CONDUCTING POLYMERS

A nonconjugated radical polymerglass with high electrical conductivityYongho Joo,1* Varad Agarkar,2* Seung Hyun Sung,1

Brett M. Savoie,1† Bryan W. Boudouris1,2†

Solid-state conducting polymers usually have highly conjugated macromolecularbackbones and require intentional doping in order to achieve high electricalconductivities. Conversely, single-component, charge-neutral macromolecules couldbe synthetically simpler and have improved processibility and ambient stability. Weshow that poly(4-glycidyloxy-2,2,6,6-tetramethylpiperidine-1-oxyl), a nonconjugatedradical polymer with a subambient glass transition temperature, underwent rapidsolid-state charge transfer reactions and had an electrical conductivity of up to28 siemens per meter over channel lengths up to 0.6 micrometers. The chargetransport through the radical polymer film was enabled with thermal annealing at80°C, which allowed for the formation of a percolating network of open-shell sites inelectronic communication with one another. The electrical conductivity was not enhancedby intentional doping, and thin films of this material showed high optical transparency.

Conducting polymers have relied on con-jugated macromolecular backbones thatare subsequently chemically doped soas to achieve high electrical conductivityvalues [for example, poly(3,4-ethylene

dioxythiophene) doped with poly(styrene sulfo-nate) (PEDOT:PSS)] (1, 2). Despite their impres-sive electrical conductivity values (3, 4), certainaspects of these macromolecules are not ideal.First, for some applications, optical transpar-ency at visible wavelengths can be difficult toachieve with extended conjugated backbones.Second, the syntheses of advanced conductingpolymers can be quite complicated with lowyields. Third, chemical doping can depend onprocessing and lead to performance variability,and the dopants can decrease the materials anddevice stability.

Charge-neutral macromolecules that achieverelatively high electrical conductivity values with-out doping could address some of these issues (5).Radical polymers (6) with nonconjugated back-bones and stable open-shell pendant groups (7, 8)can pass charge through a series of oxidation-reduction (redox) reactions between the pendantopen-shell sites (9). Because of the high densityof redox-active sites associated with these mate-rials, they have had an effect on myriad energystorage and energy conversion applications (10–14).However, the highest solid-state electrical con-ductivity value reported for radical polymers was~10−2 S m−1 (15).Despite the low reported conductivity values,

the redox reactions that allow for charge ex-change between the pendant groups are rapid(16–18), so high conductivities should be possible

with appropriatemolecular engineering of charge-transporting sites (19). We synthesized poly(4-glycidyloxy-2,2,6,6-tetramethylpiperidine-1-oxyl)(PTEO) using a ring-opening polymerizationmeth-odology, which allowed all of the radical sites inthe monomer to be conserved in the polymer.Given its flexible macromolecular backbone anda near–room temperature glass transition tem-perature (Tg), its flow temperature is well belowthe degradation temperature of the macromole-cule. Thermal annealing of the radical polymerthin film resulted in the formation of percolat-ing networks of radical sites in the solid statethat were in electronic communication with oneanother. This network formation occurred de-spite the amorphous nature of the polymer thinfilm. We achieved a >1000-fold increase in theelectrical conductivity of PTEO relative to otherreport values for radical polymers, and the ulti-mate conductivity of ~20 S m−1 is comparablewith commercially available, chemically dopedconducting polymers.We polymerized a monomer that contained

an open-shell site directly because postpolymeri-zation conversion of closed-shell pendant groupsto open-shell forms rarely achieves complete con-version and can lead to undesired by-products(20, 21). We introduced a nitroxide functionalityby reacting epichlorohydrin with 4-hydroxy-2,2,6,6-tetramethylpiperidine-1-oxyl (TEMPO-OH)in the presence of a base (Fig. 1A). This small mo-lecule, 4-glycidyloxy-2,2,6,6-tetramethylpiperidine-1-oxyl (GTEMPO), was purified (fig. S1) in orderto yield awell-definedmonomeric species (22) thatwas stable across a range of different temperatures

