Introduction

Due to the dramatic advance in complementary metal-oxide-semiconductor (CMOS) image sensors in the past decades, they are widely used in present-day smartphones and digital cameras1,2. However, one of the most critical bottlenecks of CMOS image sensors is the size that is smaller than target objects, making one-to-one direct mapping between a sensor and an object impossible without using lenses. This leads to the limitation of their use for touch-based applications such as interactive display, fingerprint security, and human-robot interaction. In this regard, thin-film phototransistor-based image sensors can be a talented alternative as they can be easily scaled up with lower process costs, provide new functionality such as flexibility, and extend the boundaries of practical applications3,4,5. Therefore, exploring channel materials of thin-film transistors (TFTs) that can bring high carrier mobility, high photoresponsivity, and mechanical flexibility is of great importance to make a new breakthrough for next-generation optoelectronic devices6,7,8.

Two-dimensional (2D) layered nanomaterials, such as transition metal dichalcogenides (TMDs), can be promising contenders for the active material of thin-film phototransistor image sensors9,10,11,12 due to their excellent optical and electrical performances as well as mechanical stability13,14,15. The optical characteristics of TMDs significantly vary with the thickness of the material13,16,17. In general, direct bandgap monolayer TMDs exhibit high photoresponsivity, whereas multilayer TMDs show relatively low photoresponsivity due to their indirect bandgap8,18. However, multilayer TMDs have been recognized as being more suitable for versatile photodetector applications than monolayer TMDs, with great benefits arising from the higher carrier density19 and the wider spectral response from near-infrared to ultraviolet (UV)18, as well as easier fabrication.

Recently, structural engineering of TMDs has drawn significant attention as it can alter the fundamental material properties20. Several approaches to overcome the limitation of the low photoresponsivity of multilayer TMDs have been reported, which include stacking of heterostructures21, decoration with quantum dots22,23, and introduction of different alignment directions24. On the other hand, nano-patterning, a relatively new technique for structural engineering, can provide a novel route to tune the optical and electrical performances of 2D materials by directly modifying the electronic states. For example, graphene, which has a significant limitation for electronic applications due to its zero bandgap, can obtain a finite bandgap through a nanoporous patterning by inducing quantum confinement effect25,26,27,28. In addition, Kim et al.29 reported that nanopore-patterned multilayer molybdenum disulfide (MoS2) presents a large photoluminescence emission peak similar to the direct bandgap monolayer MoS2. However, experimental work has not been reported for the nanoporous MoS2 device and its optoelectronic characteristics.

Here, we present the nano-patterning of multilayer MoS2 to achieve high-performance multilayer MoS2 phototransistors. The block copolymer (BCP), which consists of two homopolymers and forms periodic nanostructures through phase separation30,31,32, was used to pattern multilayer MoS2 into a nanoporous structure containing an array of periodic nanoholes. It should be noted that the photoresponsivity of our nanoporous MoS2 phototransistors (622.2 A W−1) is even higher than the direct bandgap monolayer MoS2 counterparts reported in the literature under similar optical power densities and wavelengths of incident light33,34,35. To investigate a fundamental origin of this ultra-high responsivity from the nanoporous MoS2 phototransistors, we explored the exposed edge structure of the nanoporous MoS2 by Raman spectroscopy, X-ray photoelectron spectroscopy (XPS), and scanning transmission electron microscopy (STEM). Photo-excited charge-collection spectroscopy (PECCS) measurement36 was performed to develop understanding of the newly generated trap states in the bandgap by adopting the nanohole structure on MoS2. Finally, device simulation elucidated the mechanism of photoresponse in nanoporous MoS2 phototransistors.

