Polymer-Free Carbon Nanotube Thermoelectrics ... - ACS Publications

Nov 17, 2016 - insulating polymer and control of the level of nanotube bundling in the network, which enables higher thin-film conductivity for a give...
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Polymer-Free Carbon Nanotube Thermoelectrics with Improved Charge Carrier Transport and Power Factor Brenna Norton-Baker,† Rachelle Ihly,† Isaac E. Gould,† Azure D. Avery,‡ Zbyslaw R. Owczarczyk,† Andrew J. Ferguson,*,† and Jeffrey L. Blackburn*,† †

Chemistry & Nanoscience Center, National Renewable Energy Laboratory, 15013 Denver West Parkway, Golden, Colorado 80401, United States ‡ Department of Physics, Metropolitan State University of Denver, 1201 Fifth Street, Denver, Colorado 80204, United States S Supporting Information *

ABSTRACT: Semiconducting single-walled carbon nanotubes (s-SWCNTs) have recently attracted attention for their promise as active components in a variety of optical and electronic applications, including thermoelectricity generation. Here we demonstrate that removing the wrapping polymer from the highly enriched s-SWCNT network leads to substantial improvements in charge carrier transport and thermoelectric power factor. These improvements arise primarily from an increase in charge carrier mobility within the s-SWCNT networks because of removal of the insulating polymer and control of the level of nanotube bundling in the network, which enables higher thin-film conductivity for a given carrier density. Ultimately, these studies demonstrate that highly enriched s-SWCNT thin films, in the complete absence of any accompanying semiconducting polymer, can attain thermoelectric power factors in the range of ∼400 μW m−1 K−2, which is on par with that of some of the best single-component organic thermoelectrics demonstrated to date.

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strategy takes advantage of the strong interaction between the π electron systems of s-SWCNTs and conjugated polymers, such as polyfluorenes.14,15,26 Using this approach, s-SWCNTs can be isolated in high-yield in a single step, resulting in samples with a s-SWCNT purity estimated to be >99% by spectroscopic techniques15 or even >99.98% by statistical analysis of transistor measurements.23,24 However, the dispersing polymer is difficult to remove completely, typically resulting in SWCNT networks with a ca. 1:1 polymer:nanotube ratio,27 and it has been suggested that the residual polymer hampers charge carrier transport through the s-SWCNT network.26 In this work we take advantage of a recent modification of the polyfluorene-mediated s-SWCNT enrichment, whereby the H-bonded supramolecular polymer 1,1′-(((1E,1′E)-(9,9-didodecyl-9H-fluorene-2,7-diyl)bis(ethene-2,1-diyl))bis(6-methyl-4oxo-1,4-dihydropyrimidine-5,2-diyl))bis(3-dodecylurea) (SP) is used to selectively disperse s-SWCNTs.28 After the s-SWCNTs are dispersed by the SP, the SP can be disassembled and removed by disrupting the H-bonds, allowing for the production of samples enriched with s-SWCNTs and containing no residual polymer.28 Using this approach, we

emiconducting single-walled carbon nanotubes (sSWCNTs) have recently gained attention as thermoelectric materials1−3 because of the promising physical properties engendered by their one-dimensional π-conjugated chemical structure. Thermoelectric semiconductor materials are typically evaluated using a dimensionless figure-of-merit, zT = α2σT/κ, where α is the thermopower (Seebeck coefficient), σ the electrical conductivity, T the temperature, and κ the thermal conductivity.4 Optimization of zT is complicated by the interdependence of these variables,4 although organic semiconductors often do not demonstrate the same strong correlations observed for inorganic materials.3,5−7 Additionally, nanostructured semiconductors exhibit structural and electronic properties that are favorable for improvement of zT8,9 due in part to a reduction in the strength of coupling between the individual thermoelectric parameters. Band structure calculations for SWCNTs indicate sharp van Hove singularities, suggesting the potential for a large thermopower,3,10 which has been confirmed for samples of highly enriched s-SWCNTs.1,3 The ability to isolate s-SWCNTs with well-controlled electronic structure11−16 has led to a better understanding of their true potential in a diverse range of energy17−21 or (opto)electronic22−25 applications. Enrichment strategies have focused on methods that avoid the need for covalent nanotube functionalization, and a particularly versatile enrichment © XXXX American Chemical Society

