Pristine Poly( para-phenylene): Relating Semiconducting Behavior to Kinetics of Precursor Conversion.

We investigated unsubstituted poly( para-phenylene) (PPP), a long-desired prototype of a conjugated polymer semiconductor. PPP was accessed via thermal aromatization of a precursor polymer bearing kinked, solubility-inducing dimethoxycyclohexadienylene moieties. IR spectroscopy and Vis ellipsometry studies revealed that the rate of conversion of the precursor to PPP increases with temperature and decreases with film density, indicating a process with high activation volume. The obtained PPP films were analyzed in thin-film transistors to gain insights into the interplay between the degree of conversion and the resulting p-type semiconducting properties. The semiconducting behavior of PPP was further unambiguously proven through IR and transistor measurements of molybdenum trioxide p-doped films.

Introduction

Structurally perfect poly(para-phenylene) (PPP) has remained elusive ever since its first synthesis was attempted in 1872.[1] Several solution-based direct (e.g., via oxidative coupling of benzene[2] or aryl–aryl coupling of 1,4-dihalobenzene[3]) and precursor routes (e.g., double eliminations of cyclohexenylenes)[4−6] have been developed, which only allowed to probe the properties of impure or imperfect (i.e., cross-linked or exhibiting ortho- or meta-defects) PPP. The high thermal and chemical stability raised expectations that pristine and unsubstituted PPP would be privileged for applications in organic electronics.[7,8] To date, only ill-defined PPP[8] or short oligophenylenes with up to six repeating units were characterized in (opto)electronic devices due to the insolubility of higher homologues.[8−11] Although surface-assisted syntheses complement solution-based approaches, circumventing the issue of solubility and furnishing defect-free PPP due to regioselective couplings,[12−18] the low quantity of product, the high experimental effort, and the lack of follow-up processability hinders further application. We recently developed a synthetic access to the well-defined, soluble, and high-molecular-weight PPP precursor polymer P1 with m = 25 repeat units based on kinked and thus solubility-mediating dimethoxycyclohexadienylene moieties. Additive-free thermal treatment of nanometer-thick films of P1 resulted in aromatization via demethoxylation, forming unsubstituted, insoluble, and highly fluorescent PPP layers (Scheme ).[19] In this contribution, we investigate the electronic and p-type charge transport properties of PPP thin films via photoelectron spectroscopy and analysis of thin-film transistors (TFTs). The performance in TFTs is related to the kinetics of the thermal conversion process, investigated via IR spectroscopy and ellipsometry. Finally, thin films were successfully doped with MoO3, as shown through IR and transistor measurements.

Scheme 1

Thermal treatment of precursor polymer P1 leads to insoluble pristine PPP through demethoxylation and aromatization

Experimental Section

Materials

The synthesis of the PPP precursor polymer P1 was performed according to our previously described procedure.[19] Prior to preparing thin films, the polymer was dissolved in analytical grade chloroform (CHCl3) and stirred on a hotplate at 45 °C for at least 20 min. MoO3 with a purity of 99.99% was purchased from Sigma Aldrich and used without further purification.

Sample Fabrication

For IR spectroscopic measurements, silicon wafers (1.5 × 1.5 cm2, intrinsic, σ > 5000 Ω–1 cm–1) were used with a native oxide layer as the substrate. UV–Vis measurements were performed on quartz glass substrates. Ultraviolet photoelectron spectroscopy (UPS) measurements were carried out on highly doped p-type silicon substrates covered with a native oxide. All substrates were cleaned prior to fabrication via sonication in acetone and isopropanol and dried using a nitrogen gun. To vary the thickness of the organic layer (≈30, 45, 100, and 200 nm), the concentration of the PPP precursor solutions were set to 3.5, 6, 12, and 24 mg mL–1 in chloroform. PPP precursor solutions were spin-cast at 900–4000 rpm for 60 s in a nitrogen glove box with both water content and oxygen concentration below 2 ppm. Subsequently, the backside of the substrate was cleaned with CHCl3. Thermal annealing and conversion were achieved on standard thermal hotplates inside the glove box. Film thicknesses were determined via Vis ellipsometry and a Bruker Dektak XT profilometer.

