Bottom‐Up Growth of n‐Type Polymer Monolayers for High‐Performance Complementary Integrated Circuits

Downscaling the semiconductor into ultrathin film is of vital importance to high‐performance field–effect transistors (FETs), but the high‐mobility FETs based on conjugated polymer monolayers have been rarely realized. Especially, the lack of high‐performance n‐type polymer monolayer FETs hinders the development of complementary integrated circuits. Herein, by fine‐tuning the supramolecular assembly of two thiazole flanked naphthalene diimide‐based conjugated polymers, the ≈2.5 nm‐thick monolayers with well‐defined fibrillar morphology are grown in a controllable way, where the one‐dimensional solution‐state structures are inherited. The resultant monolayer FETs show the electron mobility up to 0.25 cm2 V−1 s−1, among the record for n‐type polymer monolayer FETs. More importantly, the first demonstration of polymer monolayer complementary integrated circuits is present, and a record‐high inverter gain of 113 is achieved, which is also identical to the best polymer thin‐film inverters.


Bottom-Up Growth of n-Type Polymer Monolayers for High-Performance Complementary Integrated Circuits
Yifu Guo, Mingqun Yang, Junyang Deng, Chenming Ding, Chunhui Duan,* Mengmeng Li,* Ling Li, and Ming Liu DOI: 10.1002/aelm.202201307 but also effectively minimizing the shortchannel effects. [1] This method was first proposed and successfully utilized for Si FETs. [2] However, when the body thickness of Si is reduced below 5 nm, the resultant mobility is significantly degraded by orders of magnitude due to the thickness-fluctuation-induced scattering. [1,3] Recently, the organic FETs have been found theoretically and experimentally to exhibit thickness-independent mobility even at the thickness <5 nm. [4] In other words, the first few monolayers of organic semiconducting layer adjacent to the gate dielectric play predominant roles in charge transport for organic FETs. Therefore, the organic monolayer FETs are ideal platforms for the fundamental investigation of charge transport. [5] On the other hand, due to the direct exposure of conducting channel, organic monolayer FETs is supersensitive to the environment, holding great potentials in applications of chemical/biosensors, [6] photoresponse electronic devices [7] and health monitoring. [8] Note that the bandgap of organic semiconductors is independent of the film thickness compared to their inorganic counterparts. [9] So far, considerable achievements have been made for high-performance monolayer FETs based on conjugated small molecules. For instance, with the assistance of van der Waal interaction, a dioctylbenzothienobenzothiophene (C8-BTBT) molecular crystal was epitaxially grown on graphene and hexagonal boron nitride substrates, leading to a field-effect mobility over 10 cm 2 V −1 s −1 . [10] A nondestructive deposition method of the metal contacts was developed to fabricate 2,9-didecyldinaphtho[2,3-b:2',3'-f ]thieno [3,2-b]thiophene (C10-DNTT) monolayer FETs, and the Ohmic contact between semiconductor and metal was achieved, resulting in high intrinsic mobility of 12.5 cm 2 V −1 s −1 . [11] Furthermore, the centimetersized n-type single-crystal monolayer was also reported on hydroxyl-free polymer surface, where the electron mobility over 1 cm 2 V −1 s −1 was extracted from FETs with the observation of band-like transport. [7a] Additionally, attempts have also been made for complementary inverters on the basis of smallmolecular monolayer transistors, but a relatively low gain of ≈14 was observed at a supply voltage (V DD ) of 60 V. [7b] In comparison to conjugated small molecules, semiconducting polymers seem more suitable to large-area flexible electronics due to their better intrinsically mechanical flexibility, improved solution processability, and lower variation in Downscaling the semiconductor into ultrathin film is of vital importance to high-performance field-effect transistors (FETs), but the high-mobility FETs based on conjugated polymer monolayers have been rarely realized. Especially, the lack of high-performance n-type polymer monolayer FETs hinders the development of complementary integrated circuits. Herein, by fine-tuning the supramolecular assembly of two thiazole flanked naphthalene diimide-based conjugated polymers, the ≈2.5 nm-thick monolayers with welldefined fibrillar morphology are grown in a controllable way, where the onedimensional solution-state structures are inherited. The resultant monolayer FETs show the electron mobility up to 0.25 cm 2 V −1 s −1 , among the record for n-type polymer monolayer FETs. More importantly, the first demonstration of polymer monolayer complementary integrated circuits is present, and a record-high inverter gain of 113 is achieved, which is also identical to the best polymer thin-film inverters.