Results

Nano-scale patterning on multilayer MoS2

A nano-scale template based on the phase separation behavior of the BCP was utilized to fabricate the nanoporous structure on multilayer MoS2 films. Figure 1a shows the fabrication processes to achieve nanoporous MoS2. MoS2 multilayer sheets were separated from bulk MoS2 by a mechanical exfoliation and then located onto a Si/SiO2 substrate. SiO2 with a thickness of 10 nm was deposited on the MoS2 surface by an electron-beam (e-beam) evaporator to lead to covalent bonds with the grafting layer of a random copolymer (RCP) as well as to prevent damage to MoS2 during the etching process26,31,37. The BCP was spin-coated on the RCP film and annealed at 230 °C to produce cylindrical phase separation of polymethyl methacrylate (PMMA) and polystyrene (PS). The cylindrical PMMA phase was selectively removed by shining UV light on the BCP film and acetic acid, resulting in the nanoporous template of PS. A field emission-scanning electron microscopy (FE-SEM) image of the template shows a uniform and periodic structure of circular nanoholes with an average diameter of 21.1 ± 1.1 nm (Supplementary Fig. 1). Moreover, the nanohole dimension could be controlled from 24.7 to 30.8 nm by simply changing the treatment time of oxygen (O2) plasma reactive-ion etching (RIE), as confirmed by FE-SEM images (Supplementary Fig. 2a–d) and the linear relationship between the hole diameter and the treatment time (Supplementary Fig. 2e). Using the BCP as a template, the underlying SiO2 was perforated by sulfur hexafluoride (SF6) plasma RIE. Finally, uniform hexagon nanoholes were formed on the multilayer MoS2 sheets through two consecutive steps: (i) dry etching using boron trichloride (BCl3) plasma RIE and (ii) wet etching using potassium ferricyanide solution. After removing the remaining SiO2 layer, the patterned surface of MoS2 was observed by annular dark field (ADF)-STEM, as shown in Fig. 1b. The uniformity of the nanoholes was validated by calculating the hole diameters in a large-area ADF-STEM image (Supplementary Fig. 3a), in which the average hole diameter was turned out to be 26.2 ± 2.3 nm (Supplementary Fig. 3b). The substructure of the hole distribution was consistently characterized. The hexagonal nanoholes with an incircle radius (rh) of ~12 nm were mostly distributed into hexagonal array with a secondary incircle radius (Rh-h) of ~40 nm (Fig. 1b). This crystallographic feature of the nanoporous MoS2 film was corroborated by the fast Fourier transform (FFT) pattern analysis of the ADF-STEM image (Supplementary Fig. 3a), which is equivalent to electron diffraction of the corresponding region (Fig. 1c). The first diffraction ring shows the overall periodicity of the nanoholes, with a length repetition of ~40 nm, and the outmost faint diffraction ring shows the hexagonal shape of the nanoholes, with rh of ~12 nm. The spotty diffraction pattern indicates that the nanoporous MoS2 has a multi-domain hexagonal array of the hexagonal nanoholes. The ADF intensity profile across several hexagonal nanoholes (Supplementary Fig. 3c) supports this 2D length repetition in the nanoporous MoS2. According to STEM-energy dispersive X-ray spectroscopy (EDX) elemental mapping of Mo K, S K, and Si K peaks on the nanoporous MoS2 transferred to a copper grid (Fig. 1d), it is clarified that the chemical composition of the nanoporous structure is MoS2. The thickness of the nanoporous MoS2 nanosheet was estimated to be ~7.6 nm based on the cross-section ADF-STEM image (the bottom of Fig. 1d), which corresponds to 11 layers of MoS2.

Fig. 1: Preparation and microscopic characterization of the nano-patterned MoS2 multilayer.
figure 1

a A schematic illustration of nano-scale patterning process on multilayer MoS2. b ADF-STEM image showing an array of nanoholes in the multilayer MoS2. c FFT pattern representing a crystallographic hexagonal array of hexagon holes (from Supplementary Fig. 3a). d STEM-EDX elemental mapping (top panels) for Mo K, S K, and Si K peaks (scale bars of 50 nm) and cross-sectional view (bottom panel) of the multilayer nanoporous MoS2.

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Hexagonal nanohole edges in nanoporous MoS2

In order to confirm the material characteristics of the multilayer nanoporous MoS2 regarding the formation of nanoholes, Raman spectroscopy, and XPS were employed. In Fig. 2a, a characteristic Raman doublet of pristine MoS2 is present at 383.7 cm−1 for E12g (in-plane mode) and 407.5 cm−1 for A1g (out-of-plane mode), which corresponds to multilayer MoS238. The peak positions were rarely changed in the nanoporous MoS2, whereas the relative intensity of E12g/A1g was reduced from 0.824 for pristine MoS2 to 0.604 for nanoporous MoS2. This noticeable change is in good agreement with the increased number of edges in nanoporous MoS2 because A1g is preferentially excited for edge-terminated films compared to E12g due to their polarization dependence20,29,39,40. A decrease in the peak intensity of E12g/A1g was consistently observed for nine different nanoporous MoS2 films (inset of Fig. 2a), strongly supporting the reliability of the manufacturing process.

Fig. 2: Edge structure of the hexagonal nanohole in the multilayer nanoporous MoS2.
figure 2

a Raman and b XPS spectra of multilayer pristine and nanoporous MoS2, respectively. c, ADF-STEM image of a typical hexagonal nanohole in the nanoporous MoS2 (inset is SADP exhibiting a typical 2H-MoS2 structure). d–h Atomic-scale ADF-STEM observations of the four regions denoted by I, II, III, and IV in c (scale bars of 1 nm).

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Figure 2b shows XPS spectra of Mo 3d and S 2p core levels of pristine MoS2 and nanoporous MoS2. To compensate for sample charging, all the XPS spectra were calibrated by C 1s peak located at 284.5 eV. Most intense doublets of the Mo 3d and S 2p were observed at 229.15 eV (Mo4+ 3d5/2) and 161.96 eV (S2− 2p3/2), respectively, in both pristine and nanoporous MoS2. However, these two doublets in the nanoporous MoS2 had wider full-width at half-maximum (FWHM) than pristine MoS2. Therefore, we deconvoluted the Mo 3d peak into three component peaks corresponding to intrinsic MoS2 (i-MoS2), substoichiometric MoS2 (s-MoS2), and molybdenum oxide (MoO3) bonding41,42,43,44. The peak of s-MoS2 is attributed to the electronic structure different from stoichiometric MoS2 (ratio of S/Mo = 2) because there are fewer S atoms around Mo atoms at exposed edge sites44. The doublet of s-MoS2 at 228.55 eV (Mo4+ 3d5/2) and 231.68 eV (Mo4+ 3d3/2), respectively, was observed in both the pristine and the nanoporous MoS2, whereas the atomic ratio of s-MoS2 increased from 6.32% in pristine MoS2 to 14.41% in nanoporous MoS2 (Supplementary Table 1). In addition, as evidenced by a decrease of the ratio of S2− 2p to Mo4+ 3d (i-MoS2 + s-MoS2) from 2.00 to 1.89 after nanoporous patterning, the degree of the edge exposure introduced by formation of nanoholes on multilayer MoS2 can be quantitatively estimated.