Received: September 6, 2016 Accepted: November 11, 2016

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ACS Energy Letters prepare films of highly enriched s-SWCNTs with and without the dispersing polymer to assess the impact this polymer has on the electrical transport and thermoelectric properties. Importantly, we developed a methodology for SP disassembly different than the solution-phase method developed by Pochorovski et al.28 Our method, which relies upon SP dissolution af ter thin-f ilm deposition, is critical for realizing dramatic improvements in the thermoelectric performance of our s-SWCNT networks. We show that removal of the Hbonded supramolecular polymer af ter thin-f ilm deposition results in (1) enhanced carrier doping and electrical conductivity, (2) enhanced charge carrier mobility, and (3) a factor of ∼2.5 increase in the thermoelectric power factor. The strategy developed here should be directly applicable to improving a wide variety of emerging technologies based on thin films of s-SWCNTs, including thermoelectric devices,3 thin-film solar cells,20,21 and field-effect transistors.24,25 Three samples based on SWCNTs prepared in-house by the laser vaporization (LV) technique form the basis of this study. These are (i) “PFOBPy:LV”, a control sample of LV sSWCNTs dispersed by poly[(9,9-dioctylfluorenyl-2,7-diyl)-altco-(6,6′-{2,2′-bipyridine})] (PFO-BPy) and two samples of LV s-SWCNTs dispersed using the H-bonded SP (ii) “solid SPR” and (iii) “solution SPR”, where the SP is removed by treatment with trifluoroacetic acid (TFA) either after network deposition or in solution prior to network deposition, respectively. Figure 1A displays the ultraviolet/visible/near-infrared (UV/ vis/NIR) absorption spectra, normalized to the area under the second-exciton (S22) optical transition envelope, for the three sSWCNT inks used to prepare the nanotube networks. The similarity between the peak envelopes corresponding to both the first-exciton (S11) and second-exciton (S22) optical transitions indicates that PFO-BPy and the SP result in similar s-SWCNT distributions extracted from the starting LV material, with some subtle differences in the exact yield of the individual s-SWCNT species. In both cases, the lack of peaks in the 600− 850 nm region, characteristic of the M11 optical transitions, suggests that metallic SWCNT contamination is below the optical detection limit (i.e., < 1%). The similar species distributions are confirmed by photoluminescence excitation maps (Figure 1B), which show emission peaks of broadly similar relative intensities, corresponding to s-SWCNTs with similar chiral indices, denoted by the (n, m) indices adjacent to the circles that identify particular emission peaks.29 For the sample where the H-bonded SP is removed in solution and the s-SWCNTs subsequently redispersed in N-methyl-2-pyrrolidone (NMP), there is significant broadening and a discernible bathochromic shift (30−40 nm; 14−19 meV) of the S11 optical transitions. This points toward strong van der Waals interactions between individual s-SWCNTs30 and suggests that nanotube bundles are present even in solution. Note that the peak at ∼1930 nm corresponds to imperfect background subtraction of an absorption due to the NMP solvent. The normalized UV/vis/NIR absorbance spectra of the sSWCNT networks prepared by ultrasonic spray deposition (and subsequent postdeposition solvent treatment to remove excess PFO-BPy or fully remove SP) are shown in Figure 2A. In comparison to the inks, the S11 and S22 optical transition envelopes of the PFO-BPy:LV and solid SPR samples are broadened, indicating stronger interactions between individual s-SWCNTs, although only a small (20−25 nm; 10−11 meV) bathochromic shift of the S11 optical transitions is observed. The network prepared from the NMP dispersion (solution

Figure 1. (A) UV/vis/NIR absorbance spectra, normalized to the area under the S22 absorbance envelope, of LV SWCNT dispersions. (B) Photoluminescence excitation (PLE) maps of LV SWCNT dispersions for fluorene copolymer-wrapped tubes (top) and supramolecular polymer-wrapped tubes (bottom).

