van der Waals 2D transition metal dichalcogenide/organic hybridized heterostructures: recent breakthroughs and emerging prospects of the device

Abstract

The near-atomic thickness and organic molecular systems, including organic semiconductors and polymer-enabled hybrid heterostructures, of two-dimensional transition metal dichalcogenides (2D-TMDs) can modulate their optoelectronic and transport properties outstandingly. In this review, the current understanding and mechanism of the most recent and significant breakthrough of novel interlayer exciton emission and its modulation by harnessing the band energy alignment between TMDs and organic semiconductors in a TMD/organic (TMDO) hybrid heterostructure are demonstrated. The review encompasses up-to-date device demonstrations, including field-effect transistors, detectors, phototransistors, and photo-switchable superlattices. An exploration of distinct traits in 2D-TMDs and organic semiconductors delves into the applications of TMDO hybrid heterostructures. This review provides insights into the synthesis of 2D-TMDs and organic layers, covering fabrication techniques and challenges. Band bending and charge transfer via band energy alignment are explored from both structural and molecular orbital perspectives. The progress in emission modulation, including charge transfer, energy transfer, doping, defect healing, and phase engineering, is presented. The recent advancements in 2D-TMDO-based optoelectronic synaptic devices, including various 2D-TMDs and organic materials for neuromorphic applications are discussed. The section assesses their compatibility for synaptic devices, revisits the operating principles, and highlights the recent device demonstrations. Existing challenges and potential solutions are discussed. Finally, the review concludes by outlining the current challenges that span from synthesis intricacies to device applications, and by offering an outlook on the evolving field of emerging TMDO heterostructures.

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Sk Md Obaidulla

Sk Md Obaidulla is a Postdoctoral Scientist in the Surfaces, Interfaces & 2D Materials research group at the Institute of Physics, Croatia led by Dr Sc. Marko Kralj. He earned his PhD from the Indian Institute of Technology Guwahati in 2017 under the supervision of Prof. P K Giri. Currently, he specializes in 2D magnetic materials and transport properties at S N Bose National Centre for Basic Sciences (SNBNCBS), India. He has extensive experience on device fabrication and 2D TMD material–based optoelectronic properties. His research focuses on synthesizing novel 2D magnetic materials and fabricating field-effect devices for optoelectronic and spintronic applications.

Antonio Supina

Antonio Supina is a PhD student in the research group of Dr Aleksandar Matković in the Chair of Physics, Department Physics, Mechanics and Electrical Engineering, Montanuniversität Leoben. His research interests lie in the domain of 2D materials, with focus on the CVD synthesis of 2D TMDs as well as their characterization using SPM and optical microscopy. During his PhD, he worked on different syntheses of various novel 2D materials as well as on advanced automatic characterization and processing methods using machine and deep learning.

Sherif Kamal

Sherif Kamal is a PhD candidate in the research group of Dr Marko Kralj of the Institute of Physics in Zagreb. He researches 2D materials with the focus on epitaxial growth of atomically thick borophenes, CVD synthesis of 2D TMDs and their characterization using low-energy electron diffraction (LEED) and photoemission spectroscopic techniques such as XPS and ARPES.

Yahya Khan

Yahya Khan received his PhD from the College of Information Science and Electronic Engineering (ISEE), Zhejiang University, Hangzhou, China. Currently, he is working at the Department of Physics, Karakoram International University (KIU), Gilgit, Pakistan. His research focuses on the synthesis of 2D magnetic materials and their application in spintronic devices.

Marko Kralj

Marko Kralj is a Senior Researcher and group leader of the Surfaces, Interfaces & 2D Materials research group at the Institute of Physics. He also leads the research unit Graphene and 2D Materials within the scientific Center of Excellence for Advanced Materials and Sensing Devices. He received his PhD degree in 2003 from the University of Zagreb in Croatia. His current research focuses on the synthesis, atomic-scale characterization, and applications of various novel 2D materials.