The MoO3 may be derived from oxidation of surface and etched region of MoS2 when exposed to air. An atomic ratio of Mo6+ 3d (MoO3), located at 232.16 eV (Mo 6+ 3d5/2) and 235.29 eV (Mo6+ 3d3/2), rarely increased from the pristine (10.90%) to the nanoporous MoS2 (12.11%) due to the protective effect of the SiO2 layer on the MoS2 film. In addition, the absence of Cl 2p, B 1s, N 1s peaks (Supplementary Fig. 4a–c) verifies the protective function of the SiO2 layer, which helps to prevent chemical contaminations during the dry etching with BCl3 plasma RIE and wet etching with potassium ferricyanide solution.

To determine the atomic configuration of the hexagonal nanohole edges in the nanoporous MoS2, the atomic structures of multilayer MoS2 showing a stacking sequence as well as the edge of the nanohole should be scrutinized. Figure 2c presents a typical hexagonal shape of the nanohole in the nanoporous MoS2. The inset in Fig. 2c shows the selected area diffraction pattern (SADP) for the MoS2 multilayer region, suggesting a 2H-MoS2 structure (S.G. = P63/mmc) with an AA′ stacking configuration in which atomic columns along the [001] orientation are mixed with Mo and S atoms. To identify the edge structure of the hexagonal hole, atomic-scale ADF-STEM observation of sides and apexes of the hexagonal hole was performed, as shown in Fig. 2d–h. According to atomic configuration of the top MoS2 layer inserted in each figure, the nanohole edges mostly have a zigzag configuration of atomic columns except for a few local disordered regions.

It is worth noting that the nanohole shape was modified when further etching was carried out on the nanohole. The nanohole structure was circular when the MoS2 was dry etched with BCl3 plasma RIE (Supplementary Fig. 5a). When the MoS2 was dipped in potassium ferricyanide solution, the circular-shaped nanoholes became hexagonal structure (Supplementary Fig. 5b, c). The wet etching process changed not only the shape of the holes, but also the atomic configuration of nanohole edges. While MoS2 has two different edge types of zigzag and armchair45,46, the zigzag configuration was found more dominant in the hexagonal nanoholes. Such transformation of edge configuration can be assigned to the higher thermodynamic stability of the zigzag configuration in comparison to that of the armchair configuration47,48. It can be reasoned that the energetically less favorable armchair structure gradually changed to more favorable zigzag structure. An in-depth study of the dependence of the hole shape on the crystallographic orientation and surface energy of MoS2 is still in progress, which is beyond the scope of this study. Given the AA′ stacking configuration of the MoS2 sample, the exposed atomic terminations at every edge are regarded as a mixture of Mo- and S-terminated zigzag structures layer by layer, which can be confirmed by the same intensity profiles of Mo K and S K peaks extracted from edge to edge across the hole (Supplementary Fig. 6c) in the STEM-EDX elemental mapping data (Supplementary Fig. 6a, b).

Electronic structural characterization of nano-patterned MoS2

As opposed to a pristine MoS2, it is expected that the nanoporous MoS2 has a different band structure due to the abundant broken bonds in the MoS2 lattice represented by the exposure of zigzag edges. In order to explore the electric states of the nanoporous MoS2, PECCS measurement was conducted. The PECCS is a useful tool to probe the sub-states inside the bandgap by calculating a photoinduced threshold voltage shift (ΔVTH) from phototransistors36. Figure 3a, b show photo-excited transfer curves of the multilayer nanoporous MoS2 TFTs under monochromatic light illumination with various wavelengths and photon energy dependence of ΔVTH extracted from individual photo-excited transfer curves. In Fig. 3b, it was observed that ΔVTH values of both multilayer devices, the pristine and nanoporous MoS2 TFTs, gradually decreased as the photon energy of the illuminated monochromatic light increased up to the edge of the bandgap energy, and then increased rapidly above the bandgap energy. However, compared to the pristine MoS2 TFT, it can be seen that the VTH of the nanoporous MoS2 TFT shifted significantly, implying remarkable interband charge transitions49. The areal density of states (DOS) of the multilayer MoS2 TFTs can be extracted by distinguishing the excited charge values from the incident photon energy from the ΔVTH dependent curves. Figure 3c presents the experimentally obtained spectral DOS of both multilayer pristine MoS2 and nanoporous MoS2 TFTs. They show the charge transition peaks corresponding to the energy states assigned as T1 (~1.30 eV, indirect bandgap) and T3 (~1.85 eV, direct bandgap)49 matched with the photoluminescence spectra of MoS2. More importantly, additional peaks were observed in the photon energy region lower than the bandgap energy, which correspond to the interband charge transitions. In particular, the experimentally observed interband states in the nanoporous MoS2 were overwhelmingly more than for the pristine MoS2, which seems to have been caused by the abundant exposure of edge atoms in nanoporous structure. Here, note that the T3 transition peaks of the pristine and nanoporous MoS2 were not perfectly matched at the same energy due to the thickness difference between the two MoS2 flakes, rather than the formation of nanoporous structure.