Device Fabrication

Prepatterned bottom-gate, bottom-contact Si/SiO2 (n-doped/230 nm thermal oxide) substrates with interdigitating Au contacts (30 nm in thickness; W = 10 mm; L = 20, 10, 5, and 2.5 μm) were cleaned, coated, and annealed thermally under identical conditions with respect to the substrates used for the IR studies. Precursor concentrations of 6 and 3.5 mg mL–1 were used with a spinning speed of 900 rpm. Devices were measured inside a nitrogen-filled glove box (water/oxygen concentration < 1 ppm) using a commercial probe card station with a Keithley 4200-SCS parameter analyzer. Transistors of channel length L = 5 μm were used for the conversion study, while a channel length of L = 10 μm was used to study doping.

Mobility Calculation

The mobility μ was extracted in the saturation region based on the Shockley equation using the following formulawhere W and L are the device width and channel lenpan>gth, respectively, Cd is the dielectrics areal capacitance, Id is the drain current, Vg is the gate voltage, and Vd is the drain voltage.

UPS Measurements

The measurements were performed on a Phi VersaProbe II spectrometer with an Omicron HIS 13 n class="Chemical">helium discharge lamp (HeI: hν = 21.22 eV).

Infrared (IR) Spectroscopy

All samples for IR transmission measurements were measured using a Fourier-transform IR spectrometer (Vertex 80v) from Bruker. The complete beam path was evacuated to 3 mbar to prevent absorption from ambient air (water and CO2). A mercury–cadmium–telluride (MCT) detector and a resolution of 4 cm–1 were used for spectra acquisition, and 200 scans were averaged for each spectrum. MoO3 was deposited on a preannealed PPP layer at a rate of 0.2 nm/min (monitored by a quartz crystal microbalance) under ultrahigh vacuum (UHV) conditions, and IR spectra were measured during layer deposition.

Vis Ellipsometry

The samples used for IR spectroscopic measurements were additionally characterized with ellipsometry in the visible range (Vis ellipsometry) to determine the layer thickness and refractive index of the investigated layers. For this purpose, ellipsometric measurements with an incident angle of 70° were performed before each annealing step. The Ψ- and Δ-spectra were fitted in the range without absorption between 470 and 1050 nm by using a simple Cauchy-layer approach with a constant refractive index for the organic material.

DFT Calculations

Density functional theory (DFT) calculations were carried out using the software package Gaussian 09W.[20] We used the B3LYP functional and the 6-311G(d,p) basis set for geometry optimization and vibrational frequency calculation. For the polymeric precursor as well as for PPP, an isolated monomer unit (equivalent to three phenyl rings after aromatization) end-capped with tert-butylphenyl groups was simulated in harmonic approximation (see the Supporting Information for the structure). Dielectric functions have been simulated for each material using the calculated oscillator parameters with the software package SCOUT.[21]

Results and Discussion

Conversion Kinetics

First, we studied the time and temperature dependence of the aromatization step in this additive-free precursor route toward PPP by thin-film IR spectroscopy and Vis ellipsometry. In Figure a–c, relative transmission IR spectra of a 43 nm-thick PPP precursor layer on silicon before and after stepwise annealing at 300 °C under nitrogen atmosphere are compared to DFT-calculated relative transmission spectra of the precursor and PPP. An assignment of the strongest absorption bands to calculated vibrational modes and fitted models for the dielectric functions of the precursor and PPP are given in Figures S4 and S5 (Supporting Information). The deviations for the C–H stretching vibrations around 3000 cm–1 can be explained by the overestimation of the tert-butyl end-groups in the DFT calculations using a shortened model system (mDFT = 1 instead of mexp = 25). Importantly, the IR spectra clearly reveal a fast, complete conversion to PPP after a total annealing time of 40 min.

Figure 1

(a) DFT-calculated relative transmission spectrum of a P1 layer on silicon. (b) Experimental relative transmission spectra of 43 nm P1 on silicon as cast and for varying annealing times at 300 °C. (c) DFT-calculated relative transmission spectrum of a PPP layer on silicon.

The changes in the IR spectra with increasing annealing times reflect the gradual conversion of the polymeric precursor into PPP. By using the above-mentioned dielectric models and assuming linear superposition of vibration absorption, each spectrum for intermediate annealing steps can be fitted in a layer-stack model, accounting for the vibrational modes of the precursor and PPP (see the Supporting Information for details). This approach quantifies the fraction of PPP with a relative error of about 5%. The temperature dependence was studied by a comparison of the relative ratios at 300 and 250 °C at the respective annealing times. The influence of the layer thickness was investigated by applying different precursor concentrations of 6, 12, and 24 mg mL–1 for spin casting at 900 rpm, resulting in layer thicknesses of roughly 45, 100, and 200 nm, respectively. The estimated relative fractions of PPP are plotted in Figure a. Thermal aromatization is quantitative after 40 min at 300 °C; at 250 °C, it reaches >80% completion after 60 min, sufficient for percolative charge transport in TFT devices (vide infra). This incomplete aromatization/linearization of the polymeric backbone is interpreted as a consequence of the steric hindrance for planarization of the polymer strands in thin films.