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1. Introduction

Two-dimensional (2D) layered transition metal dichalcogenides (TMDs) have opened up new avenues due to the emergence of structural, optical, and electronic variations that allow the creation of hybrid heterostructure systems with a wide variety of optoelectronic properties. The heterostructure system comprises van der Waals (vdW) layered transition metal dichalcogenides and an organic semiconductor. Therefore, there are various ways to design semiconductor devices and achieve various band offset barriers. The vdW heterostructures are formed by vertically stacking different layers of compounds on top of each other. The very weak interlayer interactions provide complete freedom in selecting materials to form hetero stacks. This leads to separate material layers, resulting in “ideally” clean interfaces. In other words, there will be no interdiffusion of atoms at the interface, given that the transfer is conducted at room temperature.¹ Therefore, tailoring the properties of the materials in hybrid heterostructure systems via interlayer interaction either with organic semiconducting materials or by doping opens a new path for material and device engineering. Furthermore, incorporating organic semiconductors with 2D-TMDs will enable the inexpensive scaling up of fundamental research in optical and electronic applications to an industrial level.

1.1. Crystal structure and bonding in TMDs

Layered TMDs have the basic formula of MX₂, where M can be a transition metal (Mo, Cr, W, V, Ni, Mn, Fe, Sc, Nb, Ta, Ti, or Sn), and X represents a chalcogen (S, Se, or Te). There are more than 50 different layered TMD compounds (see Fig. 1).² These materials have covalent bonds within their layers and weak van der Waals forces holding the layers together, making them easy to split apart. Each layer consists of three atomic planes: one with metal atoms arranged in a hexagonal pattern, sandwiched between two planes of chalcogen atoms. Sometimes, the metal ion arrangement is distorted, causing non-planar layers, as in ReS₂, for example.³

Fig. 1 (a) A library of more than 50 distinct layered 2D TMD layered compounds. These compounds involve the transition metal (M) and the three chalcogens: sulfur (S), selenium (Se), and tellurium (Te), which are prominently featured in the periodic table. (b) Top and side views of the schematic of monolayer MX₂ (MoS₂) with hexagonal (1H) and tetragonal (1T) symmetries.

Transition metal atoms contribute four electrons to bond with chalcogen atoms, resulting in oxidation states of +4 for M and −2 for X. Chalcogen atoms’ lone-pair electrons stabilize the layer surfaces, preventing reactions with the environment. The distance between M–M bonds varies depending on the metal and chalcogen size. An MX₂ monolayer can exist in three phases: 1T, 2H, and 3R, representing different layer counts and symmetries (Fig. 1b). The numbers of X–M–X (S–Mo–S) units represents the number of layers in the unit cell and the letters indicate symmetry (T: trigonal; H: hexagonal; and R: rhombohedral).⁴

1.2. Diversity of TMDS

The MX₂ family exhibits diverse electronic properties, ranging from insulators such as HfS₂ to semiconductors such as MoS₂ and WS₂ and even metals/semimetals such as VSe₂, PdSe₂, TiSe₂, NbS₂, CrS₂, and TaS₂.^5–9 However, this review places an emphasis on TMD semiconductors with bandgaps around 1–2 eV, with a particular importance on MoS₂, WS₂, MoSe₂, and WSe₂. MoS₂ has three structural phases: 1T (metallic), 2H (semiconducting), and 3R (semiconducting). External stimuli or chemical treatments can induce structural phase transitions, making MoS₂ a versatile material for different applications. For instance, when immersed in n-butyl lithium, MoS₂ has the capability to transition from its semiconducting 2H phase to the metallic 1T phase.²

Recently, the investigation of heterojunctions involving atomically thin 2D materials, specifically transition metal dichalcogenides and organic molecules, has emerged as a rapidly expanding field of scientific and technological exploration. The primary objective behind studying these TMDO heterojunctions is to comprehend and harness the distinctive characteristics of their constituent materials, as well as the novel properties that arise from their synergistic combination. When envisioning the attributes of interfaces between dissimilar materials including TMDO heterojunctions, the initial approach often involves simplifying the system into a band energy level diagram. In a TMDO heterojunction, the relevant energy levels encompass the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO) relative to the TMD material valence band maximum (VBM) and conduction band minimum (CBM), respectively. By examining this alignment of energy levels, one can anticipate the flow of charges in both ground and excited states. The substantial hybridization of p- and d-orbitals in the TMDO system near the Fermi energy profoundly influences band alignment and charge transfer at the interfaces, thereby enabling the fine-tuning of the electronic structure and photoluminescence (PL) characteristics of TMDs.^10,11 Consequently, this methodology is frequently employed to identify the potential combinations of TMDO materials suitable for applications such as photovoltaic cells, photodetectors, diodes, and various electronic and optical devices.