Fig. 3: Optoelectronic characteristics of MoS2 TFTs.
figure 3

a Photo-excited transfer characteristics of the multilayer nanoporous MoS2 TFT under monochromatic light illumination with various wavelengths. b Dependence of ΔVTH on the incident photon energy, which was extracted from individual photo-excited transfer curves of the MoS2 TFTs. c Photon energy-dependent areal DOS profiles of the multilayer pristine and nanoporous MoS2 TFTs. The photoluminescence (PL) spectra (bold gray line) of the nanoporous MoS2 is overlaid on the graph to compare the peak position observed by PECCS and the interband transitions measured by PL.

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Multilayer nanoporous MoS2 phototransistors

In general, in-gap states can directly influence the optoelectronic properties of materials. It has been reported that the in-gap states induced by interstitial atoms in multilayer molybdenum diselenide (MoSe2) TFTs can result in significant photogating effect and high photoresponsivity under illumination49. In this study, we explored the characteristics of the nanoporus MoS2 phototransistors. Figure 4a shows a SEM image of a nanoporous MoS2 TFT, in which source and drain electrodes were positioned on the nanoporous MoS2 with a channel area of 34.28 μm2. The transfer characteristics of the nanoporous MoS2 TFT were measured under the illumination of different incident power densities (Pinc) with light wavelength (λex) of 405 nm and a drain voltage (VDS) of 1 V (Fig. 4b). The same measurement has been carried out for a pristine multilayer MoS2 TFT with a channel area of 56.28 μm2 (Fig. 4c) for comparison. Although a significant increase in OFF-state current under illumination was commonly observed in both devices (due to the conductivity increase by the photoinduced excess carriers; i.e., the photoconductive effect), only the nanoporous MoS2 phototransistor showed a negative VTH shift (i.e., the photogating effect). Consequently, the photocurrent (Iph = Iillumination – Idark) of the nanoporous MoS2 TFT presented a remarkable increase in the ON-state (i.e., at high gate voltages) (Supplementary Fig. 7a), unlike the pristine MoS2 TFT (Supplementary Fig. 7b). We also extracted photoresponsivity (R = Iph/Pinc) of both devices as a function of Pinc (Fig. 4d) (the photoresponsive characteristics of nanoporous MoS2 under illumination of λex = 638 nm are shown in Supplementary Fig. 8). The nanoporous MoS2 photodetector exhibited a photoresponsivity of 622.2 A W−1 (at Pinc = 0.8 mW cm−2 and λex = 405 nm), that is ~1240 times greater than the pristine MoS2 phototransistor. To confirm the reproducibility and reliability of the nanoporous MoS2 phototransistors, we measured the photoresponsivity of five more multilayer nanoporous MoS2 phototransistors with different channel areas. Supplementary Figure 9 shows average photoresponsivity and standard error of six nanoporous MoS2 phototransistors, presenting consistent photoresponsivity with a very small standard error. Our result was compared with several papers that report the enhancement of photoresponsivity of the multilayer MoS2 by applying various methods. Supplementary Figure 10 shows the relationship between photoresponsivity and Pinc extracted from various papers. The trend line (blue dotted line) of photoresponsivity of the nanoporous MoS2 phototransistor is higher than most other devices except one having a heterostructure with direct bandgap monolayer MoS2.The photodetection characteristics were also evaluated by specific detectivity (\({D}^{\ast }=R\sqrt{A/2q{I}_{{{{{{{{{\mathrm{dark}}}}}}}}}}}\)) and sensitivity (S = Iph/Idark), where A is the illuminated area and q is elementary charge. Significant increases in detectivity and sensitivity have been observed for the nanoporous MoS2 phototransistor, as compared with the pristine MoS2 counterpart (Fig. 4e, f). In addition, time-traced photoresponsive characteristic of the nanoporous MoS2 phototransistor was measured under repetitive illumination with λex of 405 nm and Pinc of 25 mW cm−2 (Supplementary Fig. 11). Rising and decay time were approximately 1.01 s and 4.02 s, respectively. 

Fig. 4: Photoresponsive characteristics of the multilayer nanoporous MoS2 TFTs.
figure 4

a SEM images of the nanoporous MoS2 TFT. Transfer curves of b multilayer nanoporous MoS2 and c pristine MoS2 phototransistors in the dark and under illumination of \({\lambda }\)ex of 405 nm for various Pinc. d Calculated photoresponsivity, e specific detectivity, and f Sensitivity of the multilayer nanoporous and pristine MoS2 phototransistors as a function of Pinc.