Figure 2. (A) Absorbance spectra, normalized to the area under the S22 absorbance envelope, and (B) atomic force micrographs of LV SWCNT networks.

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ACS Energy Letters SPR) also exhibits broadened optical transition envelopes compared to the corresponding ink, but in this case the S11 optical transitions also exhibit a sizable bathochromic shift of ca. 50 nm (ca. 21 meV) compared to the NMP dispersion, suggesting that further bundling occurs during the film deposition or formation steps. Atomic force microscopy (AFM) surface topography images of the three networks (Figure 2B) show that the largest bundles (ca. 43 nm) are present in the solution SPR film, consistent with the large bathochromic shift observed upon film deposition. The AFM images of the PFO-BPy:LV and solid SPR films do not indicate the presence of individual sSWCNTs (either polymer-wrapped or bare) but reveal nanotube bundles on the order of 14 and 24 nm in diameter, respectively. This observation suggests that solvent evaporation during film deposition drives s-SWCNT bundle formation. However, retaining the SP during film formation helps to limit the extent of bundling, resulting in bundles that are substantially smaller than those found when the SP is removed prior to network deposition. Prior to the postdeposition solvent treatments, the morphology of the films prepared from the PFO-BPy:LV and SP:LV dispersions are dominated by the polymer. In both cases, we are unable to measure electrical conductivity in the asdeposited film, consistent with the excess polymer in the film inhibiting charge carrier transport and our previous demonstration of negligible electrical conductivity in films of fluorenebased polymers.3 To illustrate the effects of the polymer removal process in the solid SPR sample, in Figure 3 we show AFM images and absorbance spectra before and after the TFA treatment. The AFM surface topography image of the solid SPR film prior to TFA treatment (Figure 3A, left) exhibits large, amorphous features characteristic of the excess SP in the film. The UV/vis/ NIR absorbance spectrum (Figure 3B) of the as-deposited film exhibits an absorption band at ca. 400 nm that is characteristic of the SP (c.f. the spectrum of the corresponding SP:LV ink used to spray this film, included for reference in Figure 3B). In addition, the individual peaks in the S11 optical transition envelope display only a small bathochromic shift and slight broadening with respect to the SP:LV ink, consistent with the excess SP preventing strong interactions between nanotubes even in the solid state. The Fourier-transform infrared (FTIR) absorption spectrum (Figure 3C) of the solid SPR film prior to TFA treatment exhibits absorption bands due to vibrational modes characteristic of the SP (Figure S1). All three measurement techniques illustrate that treatment of the asdeposited film with TFA results in a dramatic reduction in the polymer content in the film. The characteristic absorption peaks in the UV/vis/NIR (Figure 3B) and FTIR (Figure 3C) spectra appear to be completely absent after TFA treatment. Carbon nanotube bundles are obviously apparent in the AFM image (Figure 3A, right), and the root-mean-square surface roughness, Rq, is reduced from 15.6 to 3.4 nm after TFA treatment. In both of the spectroscopic techniques, the level of polymer is below their detection limit, suggesting nearquantitative SP removal. For instance, for the solid SPR film, there is no evidence of the dominant absorption peaks in the C−H stretching region (2800−3000 cm−1) of Figure 3C. Integration of the peak area for this region suggests that ca. 99.5% of the area is removed because of the TFA treatment. The three samples displayed in Figure 2 exhibit large variations in both the extent of residual wrapping polymer

Figure 3. (A) Atomic force micrographs, (B) UV/vis/NIR absorbance spectra, and (C) Fourier-transform infrared absorbance spectra of the film prepared from the SP:LV ink before and after (solid SPR) the TFA treatment.