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Joo et al., Science 359, 1391–1395 (2018) 23 March 2018 1 of 4

1Charles D. Davidson School of Chemical Engineering, 480Stadium Mall Drive, Purdue University, West Lafayette, IN47906, USA. 2Department of Chemistry, 560 Oval Drive,Purdue University, West Lafayette, IN 47906, USA.*These authors contributed equally to this work.†Corresponding author. Email: bsavoie@purdue.edu (B.M.S.);boudouris@purdue.edu (B.W.B.)

3400 3450 3500 3550 3600

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.)

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OO

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O

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PTEO

Fig. 1. Controlling and monitoring the amount of open-shell sites inradical polymers is a key objective in the design of high-conductivityopen-shell macromolecular species. (A) The monomer synthesis andring-opening polymerization-based synthetic scheme used to generatethe PTEO radical polymer. (B) The constant radical density (there is anoverlapping nature of the spectra acquired at different temperatures) of thePTEO macromolecule in chloroform solutions as a function of temperature,

as determined with EPR spectroscopy. (Inset) A photograph of therecrystallized monomer. (C) The EPR spectroscopy signal of a radicalpolymer thin film showing a classic Lorentzian shape, which is indicative ofsubstantial radical-radical interactions in the solid state. (D) The ATR-FTIRspectra of the oxoammonium-based TEMPOnium salt, the GTEMPOmonomer, and the PTEO polymer, indicating the absence of the oxoammoniumcation signal in either the monomer or radical polymer used in this work.

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(Fig. 1B). In the ring-opening polymerization pro-tocol, a potassium tert-butoxide initiator generatedthe PTEOmacromolecule with a number-averagemolecular weight of 2.4 kg mol−1 (23). The highdensities of radical sites present in the mono-merswere also present in the polymers (Fig. 1C),within experimental error, as determined withelectron paramagnetic resonance (EPR) spectro-scopy (24, 25). Moreover, no oxoammonium ca-tion sites within either themonomer or polymerproducts, which increase the electrical conduc-tivity of other nitroxide-bearing radical polymers(26), were detected with attenuated total reflec-tance Fourier transform infrared (ATR-FTIR)spectroscopy (Fig. 1D) at the characteristic fre-quency (n) of 1550 cm−1. Identifying these do-pants at the levels that can affect charge transportcan be difficult with standard chemical charac-terization methodologies. Thus, although thesedata suggest that no dopants are present in the~0.1% (by weight) level, we cannot definitelystate that there are no oxoammonium cation(or other dopant) sites present at lower levels.For this relatively low molecular weight ofPTEO, the main nitroxide interactions would bebetween neighboring chains in a thin film. Adensity functional theory (DFT) characterizationof the orientational dependence of TEMPOinteractions revealedmultiple favorable pairingconfigurations (fig. S2), with interaction energies

on the order of –4 kcal mol−1 (–0.17 eV). Theseinteraction energies, an order of magnitudegreater than the thermal energy, would resultin nitroxide separations between 4 and 6 Å(fig. S3).High electrical conductivity requires a rela-

tively low Tg, so that the material can be an-nealed in the molten state and away from thedegradation temperature of the radical pendantgroups. The bulky nitroxide-containing substit-uent of the PTEO prevented crystallization ofthe polymer chains (fig. S4) and resulted in aTg of ~20°C (Fig. 2A), which is well below the~150°C onset of degradation for PTEO (Fig. 2B).Thermal processing below 100°C (Fig. 2C) al-lowed us to create percolating nitroxide net-works for charge transport between electrodes,which has not been seen for other systemsbecause of the high Tg of most radical polymers,and this inability to process radical polymerthin films appropriately has contributed toreports of extremely low electrical conductivityvalues for nitroxide-based radical polymers (27).We modeled radical network formation in