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To develop further insight into the nanoporous MoS2 phototransistor exhibiting excellent optical properties, quantum transport simulations were performed considering the in-gap states that can trap electrons or holes (see Methods section for details of the simulation approach)50. As demonstrated by the PECCS measurements, the nanoporous MoS2 contained multiple states within the bandgap (Fig. 3c). However, only certain states could be major contributors to the ultra-high photoresponsivity. To test this hypothesis, we simulated drain current (IDS) vs gate voltage (VGS) behaviors of the nanoporous MoS2 phototransistor in the dark and under illumination (Pinc = 417 mW cm−2 and λex = 405 nm), assuming a single trap state (trap concentration, Pt = 5 \({{{{{{{\rm{\times }}}}}}}}\) 1025 m−3) near conduction band (EC) and valence band (EV) in Fig. 5a, b, respectively. A notable observation from the simulated IDS–VGS curves is that, with the in-gap state near EV, a significant ΔVTH can be achieved (Fig. 5b), while this phenomenon did not appear with the trap state near EC (Fig. 5a). When VGS is high, the trap state near EC is fully filled with electrons both in the dark and under illumination, due to the large number of electron carriers in EC. In contrast, while the trap state near EV is fully occupied by electrons in the dark, it can be partially filled by holes under illumination due to the increased hole concentration in the valence band with the generated excess holes, resulting in potential barrier lowering (i.e., the photogating effect). Our simulation results indicated that only the in-gap states near EV can play an important role in the photogating effect, significantly increasing the electron injection from the source. It should be noted that the simulated IDS–VGS characteristics shown in Fig. 5b manifest the same trend as the measured IDS–VGS curves (Fig. 4b), which implies the presence of in-gap states near EV in the nanoporous MoS2 TFT.

Fig. 5: Simulation of the photoresponsive behavior of the nanoporous MoS2 phototransistor.
figure 5

Transfer characteristics of the nanoporous MoS2 phototransistor assuming a single trap state (trap density of Pt =  \(5\times {10}^{25}\,{{{{{{{{\rm{m}}}}}}}}}^{-3}\)) near a EC and b EV in the dark and under illumination (λex = 405 nm, Pinc = 417 mW cm−2). c Iph variation with a single trap state at different Et within the bandgap (Et from 0.2 eV to 1.4 eV above EV) at VGS = 1 V. d Photoresponsivity as a function of Pinc at λex = 405 nm with trap states at four different energy levels (the dashed line is a guide to the eye). Et = 0.3 and 0.4 eV above (below) EV (EC). The total trap density is the same as in Fig. 5a–c (i.e., \(1.25\times {10}^{25}\,{{{{{{{{\rm{m}}}}}}}}}^{-3}\)).

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To examine the impact of trap states at different energy levels (Et), we have calculated Iph by varying Et from 0.2 eV to 1.4 eV above EV in Fig. 5c, which reveals that Iph can be significant only when the trap states are located near EV (Et–EV \(\le\) 0.8 eV). Finally, the photoresponsivity was plotted as a function of Pinc in Fig. 5d. To emulate the realistic material properties of the nanoporous MoS2 including multiple trap states, we have included four trap states at Et = 0.3 and 0.4 eV above EV and below EC with a trap concentration of 1.25 × 1025 m−3 each. The simulated photoresponsivity exhibits the same trend as the experimental results (Fig. 4d), indicating that our model could explain the underlying physics of the experimentally demonstrated nanoporous MoS2 phototransistor.

Conclusion

Nano-scale patterning on multilayer MoS2 was achieved by BCP lithography. A novel phototransistor based on the multilayer nanoporous MoS2 channel exhibited significant enhancement in photoresponsivity, specific detectivity and photosensitivity, as compared with pristine multilayer MoS2 phototransistors. A combination of Raman, XPS, and STEM-EDX measurements confirmed that the nanoporous structure, consisting of hexagonal arrays of hexagon holes, has abundant zigzag edges exposed with Mo- and S-termination. By means of PECCS experiment, it was shown that the nanoporous MoS2 has high density of trap states into the bandgap region, and the device simulation revealed that the trap states existing near EV are particularly crucial for the photogating effect and the overall optoelectronic performance of the nanoporous MoS2 phototransistors. Hence, this work suggests using nano-patterning on multilayer MoS2 for highly sensitive and responsive 2D material phototransistors.

Methods

Preparation of BCP nanotemplate

Sheets of multilayer MoS2 were mechanically exfoliated from bulk, and then positioned on a Si/SiO2 substrate. As a protective layer, SiO2 with a thickness of 10 nm was deposited on the MoS2 by e-beam evaporator. The entire substrate was then spin-coated with 1 wt% toluene solution of poly (styrene–r–methyl methacrylate) (Polymer source, Mn = 8500, Mw/Mn = 1.45) (P(S-r-MMA)) RCP at 3000 rpm. To stabilize the thin film of the RCP, the sample was annealed at 250 °C under vacuum condition for 2 h, followed by washing with toluene. The substrate was then spin-coated with 1 wt% toluene solution of poly (styrene-b-methyl methacrylate) (Polymer source, PS = 55,000, PMMA = 22,000, Mw/Mn = 1.09) (P(S–b–MMA)) BCP at 3000 rpm, and then annealed at 230 °C for 2 h. UV irradiation (VL-6.LC, Vilber lourmat) was performed for 30 min to selectively remove the PMMA from the BCP upon immersion in acetic acid for 20 min.