within each network and the resulting network morphology. To probe the effects of these changes on carrier transport and thermoelectric performance, the s-SWCNT networks were either (i) partially doped by immersion in solutions of triethyloxonium hexachloroantimonate (OA)31 in dichloroethane (with varying concentration) at 78 °C or (ii) “fully doped” by immersion of the samples in a bath of 3 mg/mL of OA in dichloroethane at 78 °C for at least 10 min, in a method similar to that recently developed within our group to afford fine control over the charge carrier density (doping level).3 The thermoelectric performance of each sample was evaluated by measurement of the thermopower and electrical conductivity as a function of carrier density, which is controlled by the extent of OA doping.3 The thermopower and power factor (PF, α2σ) are plotted versus the electrical conductivity in panels A and B of Figure 4, respectively. The dependence of the power factor on the electrical conductivity (Figure 4B) reveals several important clues regarding the effect of the polymer and nanotube bundling on charge carrier transport. First, when the PFOBPy:LV sample is compared to the solution SPR sample, a similar trend in PF versus σ emerges as well as a similar peak PF (ca. 140−150 μW m−1 K−2). This trend holds until high doping levels, whereupon the two data sets deviate: the solution PR power factor sharply decreases after the peak, predominantly because of the limited peak electrical conductivity in the network. This rapid decline can be attributed to this film’s inability to be doped to high levels efficiently (vide infra). 1214

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correlate the extent of bleaching quantitatively with s-SWCNT carrier density, especially for polydisperse samples with multiple s-SWCNT species. Figure 5A−C shows the qualitative impact of increasing carrier density on the absorbance spectra of the three samples (a more complete series of absorbance spectra is shown in Figure S2). In all cases, low doping levels result in bleaching of the S11 envelope, followed by bleaching of the S22 envelope at increasing carrier densities (Figure S2). In all cases, when the networks are fully doped, the S11 absorption envelope is completely bleached and the S22 absorption envelope is partially bleached. For the PFO-BPy:LV sample, high doping levels result in formation of positive polarons on the PFO-BPy chains (Figure 5A), as evinced by the appearance of a new spectral feature at ∼400 nm.3 In Figure 5D−F, we plot the electrical conductivity of the networks as a function of the fractional bleach (ΔA/A0) of the S11 and S22 absorption envelopes (the procedure used to calculate ΔA/A0 is discussed in further detail in the Supporting Information). For the PFO-BPy:LV network, the maximum conductivity (∼170 000 S m−1) occurs when the S11 and S22 absorption envelopes are bleached approximately 71% (Figure 5D). When the SP is removed after deposition of the network, the dopant is capable of inducing a larger bleach of the absorption bands (∼78%) and the maximum conductivity is more than doubled to ∼375 000 S m−1 (Figure 5E). In contrast, when the SP is removed in solution prior to network deposition, the maximum conductivity reaches only ∼65 000 S m−1 and the S11 and S22 absorption envelopes are bleached approximately 71% (Figure 5F). We rationalize the differences observed in Figure 5 by considering the correlation between ΔA/A0 and the hole density, as well as the necessity for OA molecules to interact directly with the s-SWCNT π-electron system to effectively inject this hole density. Because all of the samples reach their maximum electrical conductivity between 70−80% bleaching of the S11 and S22 absorption envelopes, this suggests that the hole densities are fairly similar when the samples are fully doped. However, the slightly larger extent of bleaching for the solid SPR can be rationalized if one considers that the charge carrier doping occurs by interaction of the dopant molecules with the SWCNT surface. For the PFO-BPy:LV sample, the PFO-BPy polymer partially restricts access of the dopant molecules to the SWCNT surface, whereas the excessive bundling observed for the network prepared from the LV dispersion in NMP also results in a significant reduction in the available SWCNT surface area, despite the absence of a wrapping polymer. These two effects result in a slight reduction in the fractional bleaching level when the PFO-BPy:LV and solution SPR samples reach their maximum doping level. In all cases, the electrical conductivity data in Figure 5D−F exhibit two regimes: at low doping levels the electrical conductivity is weakly dependent on ΔA/A0, followed by a transition to a region where the electrical conductivity rapidly increases with ΔA/A0 up to the maximum conductivity at full doping. We previously demonstrated, for LV SWCNTs with similar nanotube distributions to those employed here, that a ΔA/A0 value of ∼0.5−0.6 corresponds to a shift in the Fermi level of ca. 0.3−0.4 eV, which places the Fermi level close to the onset of the occupied density of electronic states, i.e., “valence band”.3 This suggests that the transition to the regime with the largest slope in the σ versus ΔA/A0 plot occurs when carrier transport is dominated by carriers close to the transport levels