Introduction
The use of ultrathin-body semiconductor is a simple but versatile strategy for the realization of low-power and highperformance field-effect transistors (FETs), which is capable of not only noticeably enhancing the electrostatic gate control www.advelectronicmat.de device-to-device performance. [12] However, the current polymer monolayer FETs generally show poor transport properties with the mobility lower than 10 −2 cm 2 V −1 s −1 in spite of few exceptions, [9,13] where the main challenges are the effective control of polymer aggregation in solution and subsequent molecular order in solid state. In particular, the high-performance n-type polymer monolayer FETs have been rarely reported, so that CMOS-like integrated circuits have not been realized based on polymer monolayers.
In this contribution, the solution-state supramolecular assembly of two thiazole flanked naphthalene diimide-based conjugated polymers with strong temperature-dependent aggregation behavior is precisely tuned. With the inheritance of onedimensional (1D) solution-state structures, the polymer monolayer with well-defined fibrillar morphology is self-organized. It is found that PTzNDI-T with stronger aggregation in solution exhibits larger coherence length and fibrillar length in solid state, resulting in the electron mobility of 0.25 cm 2 V −1 s −1 for monolayer FETs. Note that such transistor performance is not only identical to thin film FETs but also among the record for n-type polymer monolayer FETs. More importantly, the first demonstration of polymer monolayer complementary inverters is present, with the record-high gain of 113 for polymer monolayer inverters. This circuit performance is achieved at a relatively low V DD (30 V) and almost identical to the best polymer thin-film inverters. These results pave a new pathway for high-performance plastic integrated circuits in a bottom-up approach.

Results and Discussion
Two new n-type conjugated polymers based on thiazole flanked naphthalene diimides (NDI), PTzNDI-T and PTzNDI-2FT [14] (Figure 1a), are used as the model compounds (Figures S1-S3, Supporting Information). The relatively high molecular weights are obtained for both polymers (number average molecular weight of 30-35 kg mol −1 , Figures S4 and S5, Supporting Information), facilitating the intermolecular interaction and consequent self-assembly. [15] Their supramolecular assemblies in solution are first investigated using small-angle neutron scattering (SANS), where the deuterated chlorobenzene (CB-d5) and o-dichlorobenzene (DCB-d4) are used to dissolve poly mers for higher signal-to-noise ratio. As shown in Figure 1b, the collected SANS curves are analyzed by Porod law, [16] where I(q) is the scattering intensity and the Porod exponent n is the fractal dimension of scattering objects. The Porod exponent is close to 1 for all samples, indicating that at room temperature (RT) both polymers are self-organized into a 1D worm-like structure in solution, independent of solvent ( Figure 1b and Table 1). More detailed analysis using 1D model reveals that these two polymers possess almost constant radius in solution (r ≈ 1.8 nm), but PTzNDI-T exhibits doubled Kuhn length (L K , twice of persistence length) than PTzNDI-2FT, implying its longer conjugation length (Table 1 and Figure S6, Supporting Information). Such improved molecular order of PTzNDI-T can be attributed to its more planar backbone, as calculated from density functional theory (DFT, Figures S7 and S8, Supporting Information). Furthermore, it seems that CB leads to slightly larger L K than DCB for both polymers.