Fabrication of nanoporous MoS2 TFTs

The BCP nanotemplate was used as a mask to make the nanoporous MoS2. The O2 plasma RIE process (50 W, 10 s, 10 sccm) was performed to etch down to the RCP layer and increased the hole size of the nanotemplate. The SF6 plasma RIE (200 W, 15 s, 10 sccm) was utilized to imprint the nanoholes onto the SiO2 layer. MoS2 was then etched using BCl3 plasma RIE (100 W, 10 sccm) to produce the nanoporous MoS2 by varying the etching time depending on the MoS2 thickness. Then, the MoS2 was immersed in diluted potassium ferricyanide solution for 30 s, resulting in a change in the nanohole shape from circular to hexagonal. Finally, the SiO2 layer was removed by immersing the device in buffered oxide etchant.

To fabricate the nanoporous MoS2 TFTs, the pattern of source and drain electrodes was prepared using photolithography and Ti (20 nm) and Au (100 nm) were deposited. After removing unnecessary portions, the TFTs were annealed under vacuum condition at 200 °C for 2 h to enhance the electrical behaviors.

Characterizations

Raman measurement was carried out using a Micro-Raman spectrometer system (ALPHA300, WITec) with a green excitation laser. A XPS system (K-Alpha, Thermo Fisher Scientific) with monochromated Al Kα radiation was used to investigate the chemical states and stoichiometry of the MoS2 samples. The surface morphologies and atomic structures of the nanoporous MoS2 samples were observed on an aberration-corrected STEM (JEM-ARM200CF, JEOL) in ADF imaging mode at accelerating voltages of 80 and 200 kV, respectively. The angle range of the ADF detector and probe-forming convergence angle were 45–180 and 23 mrad, respectively. Wiener filtering technique inserted in commercial software was used to decrease random noise background in the obtained atomic-scale ADF-STEM images (HREM Filter Pro, HREM research). Elemental mapping of the nanoporous MoS2 was carried out by EDX (JED-2300T, JEOL) with a dual-type Si drift detector operating in the same ADF-STEM imaging mode. Each detector had an effective X-ray sensing area of 100 mm2, providing a total collection efficiency of ~10% of the total generated X-ray signals (4π sr). A semiconductor measurement system (4200-SCS, Keithley) equipped with a dark box was used for the electrical measurements. To measure the optoelectrical characteristics of the devices, laser radiation was perpendicularly incident on the MoS2 channel area (MCLS1, Thorlabs) using single mode fiber optic patch cables (S405-HP for 405 nm and SM600 for 638 nm, Thorlabs).

Optoelectronic characterization

For PECCS measurements, the photo-excited transistor characterizations were carried out using an intense monochromatic light. The photon flux of almost 5 × 1014 cm−2 s−1 was estimated from the optical power density of ~0.1 mW cm−2. The PECCS system consists of a grating monochromator providing a spectral wavelength range from 300 to 1400 nm, a Hg (Xe) light source of 500 W, and an optical fiber that delivers light to the reactive MoS2 surface of the devices. Electrical measurements were performed by electrical analyzer unit (4155C, Agilent Technologies). Photoluminescence measurements were performed by a spectrometer system (ALPHA300, WITec) with a 532 nm excitation laser.

Device simulation

Carrier transport through the nanoporous MoS2 phototransistor was simulated using the non-equilibrium Green’s function (NEGF) formalism within an effective mass approximation (m* = 0.45 m0, where m0 is the free electron mass). For the nominal device, a 50-nm-long, 6.2-Å-thick MoS2 channel and a 2.5-nm-thick SiO2 gate dielectric were used in a bottom-gate Schottky barrier (SB) field-effect transistor (FET) structure with an SB height of 0.1 eV. A power supply voltage of VDD = 0.5 V was applied. Here, we used smaller device dimensions to save simulation time; however, this will not affect the key underlying physics, and the conclusion from the simulation study will remain the same. The transport equation was iteratively solved with Poisson’s equation until a self-consistent solution was achieved. Here, we first performed a ballistic transport simulation assuming ideal MoS2 to identify the underlying physics of the photogating effect (Fig. 5a–c), and then the current degradation due to scattering in the actual nanoporous structure was taken into account by calibrating the simulation results against the experiments (Fig. 5d) using a fitting parameter (degradation factor of 8 × 10−6).