Figure 2

(a) Fraction of PPP for varying annealing times and temperatures and for different precursor concentrations as determined via IR spectroscopy. The PPP fraction corresponds to the proportion of the PPP thickness in the total thickness of the model. The layer stack is depicted in the inset. (b) Layer thickness d (top) and refractive index n (bottom) of ca. 100 nm-thick precursor layers (precursor concentration: 12 mg mL–1) for varying annealing times and different annealing temperatures (grayscales: 300 °C; redscales: 250 °C) determined through ellipsometry. (c) Normalized change in layer thickness Δdnorm. (top) and refractive index Δnnorm. (bottom) of ca. 100 nm-thick precursor layers (precursor concentration: 12 mg mL–1) with different annealing temperatures (gray: 300 °C; red: 250 °C) plotted against the fraction of PPP given in (a). Δdnorm. and Δnnorm. were derived by subtraction and subsequent division of each value by the value of the untreated precursor layer (t = 0 min).

IR spectroscopy is highly sensitive to monitor the conversion even in ultrathin layers, but because of the small vibrational absorption cross sections, this method requires extended measurement time for a high signal-to-noise ratio. UV–Vis ellipsometry is a fast analytical tool to examine the conversion of P1 to PPP. Thus, layers with varying thickness were investigated by ellipsometric measurements with respect to the annealing times and temperatures. The comparison to IR spectra assigns the ellipsometric results to molecular changes.

Figure b shows exemplarily the results obtained for a 100 nm-thick layer (top: layer thickness d; bottom: refractive index n). Graphs for thinner and thicker layers are given in Figures S10 and S11 (Supporting Information). In accordance to the gradual change in the relative fraction of PPP determined via IR spectroscopy, a gradual decrease in the layer thickness and a gradual increase in the refractive indices were observed. Since thermogravimetric analysis shows that PPP is thermally stable, mass loss during the annealing process (apart from the methoxy loss) as a cause for the decrease in layer thickness was excluded. We identified two trends:

The convern class="Chemical">sion is faster for higher annealing temperatures.

For a given annealing temperature, thicker layers convert faster than the thinner layer.

The first observation is understood in terms of thermal activation of the conversion process, but the second point is counterintuitive. The opposite behavior or at least a constant conversion rate would have been expected, accounting for the slower heat transfer in thicker films. We assume a changed film morphology or film density of thicker layers. Scanning electron microscopy images (see Figure S7), AFM images, and root-mean-square roughness (RMS) values (see Figures S8 and S9) of precursor layers with varying thicknesses show an increasing amount of voids for thicker layers compared to thinner layers. The voids likely originate from residual solvent inclusions during film formation. To exclude concentration effects on the film morphology and density, samples under increasing spin-coating speed (900, 2000, and 4000 rpm) were fabricated from all three concentrations. The refractive index decreases from thin to thick precursor layers between 1.78 and 1.63 (see Figure S12). In light of the simple Cauchy-layer approach used for analysis of the UV–Vis ellipsometric results and observed voids, the measured refractive index can be understood as an effective layer property used to estimate the effective layer density. A lower effective density of thicker layers facilitates the conversion process. During aromatization, the kinked cyclohexadienylene moieties convert into rigid linear units: This spatial rearrangement is likely facilitated within less dense films since more free space for molecular motion is available (similar to the concept of an activation volume).

Figure c shows the normalized change in layer thickness Δdnorm. and refractive index Δnnorm. plotted against the fraction of PPP determined by IR spectroscopy. Since both plots show linear dependency, they can be used as a metric to track the degree of conversion. The corresponding graphs for thinner and thicker layers are given in the Supporting Information and show the same trend.