UV-vis absorption spectroscopy is another powerful tool to investigate polymer aggregation in solution. At RT, one broad band between 500 to 850 nm is observed for PTzNDI-T in CB with the maximum absorption at 656 nm, corresponding to strong intermolecular interactions between polymer chains (Figure 1c). The use of DCB as solvent results in a small blueshift by 4 nm for this main absorption band, implying slightly reduced polymer packing (Figure 1e). This trend is also consistent with L K from SANS. By contrast, the absorption spectrum shape of PTzNDI-2FT at RT distinctly differ. On one hand, the maximum absorption edge of PTzNDI-2FT is blueshifted by around 76 nm compared to PTzNDI-T, suggesting the increased bandgap by 0.17 eV. On the other hand, the main absorption region that is indicative of polymer aggregation consists of an intense first 0-0 vibronic absorption peak at 606 nm with a less intense 0-1 vibronic progression at 661 nm for polymer solution in CB (Figure 1d), which is a typical feature of J-aggregation. [17] Compared to CB, PTzNDI-2FT in DCB exhibits a decreased absorption ratio of 0-0 to 0-1 peak, implying reduced intermolecular interaction ( Figure 1f).
All four polymer solutions show a clear temperaturedependent aggregation behavior in solution. With the elevation of solution temperature, an order-to-disorder transition [18] is observed, indicating that the aggregated phase is transformed to amorphous phase. Such behavior can also be quantitatively described through nucleation-elongation cooperative model [19] (Figure 1g,h). It is evident from UV-vis absorption that the solution temperature above 70 °C generally leads to the disappearance of the majority of aggregated state. Therefore, the elongation regime is selected from RT to 70 °C, where the degree of aggregation (α agg ) can be estimated by [19] where A max (T) is the absorbance of 0-0 peak at temperature of T (in °C), T e is critical elongation temperature, ∆H e is enthalpy release upon elongation, and R is the gas constant. For both polymers, the critical elongation temperature in CB is higher than that in DCB (Table 1), which further confirms stronger polymer aggregation in CB, well consistent with UV-vis absorption. The strong temperature-dependent aggregation behavior generally offers more feasibilities to fine control the morphology and molecular order in solid state. [20] Furthermore, a proper heating treatment to the solution before casting is capable of inducing disentanglement of polymer chains and subsequently facilitating the reorganization and growth of crystalline domains. [21] In spite of good solubility in CB and DCB at RT, the solutions of PTzNDI-T and PTzNDI-2FT are heated to 70 °C for 1 h and then cooled naturally down to RT. The selection of this heating temperature not only provides sufficient energy for chain disentanglement but also ensures the presence of aggregated states in solution (below T e ).
The structural information including orientational texture and long-range crystalline order within the solid state is gained www.advelectronicmat.de through 2D grazing incidence wide-angle X-ray scattering (GIWAXS) (Figure S9a-d, Supporting Information). It is clear from 1D profiles ( Figure S9e,f, Supporting Information) that high order (h00) reflections are evident in the out-of-plane direction for PTzNDI-T thin film, and no other peaks are visible. This observation indicates that the polymer chains are packed parallel to the surface (edge-on orientation, Figure S9g, Supporting Information), and the corresponding lamellar spacing (d lam ) is around 2.6 nm ( Table 1), independent of deposited solvents. Note that such orientational texture distinctly differs from the state-of-the-art P(NDI2OD-T2) that generally exhibits a face-on texture, in spite of similar chemical structures ( Figure S10, Supporting Information). This difference can be attributed to the much more planar backbone of PTzNDI-T than P(NDI2OD-T2), especially the torsion angle between NDI and adjacent thiazole/thiophene units ( Figures S7 and S11, Supporting Information). PTzNDI-2FT shows similar molecular orientation to PTzNDI-T, but an isotropic (010) ring reflection is observed for thin film cast from CB, which not only corresponds to the π−π stacking distance (d π ) of 0.33 nm (Table 1)