Trap model

A modified Hornbeck–Haynes model of \(\triangle \sigma =q{\mu }_{{{{{{{{\mathrm{n}}}}}}}}}\triangle {n}_{{{{{{{{\rm{PC}}}}}}}}}+q{\mu }_{{{{{{{{\mathrm{p}}}}}}}}}\triangle {p}_{{{{{{{{\rm{PC}}}}}}}}}\) was used to calculate the increase in conductivity regarding to illumination (photoconductive (PC) effect), where μn and μp are the mobilities of electrons and holes, and ∆nPC and ∆pPC are electron and hole excess densities in the conduction and the valence band, respectively. Measured carrier mobility of 2.68 cm2 V−1s−1 was used for both μn and μp. Considering charge neutrality, \(\triangle {n}_{{{{{{{{\rm{PC}}}}}}}}}=g{\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\) and \(\triangle {p}_{{{{{{{{\rm{PC}}}}}}}}}=\triangle {n}_{{{{{{{{\rm{PC}}}}}}}}}+{n}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\) were used with the trap state near EC, while \(\triangle {p}_{{{{{{{{\rm{PC}}}}}}}}}=g{\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\) and \(\triangle {n}_{{{{{{{{\rm{PC}}}}}}}}}=\triangle {p}_{{{{{{{{\rm{PC}}}}}}}}}+{p}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\) with the trap state near EV, assuming that the fermi level is located at mid-gap51. \({n}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\)and \({p}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\) are the trapped electron and hole densities, which can be calculated as \({n}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}=\frac{{P}_{{{{{{{{\rm{t}}}}}}}}}{dg}{\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}}{g{\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}+{P}_{{{{{{{{\rm{t}}}}}}}}}{d\tau }_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{n}}}}}}}}}/{\tau }_{{{{{{{{\rm{g}}}}}}}},{{{{{{{\rm{n}}}}}}}}}}\) and \({p}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}=\frac{{P}_{{{{{{{{\rm{t}}}}}}}}}{dg}{\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}}{g{\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}+{P}_{{{{{{{{\rm{t}}}}}}}}}d{\tau }_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{p}}}}}}}}}/{\tau }_{{{{{{{{\rm{g}}}}}}}},{{{{{{{\rm{p}}}}}}}}}}\), where \({P}_{{{{{{{{\rm{t}}}}}}}}}\) is the total trap concentration51. \(g(=\eta {P}_{{{{{{{{\rm{abs}}}}}}}}}/h\nu )\) is the generation rate of excess carriers, \(d\) is the thickness of the nanoporous MoS2, and \({\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\) is the carrier recombination lifetime for the photoconductive effect. \(\eta\) and \(h\nu\) are internal quantum efficiency and single photon energy, respectively. The value of η was assumed to be 0.5 for λex of 405 nm. The absorbed power density (Pabs) is calculated by \({P}_{{{{{{{{\rm{inc}}}}}}}}}\left(1-\frac{{e}^{-{a}^{\perp }d}+{e}^{-{a}^{\parallel }d}}{2}\right)\), where \({a}^{\perp }\) and \({a}^{\parallel }\) are the absorption coefficients in the vertical (\({a}^{\perp }=18.9\times {10}^{4}{{{{{{{{\rm{cm}}}}}}}}}^{-1}\) at λex = 405 nm)52 and lateral (\({a}^{\parallel }=102.5\times {10}^{4}{{{{{{{{\rm{cm}}}}}}}}}^{-1}\) at λex = 405 nm) directions53, respectively. \({\tau }_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{n}}}}}}}}/{{{{{{{\rm{p}}}}}}}}}\) and \({\tau }_{{{{{{{{\rm{g}}}}}}}},{{{{{{{\rm{n}}}}}}}}/{{{{{{{\rm{p}}}}}}}}}\) are the trapping and detrapping times of electrons/holes, respectively, and the \(\frac{{\tau }_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{n}}}}}}}}}}{{\tau }_{{{{{{{{\rm{g}}}}}}}},{{{{{{{\rm{n}}}}}}}}}}\) and \(\frac{{\tau }_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{p}}}}}}}}}}{{\tau }_{{{{{{{{\rm{g}}}}}}}},{{{{{{{\rm{p}}}}}}}}}}\) ratios can be obtained by \(1/\left(\frac{{P}_{{{{{{{{\rm{t}}}}}}}}}}{{N}_{{{{{{{{\rm{C}}}}}}}}}}{{\exp }}\frac{\triangle {E}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{C}}}}}}}}}}{{kT}}\right)\) and \(1/\left(\frac{{P}_{{{{{{{{\rm{t}}}}}}}}}}{{N}_{{{{{{{{\rm{V}}}}}}}}}}{{\exp }}\frac{\triangle {E}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{V}}}}}}}}}}{{kT}}\right)\)51, where NC and NV are effective density of states at EC and EV, respectively. ∆Et,C and ∆Et,V are defined as EC–Et and