Electronic Characterization

Electronic characterization of insoluble and infusible polymers, for example, unsubstituted PPP, is difficult because of its lack of processability, which is circumvented by our precursor route. As pristine thin PPP films are accessible for the first time, we further investigated the electrical properties of neat PPP and precursor polymer films. The optical absorption spectrum shows a bathochromic shift of 82 nm after annealing (see Table and Figure S1a,b). This corresponds to a change in the bandgap energy from 3.95 to 2.90 eV, as extrapolated from a Tauc plot (see Figure S1c). Ionization potentials (IPs) of P1 and fully converted PPP were studied via UPS (see Figure S2 for spectra). Due to signs of charging, UPS spectra of P1 could not be analyzed in detail, symptomatic for insulating materials. The lack of conjugation and the large bandgap suggest insulating properties of the precursor in a TFT setup (see Figure ). Conversion to PPP results in an IP of 5.80 eV, close to values obtained for substituted PPPs (typically 5.8 to 6 eV).[22] Although a high injection barrier to conventional electrode materials (e.g., Pt, Au, or monolayer-modified electrodes) is anticipated, PPP should be feasible as a hole transport material.

Table 1

Optical and Electrical Properties of P1 and PPP

compound	λmax,abs. (nm)	Egap (eV)	IP via UPS (eV)
P1	268	3.95
PPP	350	2.90	5.80

Figure 3

Layout of bottom-gate, bottom-contact transistors employed. To study the effect of doping (vide infra), a thin MoO3 layer was evaporated on the top (not shown).

Layout of bottom-gate, bottom-contact trann class="Chemical">sistors employed. To study the effect of doping (vide infra), a thin n class="Chemical">MoO3 layer was evaporated on the top (not shown).

To demonstrate the semiconducting properties of PPP and to understand the interplay between the conversion process and electrical properties of the resulting films, PPP and its precursor were investigated in p-type TFTs. The high temperatures necessary for precursor conversion limit the choice of device layouts and dielectric materials. Considering the required high annealing temperatures, a simple bottom-contact, bottom-gate structure (see Figure ) employing gold as an electrode material due to its high work function[23] and a 230 nm-thick layer of thermally grown SiO2 dioxide as a dielectric material gave the most robust devices. Although SiO2–semiconductor interfaces adversely influence device characteristics through deep traps,[24,25] SiO2 exhibits low gate leakage and high thermal stability. The measured transfer characteristics of a PPP device obtained by spin coating of a P1 solution (6 mg mL–1) and curing for 1 h at 300 °C to ensure complete conversion (see Figure ). Initial transistor measurements exhibit a strong shift in turn-on voltage on the back sweep from high to low voltages, attributed to the filling of surface traps at the PPP/SiO2 interface (see the Supporting Information).[24] After three complete transfer IV sweeps, the turn-on voltage of the device has nearly reached a stable operation point, greatly decreasing the observed hysteresis. The trap states are slowly depopulated under ambient conditions inside a glove box over the course of several days so that a repeated measurement shows a hysteresis comparable to freshly prepared devices. To minimize the effect of traps, in further analyses, only the third transfer curve of pristine devices is evaluated. The calculated saturation region mobilities are in the order of 10–7–10–6 cm2 V–1 s–1 (>10 devices) with peak mobilities of 5 × 10–6 cm2 V–1 s–1. In contrast to PPP, the unconverted precursor P1 expectedly does not show any sign of charge transport in a reference transistor.

Figure 4

Third measured transfer characteristic of an exemplary PPP transistor (L = 5 μm, cured for 1 h at 300 °C): Drain current (Id) as a function of gate voltage (Vg).

Third measured transfer characteristic of an exemplary n class="Chemical">PPP trann class="Chemical">sistor (L = 5 μm, cured for 1 h at 300 °C): Drain current (Id) as a function of gate voltage (Vg).

Previous investigations of unsubstituted PPP regarding its charge transport were unavailable for comparison since the synthetic access of structural well-defined PPP was developed just recently. Unsubstituted para-phenylene oligomers (n = 4, 5, and 6) gave high hole mobilities in transistors (10–2–10–1 cm2 V–1 s–1).[26] We attribute the discrepancy to the herein obtained mobilities to the fact that these devices were processing via vacuum sublimation, which is known to result in higher order and thus a higher mobility. Recent computational results on cyclic oligophenylenes estimated that mobilities up to 2 cm2 V–1 s–1 can be reached.[27] Note that neither the precursor nor the final PPP exhibited n-type semiconducting properties in our devices. Electron transport and injection into PPP are very unlikely for two reasons: First, the electron affinity of 2.90 eV (IP-Egap) suggests electron transport to be highly impeded by the ubiquitously occurring trap states located at 3.6 eV.[28] Second, charge injection into such a high-lying LUMO requires very low work function electrodes known to be very unstable due to degradation at the required high temperatures needed for conversion.