However, the model of energy level alignment does not fully encompass many of the pertinent details associated with the TMDO heterojunction. For instance, the proximity of a molecule to the TMD material can induce shifts in the energies of bands and orbitals via the electrostatic and hybridization effects. Moreover, apart from the primary HOMO and LUMO, other molecular orbitals have the potential to contribute to the photophysical dynamics observed in the TMDO heterojunction. Additionally, in the heterojunction, electrons and holes may not exist as free carriers but instead primarily manifest as bound electron–hole pairs, resulting in a range of distinct physical phenomena. The simplistic energy-level alignment approach also fails to account for details pertaining to magnetism or the molecular spin states, molecular orientation, thin-film morphology, and chirality. These extra molecular characteristics provide significant opportunities for further modulation of the properties of TMDO heterojunctions.¹² As a result, the combination of transition metal dichalcogenides and organic semiconductors in heterostructures provides a versatile platform for the exploration of novel optical and electronic properties in multifunctional devices. This is due to the occurrence of unique physics at the interfaces of TMDO. Furthermore, various TMDO heterostructures showcase improved optical and electrical performance when compared to individual TMD layers or organic entities.^13,14 From an application perspective, both 2D-TMDs and organic semiconductors are in high demand for the development of next-generation flexible and wearable electronics, chemical/biological sensors, detectors, displays, and related technologies. The combination of the efficient light absorption capabilities of organics with the high carrier mobility of inorganics holds the potential to enable cost-effective and efficient optoelectronic devices. Unfortunately, the optoelectronic properties of TMDO heterojunctions and TMDs doped with organic molecules have not yet been fully exploited and the underlying physics is still not well understood. Studying heterojunctions and the effects of doping can play a crucial role in determining the function of hybrid systems. Thus, incorporating 2D-TMDs with organic semiconductors would lead to a plethora of opportunities to modulate the optoelectronic properties and different unusual physical and chemical properties that have not been realized before. In addition, the atomic-scale control of organic molecular arrangements on recently reported TMDO interfaces has been found not well controlled. This limitation is noteworthy since the optoelectronic characteristics of organic molecular crystals depend significantly on their aggregation and molecular configurations.^14,15 This review provides an overview of the ultrathin and well-controlled molecular arrangements on the interface through various synthesis techniques and explores how the molecular aggregations would affect the light–matter interactions in the hybrid heterostructure system. Typically, very thin organic molecular layers with a thickness less than 5 nm have a clean interface and high excitonic densities, accompanied by high binding energies. These distinctive properties enable them to effectively interact with excitonic states in TMD materials. As a result, TMDO heterostructures hold great potential as viable candidates for the development of optical devices including photodetectors and light-emitting devices.¹⁶

In addition to that, organic molecules have compatible features with TMDs. The technologies based on many novel applications are challenging because of demands such as large-area coverage, optical transparency, flexibility, and easy processing. Here, combining organic semiconductors with TMDs can fulfill the desired requirements, as organic semiconductors can be deposited by simple fabrication at low temperatures on almost any substrate. Furthermore, this offers a vast array of functions, from memory switching, wide range absorption (visible to far-infrared), light emission, to chemical/biological light-sensors, and photodetectors. Organic molecules have large charge carrier effective masses, and low dielectric polarizability compared to traditional inorganic semiconductors. Therefore, TMDO materials would be established with improved or unusual novel physical as well as optoelectronic properties. TMDs sometimes show intrinsic volatile-chalcogen (S, Se, and Te) vacancies; however, by doping the TMDO heterostructure, the vacancy sites can be reduced, and the optoelectronic properties can be remarkably modulated. It can be noted that despite the vdW interaction, the vdW interfaces are extremely complex and poorly defined in cases when they are buried by the over layers, making it hard to quantify the active junction sites and the physical properties such as those in field-effect transistors (FETs). To resolve this issue, we can use atomically thin TMDs to make TMDO heterostructures, which eliminate all unwanted vdW junctions. Organic layers alone often show amorphous grains in the active channel, and as a result, the transport properties are always realized from hopping and not from band transport.¹⁷