Figure 4. (A) Thermopower and (B) thermoelectric power factor as a function of the electrical conductivity for the PFO-BPy:LV, solid SPR, and solution SPR thin-film networks. The solid lines are generated by a fit of the thermopower to a ln(σ/σmax) dependence7 and are used here as a guide to the eye rather than to infer a particular charge transport mechanism. The error bars are discussed in further detail in the Supporting Information.

Second, and most dramatically, the peak power factor of ∼400 μW m−1 K−2 for the solid SPR film more than doubles that of the previous two films, placing the performance of the ca. 1.3 nm diameter LV networks in the same range as highperformance semiconducting polymers such as PEDOT:PSS. To understand the trends observed in Figure 4, we analyze the absorption spectra of the doped s-SWCNT networks as a function of OA doping level in Figure 5. In our previous study we demonstrated that the extent to which the exciton transitions are quenched is proportional to the hole density injected into the occupied electronic states of the s-SWCNTs.3 This dependence is expected a priori from the phase space filling (PSF) effect for an excitonic semiconductor, where the exciton oscillator strength is reduced proportionally by the presence of charge carriers of either sign because of a reduction of the number of single-particle states that contribute to the bound electron−hole state.32,33 In time-resolved measurements, PSF enables the temporal tracking of carrier populations in SWCNTs, based on the magnitude of the S11 exciton bleach.34−36 In steady-state measurements, this correlation allows for the extent of absorption bleaching to serve as a comparative measure of the relative carrier density within each s-SWCNT network at a given conductivity. This relationship is useful because a number of assumptions must be made to 1215

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Figure 5. (left) Representative normalized absorbance spectra for the (A) PFO-BPy:LV, (B) solid SPR, and (C) solution SPR thin-film networks. (right) Electrical conductivity as a function of the doping-induced bleach of the S11 and S22 peak enveloped for the (D) PFOBPy:LV, (E) solid SPR, and (F) solution SPR thin-film networks.

Table 1. Transport Properties of Doped LV SWCNT Thin-Film Networks slope [S m−1] sample

σmax [S m−1]

(ΔA/A0)maxa

(ΔA/A0)transb

region 1c

region 2d

μFET [cm2 V−1 s−1]

(α2σ)max [μW m−1 K−2]

PFO-BPy:LV solid SPR solution SPR

169 000 376 000 69 000

0.71 0.78 0.71

0.48−0.58 0.43−0.53 0.49−0.61

64 000 112 000 42 000

521 000 787 000 195 000

35.4 55.7 3.5

152 398 139

Value of the fractional bleach (ΔA/A0) at the point of “full doping”. bRange of the fractional bleach (ΔA/A0) at the transition between the low- and high-doping regimes. cSlope of the low-doping regime in the plot of electrical conductivity (σ) versus fractional bleach (ΔA/A0) of the S11 and S22 absorption envelopes. dSlope of the high-doping regime in the plot of electrical conductivity (σ) versus fractional bleach (ΔA/A0) of the S11 and S22 absorption envelopes. a

mobility is estimated using the slope of the source-drain current versus gate voltage.37 The solid SPR sample exhibits the largest slopes (for both regimes), suggesting that removing the wrapping polymer reduces the barriers to carrier transport and results in high carrier mobility. With that said, the solution SPR sample exhibits the smallest slopes, indicating that removal of the wrapping polymer is not sufficient to maintain a large carrier mobility. In this case, the junction resistance between SWCNT bundles has been shown to increase with bundle size,38 suggesting that any potential mobility improvements attained by removing the polymer in the solution SPR sample must be negated by the dramatically larger bundle sizes (Figure

within the occupied density of electronic states. The electrical conductivity (σ) is the product of carrier density (n) and mobility (μ) according to σ = neμ