PTEO with Monte Carlo (MC) simulations ofthe annealing process using a Hamiltonianparameterized with DFT calculations and withan implicit treatment of the polymer (supple-mentary materials). The configurations werethen processed to characterize the radical net-

works and their degree of percolation (b) ac-cording to

b ¼ 1

3ðsx þ sy þ szÞ ð1Þ

Here, si is the span of the largest network ineach lattice dimension, with b ranging from 0 (nopercolation) to 1 (complete percolation acrossthe lattice). The simulations revealed a dramaticeffect of annealing on the percolation behavior atall radical densities. Random distributions of theradical groups (an amorphous as-cast film) didnot form percolating networks (Fig. 2D). Onlyafter annealing did discontinuous subnetworks(Fig. 2E) combine through aggregation to formpercolating networks (Fig. 2F). Moreover, onlyabove a critical radical density could percolat-ing networks form at all.We monitored this transition for PTEO from

the low-charge-transport regime (~10−9 S m−1) tothe high-charge-transport regime (~10 S m−1)in real time as thematerial crossed from the as-spin-coated glassy state (the thin film was castat ~20°C) into a liquid-likemolten state (Fig. 3A).In order to evaluate charge transport in thisquenched state, the films were transferred toan inert atmosphere vacuum probe station andheld at a temperature of 100 K. Heating of thesample occurred inside of the inert atmosphere

Joo et al., Science 359, 1391–1395 (2018) 23 March 2018 2 of 4

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TEMPO Density

Fig. 2. PTEO flow, thermal, and percolation properties. (A) The glasstransition temperature at ~20°C for the PTEO macromolecule is clear, andthere is no melting or crystallization transition. (B) The onset degradationtemperature of PTEO is ~150°C. (C) No degradation of PTEO was observedwhen the material was held at 80°C for 2 hours. (D) Percolation behavior as afunction of radical density and annealing summarized from the Monte Carlosimulations.The average values are shown as solid lines, and standard

deviations are shown as the shaded regions. (E and F) Typical configurationsfor the (E) unannealed and (F) annealed films at a radical density of2.5 × 1021 cm−3. Independent networks are drawn with different colors, andany networks composed of less than five molecules are rendered gray andtransparent. For example, in (F) almost all of the nitroxide groups are colorednavy blue, indicating that there is one large continuous percolation networkfor charge transport after annealing has occurred.

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vacuum probe station in order to capture thecharge transport ability at low temperaturesbefore bringing the PTEO thin films near Tg, sothe changes in conductivity appear to be causedby changes in the nanoscale structure as op-posed to changes in chemical oxidation. In theseexperiments, the sample was allowed to reachthe desired temperature and held at that tem-perature for 30 min before collecting the elec-trical data. Then, the temperature of the samplewas moved to the next temperature. Moreover,in a separate experiment, replicate devices wereannealed at 80°C for 2 hours in inert atmo-sphere conditions but without evaluating theirlow-temperature electrical properties (withoutthe “1st heating” scan in Fig. 3A), and the con-ductivity of thin films processed in this mannerbegan and remained at ~10 S m−1 (in the samemanner seen for the “2nd heating” scan of Fig.3A), indicating that orderingwas occurring thoughthermal annealing (fig. S5).Once local order was created within the melt,

it was retained as the thin film was cooled backinto the glassy state, and this enabled rapidcharge transport at room temperature both forcooling (Fig. 3A, red squares) and heating (Fig.3A, blue triangles) of the thin film. AboveTg, thinfilms displayed thermally activated transport withan activation energy of ~90 meV (Fig. 3B), whichis consistent with the increased molecular mo-tion. Moreover, as has been observed with otherradical polymers (22, 26, 28), the transport wasindependent of temperature below Tg (Fig. 3C).To study the role of defects, such as impurities