Et–EV, respectively. The recombination lifetime (\({\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}=100\) ps) was adopted from a previous report54. Finally, photoconductive current (\({I}_{{{{{{{{{\mathrm{ph}}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}\)) was obtained by \({I}_{{{{{{{{{\mathrm{ph}}}}}}}}},{{{{{{{\rm{PC}}}}}}}}}=\frac{1}{L}{V}_{{{{{{{{\rm{DS}}}}}}}}}\triangle \sigma\). The photogating (PG) effect due to barrier lowering can be simulated by considering trapped holes (i.e., total trap concentration – electrons in the trap state) as \({p}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PG}}}}}}}}}={P}_{{{{{{{{\rm{t}}}}}}}}}(1-f)\), where f is the probability of electron occupation at a given trap state, which can be calculated by \(f=\frac{{v}_{{{{{{{{\rm{th}}}}}}}}}{\sigma }_{{{{{{{{\rm{n}}}}}}}}}{n}_{{{{{{{{\mathrm{C}}}}}}}}}+{e}_{{{{{{{{\rm{p}}}}}}}}}}{{v}_{{{{{{{{\rm{th}}}}}}}}}{\sigma }_{{{{{{{{\rm{n}}}}}}}}}{n}_{{{{{{{{\mathrm{C}}}}}}}}}+{e}_{{{{{{{{\rm{n}}}}}}}}}+{v}_{{{{{{{{\rm{th}}}}}}}}}{\sigma }_{{{{{{{{\rm{p}}}}}}}}}{p}_{{{{{{{{\mathrm{V}}}}}}}}}+{e}_{{{{{{{{\rm{p}}}}}}}}}}\)54. nC and pV are concentration of electrons in the conduction band and concentration of holes in the valence band, respectively, and vth is thermal velocity. en and ep are the probabilities of electron and hole emission from the trap state, respectively, determined by \({e}_{{{{{{{\rm{n}}}}}}}}={v}_{{{{{{{\rm{th}}}}}}}}\,{\sigma }_{{{{{{{\rm{n}}}}}}}}{n}_{0}{{\exp }}\frac{{E}_{{{{{{{{\rm{t}}}}}}}}}-{E}_{{{{{{{{\rm{F}}}}}}}}0}}{{kT}}\) and \({e}_{{{{{{{{\rm{p}}}}}}}}}={v}_{{{{{{{{\rm{th}}}}}}}}}\,{\sigma }_{{{{{{{{\rm{p}}}}}}}}}{p}_{0}{\exp }\frac{{E}_{{{{{{{{\rm{F}}}}}}}}0}-{E}_{{{{{{{{\rm{t}}}}}}}}}}{{kT}}\), where σn and σp are capture cross-sections for electrons and holes, respectively54. n0 and p0 are the concentration of free electrons in the conduction band and the concentration of free holes in the valence band, respectively, under equilibrium, which are defined as \({n}_{0}={N}_{{{{{{{{\mathrm{C}}}}}}}}}{{{{{{{\rm{ln}}}}}}}}\left(1+{{\exp }}\left\{\left({E}_{{{{{{{{\mathrm{F}}}}}}}}0}-{E}_{{{{{{{{\mathrm{C}}}}}}}}}\right)/{kT}\right\}\right)\) and \({p}_{0}={N}_{{{{{{{{\mathrm{V}}}}}}}}}{{{{{{{\rm{ln}}}}}}}}\left(1+{{\exp }}\left\{\left({E}_{{{{{{{{\mathrm{V}}}}}}}}}-{E}_{{{{{{{{\mathrm{F}}}}}}}}0}\right)/{kT}\right\}\right)\), where EF0 is the fermi level under thermal equilibrium assumed to be at mid-gap in this study. For n-type transistors, the electron concentration (nC) is obtained by \({n}_{{{{{{{{\mathrm{C}}}}}}}}}={n}_{{{{{{{{\rm{inj}}}}}}}}}+\triangle {n}_{{{{{{{{\rm{ph}}}}}}}}}\), where ninj is the injected electron concentration obtained from the NEGF simulation. Under illumination, the excess electron concentration becomes a non-zero value, which can be modeled as \(\triangle {n}_{{{{{{{{\rm{ph}}}}}}}}}=g{\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PG}}}}}}}}}/d\), where the recombination lifetime for the photogating effect (\({\tau }_{{{{{{{{\rm{r}}}}}}}},{{{{{{{\rm{PG}}}}}}}}}\)) was assumed to be 1 μs. The hole concentration (pV) is mainly determined by the excess hole concentration (\(\triangle {p}_{{{{{{{{\rm{ph}}}}}}}}}\)), which is the same as \(\triangle {n}_{{{{{{{{\rm{ph}}}}}}}}}\). For a given trap state, f is a strong function of \(\frac{{\sigma }_{{{{{{{{\rm{n}}}}}}}}}}{{\sigma }_{{{{{{{{\rm{p}}}}}}}}}}\), which was assumed to be a logistic function of energy in a logarithmic scale with maximum and minimum values of 103 and 10−3 at EC and EV edges, respectively. Since the trap sites are negatively charged in the presence of electrons, the concentration of trapped electrons (\({n}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PG}}}}}}}}}={P}_{{{{{{{{\rm{t}}}}}}}}}-{p}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PG}}}}}}}}}\)) was considered in Poisson’s equation as follows: \({\nabla }^{2}{{{{{{{\rm{V}}}}}}}}=-\frac{q}{\varepsilon }\big\{{N}_{{{{{{{{\rm{D}}}}}}}}}-{N}_{{{{{{{{\rm{A}}}}}}}}}+{p}_{{{{{{{{\mathrm{V}}}}}}}}}-\big({n}_{{{{{{{{\mathrm{C}}}}}}}}}+{n}_{{{{{{{{\rm{t}}}}}}}},{{{{{{{\rm{PG}}}}}}}}}\big)\big\}\).