Physics and Materials Science Department of BNL for the collection and discussion of GIWAXS data.
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REFERENCES
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