Angular-dependent IR measurements did not indicate a preferred orientation of the PPP polymer strands, neither in as-processed nor tempered films (see Figure S13). Likewise, grazing-incidence wide-angle X-ray scattering (GIWAXS) and 1D XRD measurements of converted and unconverted films do not show any molecular order (see Figures S15 and S16). In combination with the SEM and AFM pictures, we thus conclude that the obtained PPP films are amorphous on the herein investigated scales. Apart from a few notable examples, a high molecular order is the prerequisite for efficient charge transport in transistors.[29−31] Thus, the herein observed modest performance in transistors is, in part, explicable by the absence of molecular order in the films. Prealignment of the precursor as a way to relieve this issue was investigated through the use of advanced processing techniques (see the Supporting Information for details). We were unable to obtain any signs of thin-film order, which we attribute to the chemical structure of the precursor. We further expect that thermal conversion of ordered precursor films would regardless introduce a high degree of disorder due to the completely different molecular shape of P1 (coiled structure) and PPP (rigid rod), impeding efficient charge transport. In conclusion, we demonstrated the conversion of the insulator P1 into a semiconducting material whose charge transport capabilities are inherently limited through its conversion. As such, its mobility does not meet those of high-performance materials optimized for efficient charge transport. We plan to address this issue by minimizing conformational change throughout the conversion process by redesigning our precursor polymer (vide infra).

Influence of Conversion on Charge Transport

We follow the transition from an insulator to a semiconductor by studying the conversion process in devices with a film thickness of 45 nm of P1 annealed at 300 °C for different periods of time. The thinnest layers examined with IR were chosen for two main reasons: (a) we expect films with a higher density and homogeneity to show better charge transport in transistors, and (b) the slower conversion allows us to investigate samples with a lower amount of converted PPP. An overview of the results of 49 working transistors fabricated in three batches with annealing times ranging between 20 s and 60 min can be found in Figure . As long as a lower limit of 2 min of annealing is surpassed, which corresponds to a conversion of ≈80% (vide supra), the maximum current, turn-on voltage, and maximum transconductance are nearly unvaried. Prolonged annealing does not influence the device parameters thereafter since it does not increase the molecular order in the thin films.[30] Even for extended annealing times up to 120 min, the transistors demonstrate a stable performance. As no significant order was achieved (vide supra), the degree of conversion consequently has only a minor impact on device performance, as long as most precursor polymer strands are aromatized to form sufficient percolation pathways. Above the percolation threshold, the inclusion of unconverted PPP is highly unlikely to cause an electronic impact due to its inertness and high IP. In some cases, the inclusion of insulators in a semiconductor improves charge transport.[32,33] Thus, the obtained PPP possesses a high tolerance for an unconverted material. The compatibility with respect to prolonged thermal treatment underlines its high thermal stability as device parameters are not detrimentally influenced. Future research will be directed toward molecular engineering an anti-configured precursor to increase the order in the precursor polymer and, subsequently, the final PPP.

Figure 5

(a) Turn-on voltage of the fabricated transistors as a function of annealing time. (b) Transconductance for transistors annealed for different time periods. (c) Maximum current through the same transistors as a function of annealing time.

Doping

The possibility to alter the charge carrier transport properties through controlled doping is one of the hallmark characteristics of semiconductors. We utilize doping experiments to further underline the semiconducting nature of the obtained PPP films. The chosen dielectric (SiO2), the nonideal film morphology, and the assumed high injection barrier contribute to the observed high turn-on voltages. Successful p-doping of devices should significantly decrease this quantity while increasing the off-current of the devices. Furthermore, doping of organic materials results in the formation of polaron signatures, of which can be detected in the IR, offering an independent way to prove the charge transfer as an indicator for successful semiconductor doping. Thus, PPP thin films doped with MoO3 were first investigated with IR spectroscopy. The metal oxide was thermally deposited on the top of a 53 nm-thick, fully converted PPP layer, and IR spectra were measured continuously during the evaporation process. In Figure , the relative transmission spectra for increasing MoO3 coverage up to 2.0 nm, referenced to the spectrum of the bare PPP layer, are shown. With increasing MoO3 coverage, the intensity of the characteristic broad vibrational modes of the metal oxide below 1000 cm–1 increases continuously. IR-activated vibrational (IRAV) modes between 1000 and 1600 cm–1 and a broad polaronic absorption band around 4500 cm–1 appear. These features saturate for a coverage of around 1.6 nm of MoO3. As MoO3 does not exhibit any absorption features above 1000 cm–1, these intense absorption lines indicate the interaction of the metal oxide with the PPP. Due to the high work function of MoO3 (6.7 eV),[34] holes transfer from MoO3 to PPP, which gives rise to the observed IRAV modes and IR polaron. Similar observations were made for doped organic semiconductors and the sequential deposition of both MoO3 onto layers of CBP (4,4′-bis(N-carbazolyl)-1,1′-biphenyl) and small molecular dopants onto the polymer P3HT.[35−38] Therefore, a p-type doping of PPP with MoO3 seems reasonable, and an enhanced electrical conductivity upon doping is expected (vide infra).