1.3. Types of organic compounds incorporated with TMDs and their requirements

Combining organic layers with 2D TMDs is an area of active research with various applications including electronic devices, detectors, and sensors. The choice of organic layer depends on the specific application and the desired properties. Here are some examples of organic molecules that can be combined with 2D TMDs and their requirements: (i) organic semiconductors: organic semiconductors can be combined with semiconducting 2D TMDs to create heterostructures. The choice of organic semiconductors should match the electronic properties and band energy alignment (type-I/II/III) of the TMD. For example, interlayer exciton (ILX) emission, a type-II band energy alignment with the p–n junction, is necessary, where the organic semiconductor could be p-type and the TMD be n-type (vice versa). Similarly, this p–n junction and ILX emission concept is also applicable for other TMD/organic semiconductors.^13,18,19 Again, for lasing, giant PL emission enhancement type I is needed.^14,20 Moreover, these heterostructures find application in electronic and optoelectronic devices including field-effect transistors (FETs) and photodetectors. (ii) Organic ligands: organic ligands are commonly used to functionalize the surface of 2D TMDs. These ligands can be small molecules or polymers with functional groups (e.g., amines, thiols, and carboxylic acids) that can chemically bind to the TMD surface.^21,22 Organic ligands can improve the dispersibility of TMDs in solvents, enhance stability, and provide specific chemical functionalities for targeted applications. (iii) Polymer matrices: polymers with good solubility and compatibility with TMDs are often used to form 2D layered structures. These polymers should have functional groups that can interact with TMDs via non-covalent interactions or covalent bonding. Polymer-TMD composites are used in flexible electronics, coatings, and membranes, where the polymer matrix can provide mechanical support and enhance material properties. (iv) Organic dyes: organic dyes can be used for PL studies when combined with TMDs. The absorption and emission properties of the organic molecules should complement the TMD's properties. This combination is used in PL spectroscopy and imaging to study the electronic and optical properties of TMDs.²³ (v) Conjugated polymers: conjugated polymers can be combined with TMDs to create hybrid materials with unique electronic properties. The polymer should be chosen to match the TMD's electronic band alignment whether type I, type II, or type III. These hybrids find use in photovoltaics, light-emitting transistors, and other organic electronics.^24–28 In all the cases, the choice of organic molecules should consider their compatibility with the 2D TMD surface, the intended application, and the desired chemical and physical properties of the resulting composite material. Surface functionalization techniques such as chemical modification or physical adsorption are often employed to achieve the desired interactions between organic molecules and 2D TMDs.