or traps, we performed a fragility analysis on theannealed MC networks by removing a randomdistribution of the TEMPO sites from the an-nealed MC networks (those shown in Fig. 2, D toF) and recalculating the percolation behavior(Fig. 4A). Defect introduction has adverse effectsat all TEMPO densities but was most sensitivenear a TEMPO density of ~1.7 × 1021 cm−3. Weestimate that the TEMPO number density inPTEO is at least 1.75 × 1021 cm−3, on the basis ofprevious characterizations of chemically similar

polymers (29), so our PTEO films should bestrongly sensitive to the levels of defect incorpo-ration and film processing. Over long enoughlength scales, defects will render the networksnonpercolative. Given the length scale limita-tions of our simulations, we can place a lowerlimit for the percolation length scale of ~20 nmfor radical densities greater than 1.75 × 1021 cm−3.The lack of crystallinity in these amorphous

conductors made it difficult to experimentallydetermine the length scale at which the percola-tion of nitroxide groups ended.We estimated thesize of high-charge-transport domains by chang-ing the lengths between the two electrodes (Fig. 4B).At channel lengths >0.7 mm, the conductivitywas low, ~10−4 S m−1 (21, 26). However, forchannel lengths of 0.6 mm or less, the conductiv-ity reached ~20 S m−1. These data suggest that

the nitroxide percolation network had a charac-teristic length scale of ~600 nm. Thus, if thehigh-charge-transport domains do not bridgeacross the electrodes, the conductivity is limitedby the low-charge-transport domains. Bare chan-nels and those containing polystyrene (PS) werefabricated andhad conductivities below the detec-tion limit (<10−12 S cm−1) of our system (fig. S6).Similar experiments performedwith PEDOT:PSSshowed some contact resistance but otherwiseexhibited high conductivity across this length-scalerange, as expected, and the variation of its higherelectrical conductivity for the same channel lengthswas only twofold (Fig. 4B).We evaluated how the chemical nature of the

thin films affected the observed transport behav-ior. When the radical group within the PTEOrepeat unit was intentionally quenched to form

Joo et al., Science 359, 1391–1395 (2018) 23 March 2018 3 of 4

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Cooling2nd Heating

1st Heating

Fig. 3. PTEO conducts charge after annealing has occurred. (A) Elec-trical conductivity as a function of temperature for PTEO in a 0.5-mmchannel. Real-time annealing of the thin film allowed for local order toappear within the film so that the electrical conductivity was altered by10 decades. (B) Above the glass transition temperature, there is thermally

activated transport, whereas (C) below the glass transition temperature,the electrical conductivity is temperature-independent. The data pointsrepresent the average value measured for four different PTEO films, andthe error bars represent the standard deviation from this average. If noerror bars are present, the error is within the size of the point.

500 nm

Au

Au

Channel

1 10

10-5

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105

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0255075100Defect percentage

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2.50 x 1021 cm-3

TEMPO Density

Con

duct

ivity

(S

m-1

)Fig. 4. The charge transport behavior of radical polymers is dictated by percolating domainsof paired nitroxide groups. (A) Fragility analysis performed on the annealed networks as a functionof radical density and defect density. (B) The electrical conductivity at T = 350 K for PTEO as afunction of channel length (black squares) shows that these localized domains were less than orequal to 600 nm in size in experimental practice. The data points represent the average of themeasurements conducted on four different devices of that channel length, and the error barsrepresent 1 standard deviation from this average. If no error bar is shown, the error is withinthe size of the point. Equivalent experiments were performed with PEDOT:PSS (red squares) servingas the channel material, and there is almost no dependence on the conductivity as a functionof channel length. (Inset) A scanning electron microscopy image of a typical channel used inthese studies, showing the spacing between the gold electrode contacts before film deposition.