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but also indicates the random orientation of π−π stacking. Further analysis is conducted for (100) peak, and it is found that the related coherence length (L CH ) of PTzNDI-T is larger by ≈2 nm than that of PTzNDI-2FT due to its higher planarity of backbones ( Figures S7 and S8, Supporting Information). Note that this trend in solid state is in excellent agreement with the solution-state supramolecular nanostructure, where PTzNDI-T possesses higher L K in solution ( Table 1).
The single molecular layer (monolayer) of these two polymers is fabricated by drop-casting onto SiO 2 /Si substrate that is functionalized using a hexamethyldisilazane (HMDS) self-assembled monolayer. It is evident from atomic force microscopy (AFM) that the 1D supramolecular nanostructure in solution is inherited in the polymer monolayers, with the generation of fibrillar microstructure (Figure 2). For PTzNDI-T monolayer deposited from CB, the average length of nanofibrils (L ML ) reaches around 426 ± 77 nm (Figure S12, Supporting Information and Table 1). Furthermore, the thickness of each single molecular layer is around 2.5 nm, in excellent agreement with d lam from GIWAXS ( Figure S9, Supporting Information and Table 1). This result proves that (i) the 2.5-nm-thick ultrathin film consists of only one polymer monolayer; (ii) polymer chains are packed in an edge-on way within the monolayer. When deposited from DCB, smaller L ML (270 ± 57 nm) is observed for PTzNDI-T monolayer. In comparison to PTzNDI-T, the fibril dimension of PTzNDI-2FT monolayer is pronouncedly reduced (Figure 2c,d, Figure S12, Supporting Information and Table 1), well consistent with the trends of L CH in thin film and L K in solution. This observation further supports the predominant roles of solution-state supramolecular assembly in morphology of solid state. [22] In short, the temperaturedependent aggregation behavior of conjugated polymers allows the fine control of 1D worm-like nanostructures in solution, consequently resulting in the controllable growth of 1D fibrillar monolayer with edge-on orientation (Figure 2f).
The charge transport properties of polymer monolayers are investigated by fabricating FETs with the top-gate bottom-contact architecture and CYTOP as gate dielectric. Both transfer and output characteristics demonstrate that all polymer monolayer FETs show typical linear/saturation behaviors (Figure 3a-c). For PTzNDI-T monolayer FET from CB, the off-state current is on the order of 10 −12 A, resulting in an on/off current ratio of 10 6 ( Figure 3a). Furthermore, the square root of drain current (I DS ) presents a good linear relation with gate voltage (V GS ), ensuring the accuracy of mobility estimation. [23] The extracted saturated electron mobility is 0.18 ± 0.05 cm 2 V −1 s −1 with the maximum value of 0.25 cm 2 V −1 s −1 . Note that such performance is comparable to the PTzNDI-T thin-film FETs, [14] providing further evidence that the monolayer of semiconductor close to the dielectric interface plays the predominant role in charge transport of planar FETs. [4a] It is also found that larger V DS leads to slightly higher electron mobility ( Figure S13, Supporting Information). When deposited from DCB, I DS of the resultant monolayer FET is slightly lower than that from CB, contributing to the average mobility of 0.15 ± 0.05 cm 2 V −1 s −1 . Identical solvent-dependent performance is also observed for PTzNDI-2FT monolayer FETs (Figure 3b). These trends can be ascribed to the solventdependent fibril length as determined from AFM ( Figure 2). Figure 3d summarizes the electron mobilities of all four types of monolayer FETs. The electron mobility of PTzNDI-2FT monolayer devices is 0.16 ± 0.05 cm 2 V −1 s −1 (maximum value of 0.25 cm 2 V −1 s −1 ) for CB samples and 0.12 ± 0.01 cm 2 V −1 s −1 (maximum value of 0.13 cm 2 V −1 s −1 ) for DCB samples. These performances are slightly lower than that of PTzNDI-T monolayer, which can be attributed to smaller coherence length and fibril length ( Table 1). Note that the fibril dimension of polymer monolayer is remarkably reduced on top of modified electrode ( Figure S14, Supporting Information), which could be one of the main reasons for the relatively large contact resistance of around 100 Ω m ( Figure S15, Supporting Information). The short channel effect is observed when scaling the channel length to a few micrometers ( Figure S16, Supporting Information). Additionally, negligible hysteresis behavior is observed for these top-gate polymer monolayer FETs ( Figure S17, Supporting Information). It has to be emphasized that such high electron mobility are among the mobility record for n-type polymer monolayer FETs, [13,27] which is highly desired for organic CMOS-like circuits.
With the combination of previously reported p-type PffBT4T-2DT polymer monolayers, [9,28] the complementary integrated circuits can be realized using PTzNDI-T and PffBT4T-2DT polymer monolayers (Figure 4a). Figure 4b shows the input-output characteristics of resultant CMOS-like inverter based on these n-type polymer monolayers, where V DD = 30 V is applied and a clear voltage amplification is observed. To the best of our knowledge, this is the first demonstration of complementary integrated circuits fabricated using polymer monolayers. The corresponding gain reaches 113, which is one order of magnitude higher than the polymer monolayer inverter with zero-V GS load architecture. [9] More importantly, such circuit performance is also among the gain record for polymer inverters with various architectures including ambipolar-, CMOS-, and resistor-type designs (Figure 4d). It has to be emphasized that the high inverter gain is typically achieved at very high V DD (≥50 V) for reported polymer thin-film inverters, especially for the devices based on the ambipolar FETs, while the polymer monolayer inverters present in this contribution apply a more reasonable V DD (30 V) (Figure 4d). Additionally, the CMOS-like inverter based on PTzNDI-2FT monolayer also shows good performance with the gain of 80, and only a small hysteresis effect is observed (Figure 4c, Figure S18, Supporting Information).

Conclusions
In conclusion, the aggregation behaviors of two n-type conjugated polymers are precisely tuned, and 1D worm-like supramolecular assembly in solution is revealed. Such controllable aggregation results in the growth of polymer monolayers with well-defined 1D fibril nanostructure. The resultant polymer monolayer FETs exhibit identical performance to thin-film device, and the electron mobility of 0.25 cm 2 V −1 s −1 is achieved, which is also among the record for n-type polymer monolayer FETs. More importantly, the first demonstration of polymer monolayer complementary inverters is realized based on these n-type polymer monolayers. Note that the record-high gain of 113 is obtained for such polymer monolayer complementary inverters, almost identical to the best polymer thin-film inverters but with more reasonable V DD . These results not only provide further understanding of structure-property relation for conjugated polymers but offers new possibilities for imperceptible wearable devices and electronic skins using polymer molecular layer. www.advelectronicmat.de

Experimental Section
All experimental details can be found in Supporting Information.

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.