1.4. Importance of band alignment in determining device performance

The band alignment (Type-I/II/III) between transition metal dichalcogenides (TMDs) and organic materials at heterojunctions is of paramount importance for determining the performance of various electronic and optoelectronic devices. The impact of band alignment on device performance in TMD/organic hybrid systems can be summarized as follows: (i) efficient charge transport, (ii) reduced energy barriers, (iii) enhanced photoemission, (iv) tunable emission, (v) enhanced sensing performance, and (vi) quantum confinement effects. For example, a resonance energy transfer (RET) from a red-emitting monolayer WS₂ (1L-WS₂) to a layer of NIR-emitting organic dye molecule–based type-I heterostructure.²⁹ It is found that the total PL yield of the heterostructures is up to a factor of eight higher than the PL yield of pristine 1L-WS₂. A strong decrease in the PL of the dye is observed on 2L-WS₂ when it gives the type II heterostructure. A PTCDA/MoS₂ type I heterostructure results in enhanced PL emission.²⁰ Again, the PL emission spectra of the DY1/MoS₂ type I heterostructure reveal the quenching of DY1 emission and the enhancement of MoS₂ emission. This observation points to a robust electronic interaction between these two materials.³⁰ In type II, in particular p–n junction, TMDO heterostructures, generally the PL emission peak quenches for individual layers; however, recently an extra ILX emission peak from the NIR to IR range has been observed. Due to their spatially indirect nature, interlayer excitons possess long-lifetimes, attributed to the reduced overlap between electron and hole wavefunctions,^31,32 extended spatial coherence of the excitons,³³ quantum emitter,³⁴ and high-temperature exciton condensate.^35,36 Most recently, Thompson et al., has unveiled, in both theoretical models and experimental observations, distinct features of the entire intra- and interlayer exciton landscape in the PL spectra. In particular, Notably, they have identified a significant transfer of intensity from the intralayer TMD exciton to a series of energetically lower interlayer excitons with decreasing temperature.¹⁹ Nevertheless, it is worth noting that type III alignment has not been reported thus far, and its occurrences in 2D/2D heterostructures are relatively scarce. An artificial neural network simulation, utilizing electronic synaptic arrays and trained on handwritten digit datasets, showcased an impressive recognition accuracy of 91.3%.³⁷ However, the organic layer could fill the gap in near future. Therefore, the careful engineering of band alignment between TMDs and organic materials is fundamental for optimizing the performance of hybrid electronic and optoelectronic devices. It facilitates efficient charge transport, reduces energy barriers, enhances light emission, and allows for the tailoring of energy levels to meet the specific needs of diverse applications, ultimately leading to improved device functionality and efficiency.

This review overviews the most recent advent in two-dimensional transition metal dichalcogenide/organic-enabled interlayer exciton emission and emission modulation covering a range of 2D-TMD materials and organic layers and how the band energy alignment affects it. We discuss the suitability of 2D-TMDO material properties for synaptic devices. Their operational principles are thoroughly examined, and this is followed by the presentation of up-to-date device applications, encompassing field-effect transistors, detectors, and phototransistors. In this review, we initially present the remarkable characteristics exhibited by different TMDs and organic semiconductors (molecules and polymers), and we explore the emerging applications enabled by TMDO hybrid heterostructures in Section 1. In Section 2, we highlight some physical and chemical synthesis routes and fabrication processes using various experimental techniques for 2D-TMDs and organic layers, either individually or for hybrid heterostructures. We also outline the challenges and propose ways to improve the experimental techniques for TMDO hybrid heterostructure fabrication. In Section 3, we provide the band energy alignment schemes of various TMDs with respect to a library of organic molecules and polymers. We also discuss the various types of band alignment, namely, type I, type II, and type III. We discuss the band bending and charge transfer processes in the TMDO heterojunction system from the perspective of band structure and molecular orbital theory. In Section 4, we provide well-comprehensive discussions about the most recent breakthrough in the novel room-temperature interlayer exciton emission and its modulation in the TMDO hybrid heterojunction. Unlike the typical TMD-TMD heterostructure, the TMDO system is usually far from lattice matching. In this section, we highlight the up-to-date research progress in the modulation of emission signals in TMDO constituents from the perspective of photoinduced charge transfer, energy transfer, doping, defect healing, and phase engineering. In the main context, we include a table summarizing molecular structures with the HOMO/LUMO level, type of band alignment, and advancement of emission response of diverse TMDO hybrid heterostructures. In addition, we also discuss the emission modulation in the hybrid structure beyond the band energy alignment. Next, in Section 5, we overview the recent progress in 2D-TMDO material–based optoelectronic synaptic devices for neuromorphic applications, covering an array of 2D-TMDs and organic materials, which are investigated for them. The 2D-TMD and organic material property compatibility for the synaptic devices is discussed, and their operating principles are revisited, followed by the current device demonstrations. We also present, in detail, the advancement of the performances of field-effect transistors—a key component of integrated electronic circuits, anti-ambipolar transistors, phototransistors, and a contemporary breakthrough—photo-switchable superlattices. We summarize the up-to-date research on improved FET performances including organic doping, phase engineering, and hybrid heterostructures in a table within the main context. In section 6, we highlight one of the real-life applications—the photodetectors based on TMDO hybrid heterostructures. In addition, we pinpoint the present challenges and possible solutions for TMDO detectors. Finally, in Sections 7 and 8, we conclude this review by outlining the current challenges from synthesis to device applications and outlooks on this topic of emerging TMDO heterostructures.