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the N-OH functionality (PTEO-OH), the electricalconductivity for the 0.5-mm channel decreased to~10−7 S m−1 (fig. S7A). In order to form an elec-tron donor-acceptor system within the radicalpolymer thin film, we intentionally doped thePTEO film with 4-acetamido2,2,6,6-tetramethyl-1-oxopiperidinium tetrafluoroborate (TEMPOnium)at a high loading (10%, by weight) (26). This in-tentional doping had little impact on the ulti-mate electrical conductivity of the film, and theconductivity showed the same marked increaseafter thermal annealing as was observed for theundoped system (fig. S7B). Moreover, makingcomposite blends of PTEO and PTEO-OH atchannel lengths of 0.5 mm allowed for experi-mental verification of the computational fra-gility analyses shown in Fig. 4A. That is, for thinfilms that conained >90% (by weight) PTEO,the electrical conductivity was ~10 S m−1, as ex-pected; however, when >20% (by weight) of thepolymer composite was composed of inactivePTEO-OH, the conductivity of the thin film fellto ~10−7 S m−1 (fig. S7C). Thus, charge exchangefrom the injecting electrode to the collectingelectrode depends on the ability of the nitroxidegroups to form a percolating structure and noton any inherent limitations of the radical self-exchange reactions.Last, the nonconjugated nature of the radical

polymers led to weak absorption profiles of thematerial both in solution and as thin films. Thatis, the solution absorption spectrum showed theoft-observed signal for nitroxide-based radicalpolymers (fig. S8). However, the ~1-mm-thickPTEO film showed only minimal absorption inthe visible spectrum, with ≥98% transmissionat wavelengths of l ≤ 500 nm and 100% trans-mission at longer wavelengths. These thin films

maintained their high electrical conductivityover multiple weeks (fig. S9) when they wereexposed to ambient conditions. Thus, these rad-ical polymer films present as relatively high-electrical-conductivity materials with high opticaltransparency and ambient stability.

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ACKNOWLEDGMENTS

We thank D. A. Wilcox for assistance in acquiring thethermogravimetric analysis data. Funding: The moleculardesign and characterization work was supported by the NSFthrough the CAREER Award in the Polymers Program (award1554957; program manager, A. Lovinger), and the advancedmicroelectronic device fabrication was supported by theAir Force Office of Scientific Research through the OrganicMaterials Chemistry Program (grant FA9550-15-1-0449;program manager, K. Caster). Authors contributions: V.A.designed and conducted all experiments associated withmolecular synthesis. Y.J. conducted experiments regardingthe experimental optoelectronic properties. S.H.S. designed theinitial electronic characterization experiments and fabricatedthe nanoscale electronic devices. B.M.S. designed andconducted all efforts related to the computational work.B.W.B. conceived the project, designed experiments, anddirected the research. All authors contributed to the writingand editing of the manuscript. Competing interests: Nonedeclared. Data and materials availability: All data needed toevaluate the conclusions in the paper are present in the paper orthe supplementary materials.

SUPPLEMENTARY MATERIALS

www.sciencemag.org/content/359/6382/1391/suppl/DC1Materials and MethodsFigs. S1 to S9References (30–34)

1 September 2017; resubmitted 29 December 2017Accepted 21 February 201810.1126/science.aao7287

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A nonconjugated radical polymer glass with high electrical conductivityYongho Joo, Varad Agarkar, Seung Hyun Sung, Brett M. Savoie and Bryan W. Boudouris

DOI: 10.1126/science.aao7287 (6382), 1391-1395.359Science 

, this issue p. 1391; see also p. 1334Sciencepercolation networks for electrons.Perspective by Lutkenhaus). The polymer has a low glass transition temperature, allowing it to form intermolecular

synthesized a redox-active, nonconjugated radical polymer that exhibited high conductivity (see theet al.to process. Joo the electrons, charge carriers can move freely. These conjugated backbones can also make the polymers rigid and hard

Conducting polymers usually contain backbones with multiple bonds. After chemical doping to remove some ofMoving charges with radicals

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