(1)

where e is the elementary charge. Because ΔA/A0 is proportional to n, the slopes of these two regimes provide an indication of the relative charge carrier mobility in the sSWCNT networks. We note that this dependence is analogous to the general methodology for extracting field-effect mobilities from field-effect transistor (FET) transfer curves, where the gate voltage is proportional to charge carrier density and the 1216

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ACS Energy Letters 2B) in this sample. These observations imply that a network consisting of small bundles of SWCNTs that are free of wrapping polymer is the optimum morphology for highly conductive SWCNT films. To confirm that these observations can be attributed to changes in the charge carrier mobility, we fabricated SWCNT FETs using an identical process as for the films employed in the optical and electrical transport measurements. Figure S5A shows representative transfer curves obtained for the three LV SWCNT samples, and the extracted FET mobilities are included in Table 1. Figure S5B,C illustrate the correlation between the FET mobility, μFET, and the slopes for the two regimes of the σ versus ΔA/A0 plot. These data confirm that the solid SPR sample exhibits the highest charge carrier mobility, whereas the solution SPR sample has the lowest, consistent with the proposed origin for the slopes shown in Figure 5D−F. The trends in Figure 5 allow us to make some informed conclusions on the differences in thermoelectric properties shown in Figure 4. The enhanced performance for the solid SPR sample arises primarily from three effects. First, the electrical conductivity measurements shown in Figure 5D−F demonstrate that residual polymer within these films is essentially an inert filler, although its presence presumably blocks certain areas of tube and bundle surfaces from being directly accessible to dopant adsorption. Thus, removing this polymer generates networks composed solely of the active sSWCNT transport phase and allows for better surface accessibility, and hence doping efficiency, of the OA molecules. Second, removing the polymer in the solid state increases the charge carrier mobility within the network, relative to a control sample containing residual PFO-BPy in a roughly 1:1 mass ratio with the SWCNTs. Third, removing the SP in the solid-state, as opposed to solution-phase removal, enables the retention of small bundle sizes in the final network, similar to (but slightly larger than) the bundle size within the PFOBPy:LV network. The larger bundle size of the solution SPR networks reduces the mobility, consistent with previous studies,38 which in turn limits the ultimate conductivity of these networks. Because the carrier concentration of the fully doped solution SPR network is only slightly smaller than that of the fully doped solid SPR network, the dramatically lower maximum conductivity of the solution SPR sample (64 000 S m−1 versus 375 000 S m−1 for the solid SPR sample) is a direct result of this reduced mobility. Ultimately, the increased surface area associated with the small bundles of the solid SPR network (relative to the solution SPR network), the better accessibility of dopant to SWCNT/bundle surfaces, and the higher carrier mobility (relative to the PFOBPy:LV network) accounts for the ability of the solid SPR network to achieve extremely high electrical conductivities that reach ca. 375 000 S m−1. These results lead to a general conclusion, and some important design principles, for improving s-SWCNT thermoelectric thin films. The primary general conclusion we draw from our data is that improving the charge carrier mobility generates large improvements in the attainable thermoelectric power factor. This conclusion can be easily understood by considering the dependences for the electrical conductivity and thermopower on carrier density. While electrical conductivity is directly proportional to carrier density (eq 1), the Seebeck coefficient is inversely proportional to carrier density and for degenerately doped semiconductors is given by