Figure 6

Relative transmission spectra of 53 nm PPP for various MoO3 coverages. The spectrum of the pristine PPP layer was used as a reference. Spectral changes due to charge transfer from PPP to MoO3 are marked in gray. Vibrational modes below 1000 cm–1 can be attributed to MoO3 (light red).

A set of transistors (L = 10 μm) with a 30 nm fully converted PPP layer (120 min, 300 °C) were fabricated to study doping. The layer thickness of the following evaporated MoO3 layer was adjusted to 4 nm to ensure saturated doping (vide infra). Devices with a MoO3 layer exhibit several signs of successful doping, whereas reference devices with only a layer of MoO3 (without the polymer) show no transistor current (comparison in Figure ). The herein used transistors rely on the accumulation of charge carriers at the semiconductor/insulator interface to conduct a current. Meijer et al. demonstrated that strong doping of the semiconductor in such devices induces another conduction channel in the bulk of the semiconductor.[39] This additional channel can be depleted by an increasing gate voltage, decreasing its conductivity. Our transistors follow this model closely, pointing toward a gate voltage-dependent conduction composed of both an organic field-effect transistor channel conduction and a bulk conduction. The induced doping decreased the turn-on voltage of the TFT and reduced the shift of the turn-on voltage for repeating measurements, while the saturation regime mobility was unchanged. Yet, this device improvement comes on the expense of the channel resistance in the off-state. Although doping of the layers is not expected to be homogeneous at the interface, successful doping of PPP was demonstrated.

Figure 7

An exemplary transfer curve of an undoped and doped transistor. The dotted line shows an extrapolation of the saturation regime of the doped transistor according to the Shockley model. Doping induces a strong increase in off-current and a shift in turn-on voltage. Furthermore, a bulk conduction channel opens (which can be depleted, vide infra).

To investigate the stability of n class="Chemical">PPP, a set of doped and undoped devices were stored for 6 months in the glove box and for an additional year under ambienpan>t conditions in the dark. The majority of devices yielded idenpan>tical characteristics compared to freshly prepared ones. Data of an exemplary transfer measuremenpan>t are found in the Supporting Information. Thus, we conclude that n class="Chemical">PPP has a high shelf-life stability.

Conclusions

In summary, we have investigated the aromatization process of the precursor polymer P1 to pristine unsubstituted poly(para-phenylene) and its electrical properties. Based on in-depth IR and ellipsometric studies in the visible range, the conversion kinetics could be traced in time for differently prepared films and for different temperatures. The transition of thicker, less dense films is faster compared to thinner, denser films. This behavior suggests a space-demanding process, which is also reflected in its temperature dependence, since the reaction proceeds faster at elevated temperature.

The change of the refractive indices obtained from ellipsometry turned out to be an alternative and even faster metric to track the conversion, which has a strong impact on the electric properties of the thin film. Whereas the untreated precursor polymer exhibits no sign of charge transport, PPP films show clear p-type semiconducting properties in thin-film transistors. Due to the precursor route, we demonstrate that unsubstituted PPP possesses semiconducting properties. The rate of conversion of the precursor polymer has only a weak influence on device performance even for conversions exceeding ≈80%. PPP films were successfully p-doped using MoO3. Polaron and IRAV absorption bands in IR indicated successful doping, which finally reduced the turn-on voltage in respective TFTs. Future effort will be directed toward improving the order of both the precursor and resulting PPP thin films and to understand the aromatization process to develop material with improved charge carrier transport behavior. We will focus on the synthesis of alignable linear precursor polymers and on a sidechain engineering of the precursor polymer to lower the conversion temperatures.