α=

⎛ π ⎞2/3 m*T ⎜ ⎟ ⎝ 3n ⎠ 3eh

8π 2kB 2 2

(2)

where kB is the Boltzmann constant, e the elementary charge, h Planck’s constant, m* the charge carrier effective mass, and T the absolute temperature.4 Equation 1 indicates that for a higher hole mobility a given conductivity can be reached at lower carrier density. In turn, eq 2 implies that at any given conductivity, the s-SWCNT thin film with the highest hole mobility will also display the highest Seebeck coefficient. Figure 5 provides a simple way to visualize this effect for any series of s-SWCNT thin films, and Table 1 summarizes key metrics from this simple absorption analysis that can be used to quickly screen the effects of a given change on the expected SWCNT TE performance. The shaded regions in Figure 5D−F indicate the doping levels at which the maximum power factor is observed for each sample. The magnitude of the maximum PF for this series of samples is inversely proportional to the ΔA/A0 at which this power factor occurs, implying lower carrier density for better-performing samples. Also, as discussed above, the maximum attainable conductivity, the slope of the conductivity versus ΔA/A0, and the field-effect mobility all correlate directly to the maximum power factor. Ultimately, as demonstrated here for the solid SPR sample, higher carrier mobility essentially shifts the α versus σ curve in Figure 4A up and to the right, translating directly to a dramatic increase in the achievable power factor (Figure 4B). In terms of design principles, our observations suggest that to achieve a highly conductive s-SWCNT network from a polymer-enriched dispersion one needs to limit the level of nanotube bundling in the network and eliminate as much of the insulating polymer as possible. The use of a removable SP for sSWCNT enrichment and its subsequent dissolution af ter thinf ilm deposition produce dramatic improvements to the thermoelectric performance of our s-SWCNT networks. These advances push the performance of the ca. 1.3 nm diameter LV networks into the same range recently demonstrated for both HiPCO s-SWCNTs3 (⟨d⟩ ≈ 1.1 nm, with residual fluorene-based copolymer) and high-performance PEDOTbased organic thermoelectrics.39 While the thermoelectric power factor was recently suggested as an excellent metric by which to make initial judgments regarding the potential performance of new thermoelectric materials,40 it is also important to evaluate the impact of polymer removal on the thermal conductivity of the LV SWCNT networks in the solid SPR sample. Here we employ the micromachined Si−N thermal isolation platform used in our previous study3 to measure the thermal conductance in a solid SPR network prepared identically to those employed in the optical and electrical transport measurements. Figure 6 shows the thermal conductivities extracted using this measurement technique for the solid SPR sample immediately after TFA treatment and after heavy doping using OA, where the electrical conductivity has been increased by 3 orders of magnitude. Figure 6 also shows the thermal conductivity of the same sample after intentional dedoping under vacuum. These data suggest that the thermal conductivity of the solid SPR network lies in the range of 2.45−3.85 W m−1 K−1, close to the values obtained previously for a PFO-BPy:LV network.3 These values allow us to estimate a material zT of ca. 0.031−0.048 at 298 K, similar to the upper values estimated for HiPCO s-SWCNTs in our previous work.3 1217

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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/acsenergylett.6b00417. Experimental methods; FTIR spectrum of the supramolecular polymer for comparison with as-deposited and treated films; description of the sources of errors in the measured thermoelectric properties and their propagation to yield the errors bars shown in Figure 4; description of area determination for calculation of ΔA/A0 and complete series of absorbance spectra of the s-SWCNT networks as a function of doping level; fieldeffect transistor (FET) measurements (PDF)

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Figure 6. Thermal conductivities of a solid SPR film measured using a micromachined Si−N thermal isolation platform. We note that the “undoped” sample in this case is actually lightly doped by the TFA treatment. Here we define “undoped” as the fact that the film was not treated with OA.

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: jeff[email protected]. Notes

The authors declare no competing financial interest.

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Looking forward, it is exciting to envision the application of similar methodologies to s-SWCNTs with diameters in the range of 1.0−1.1 nm, which we recently demonstrated can produce power factors of ca. 350 μW m−1 K−2, even at a polymer:SWCNT mass ratio of ca. 1:1. If power factor enhancements similar to those in the present study for ca. 1.3 nm diameter LV SWCNTs could be achieved for the smaller diameter nanotubes, this would push the performance of doped s-SWCNTs above most organic TE materials demonstrated to date. The recent demonstration of high-yield s-SWCNTs extraction using various highly selective, removable dispersants25,41−45 suggests a viable path forward for such materials. We have used 1.3 nm diameter LV s-SWCNT networks as a test platform to elucidate the impact of residual fluorene-based wrapping polymer on charge carrier transport and thermoelectric power factor. Our study demonstrates that removal of the wrapping polymer increases charge carrier mobility, which in turn substantially increases the maximum conductivity and thermoelectric power factor that can be attained for a given sSWCNT network. Additionally, we demonstrate that the correlation between network conductivity and excitonic bleaching as a function of incremental charge transfer doping can be used as a tool for tracking changes to charge carrier mobility within a series of s-SWCNT networks (with identical population distribution). This technique circumvents the inherent difficulties of accurately determining charge carrier densities for s-SWCNT films by substituting the easily measured fractional absorption bleaching as a proxy for carrier density. Finally, the data shown here encourage the exploration of strategies that produce highly individualized SWCNTs and/ or very small SWCNT bundles within the s-SWCNT network. The strategies demonstrated here should enable immediate improvements in a variety of technologies that require highly conductive, enriched s-SWCNT thin films, including thermoelectric devices, photovoltaic solar cells, and field-effect transistors.

ACKNOWLEDGMENTS The investigation of the thermoelectric properties of the SWCNT networks was performed under a grant from the Laboratory Directed Research and Development (LDRD) Program at the National Renewable Energy Laboratory (NREL). The development of the s-SWCNT separations and synthesis of the supramolecular polymer at NREL was funded by the Solar Photochemistry Program, Division of Chemical Sciences, Geosciences, and Biosciences, Office of Basic Energy Sciences, U.S. Department of Energy (DOE). NREL is supported by the DOE under Contract No. DE-AC3608GO28308. B.N.-B. and I.E.G. received support from the U.S. Department of Energy, Office of Science, Office of Workforce Development for Teachers and Scientists (WDTS) under the Science Undergraduate Laboratory Internships (SULI) Program. We thank Prof. Barry Zink (University of Denver) for insightful discussion regarding the transport measurements. The U.S. Government retains (and the publisher, by accepting the article for publication, acknowledges that the U.S. Government retains) a nonexclusive, paid up, irrevocable, worldwide license to publish or reproduce the published form of this work, or allow others to do so, for U.S. Government purposes.

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REFERENCES

(1) Nakai, Y.; Honda, K.; Yanagi, K.; Kataura, H.; Kato, T.; Yamamoto, T.; Maniwa, Y. Giant Seebeck Coefficient in Semiconducting Single-Wall Carbon Nanotube Film. Appl. Phys. Express 2014, 7, 025103. (2) Piao, M.; Joo, M.-K.; Na, J.; Kim, Y.-J.; Mouis, M.; Ghibaudo, G.; Roth, S.; Kim, W.-Y.; Jang, H.-K.; Kennedy, G. P.; et al. Effect of InterTube Junctions on the Thermoelectric Power of Mono-Dispersed Single Walled Carbon Nanotube Networks. J. Phys. Chem. C 2014, 118, 26454−26461. (3) Avery, A. D.; Zhou, B. H.; Lee, J.; Lee, E.-S.; Miller, E. M.; Ihly, R.; Wesenberg, D.; Mistry, K. S.; Guillot, S. L.; Zink, B. L.; et al. Tailored Semiconducting Carbon Nanotube Networks with Enhanced Thermoelectric Properties. Nat. Energy 2016, 1, 16033. (4) Snyder, G. J.; Toberer, E. S. Complex Thermoelectric Materials. Nat. Mater. 2008, 7, 105−114.

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EXPERIMENTAL METHODS Experimental details are provided in the Supporting Information. 1218

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DOI: 10.1021/acsenergylett.6b00417 ACS Energy Lett. 2016, 1, 1212−1220