Recently, a review article has been published on the recent advances in 2D organic–inorganic heterostructures. It focuses on the integration of 2D organic materials with 2D inorganic materials, emphasizing the combination of organic and inorganic layers, synthesis techniques, and the importance and range of device applications including FETs, photodetectors, solar cells, and neuromorphic computing devices. We include this information in the review article with proper citation.³⁸

2. Fabrication of TMDO heterostructures

The morphology and molecular orientation within TMDO heterojunctions are influenced by the specific fabrication techniques employed. Typically, the fabrication process involves initially creating a film of TMD material, followed by the deposition of the organic layer on top (or vice versa, such as in the case of self-assembled monolayers—SAMs).^39,40 The choice of the preparation method for the 2D-TMD material is generally made based on the intended application and the characteristics of the substrate. Thin-film fabrication methods for TMDO hybrid materials can be broadly classified into two categories: solution-phase and gas-phase film deposition. Solution-phase techniques include drop casting, spin coating, blade coating, and printing. In this context, several common approaches for fabricating TMDs and organic layers are briefly discussed.

2.1. Mechanical transfer (top-down method)

Top-down methods, commonly known as exfoliation techniques, are utilized to isolate single or a few layers of transition metal dichalcogenides (TMDs) from bulk crystals (see Fig. 2a). These methods encompass mechanical exfoliation conducted under ambient conditions or in an inert atmosphere, and chemical or electrochemical exfoliation performed in a liquid phase. By stacking mechanically exfoliated monolayer flakes, various TMDO heterostructure (transition metal dichalcogenide/organic) materials can be synthesized. The procedure commences with the exfoliation of TMD flakes onto an ultra-soft viscoelastic polydimethylsiloxane (PDMS) stamp via mechanical exfoliation aided by high-quality Nitto tape. Subsequently, the PDMS stamp, with the flakes facing away from the glass slide, is affixed to a desired glass substrate, and optical microscopy is employed to align the flakes with the target substrate. Via controlled manipulation using a manipulator, the flakes are gently brought into contact, after which the glass slide is gradually elevated, resulting in the separation of the TMD flakes from the PDMS stamp and their specific positioning on the substrate.⁴¹

Fig. 2 Commonly used techniques for TMDO synthesis: (a) exfoliation, (b) chemical vapor deposition (CVD), (c) molecular beam epitaxy (MBE), and (d) spin-coating.

Moreover, mechanical exfoliation has demonstrated its effectiveness in producing high-quality monolayers of transition metal dichalcogenides (TMDs). Although extensively employed in fundamental research and fabrication, this method is primarily utilized for concept testing in a limited number of heterostructure-based devices. It is accompanied by certain limitations such as small flake sizes, variations in layer thicknesses, and low production throughput. Consequently, there is a growing demand for scalable approaches that enable precise control over the lateral size and thickness of TMD layers, ensuring excellent reproducibility for practical applications. Chemical vapor deposition (CVD) growth methods are commonly employed to generate large-area TMD crystals, effectively addressing the scalability requirements for practical implementation.

2.2. CVD (bottom-up method)

Since the exfoliation method tends to yield relatively small flakes, direct growth methods are employed in certain applications that require large-area 2D films. Therefore, utilizing bottom-up techniques enables the direct synthesis of TMD layers onto a desired substrate, provided that the substrate possesses a high melting temperature, such as SiO₂, sapphire, or mica.⁴² Presently, the chemical vapor deposition (CVD) technique stands as a widely employed method for synthesizing 2D TMDs (see Fig. 2b). Later developments enabled the synthesis of micro- and millimeter-sized single-crystal TMD domains. Similarly, CVD is among the most popular means of preparing TMDs with controlled stoichiometry, layer number, grain boundaries, and domain size. The CVD method can be categorized into two distinct ways—the ‘single-vapor zone’ method, also known as the sulfurization method, and the ‘double vapor zone’ method. The ‘single vapor deposition’ method involves the initial deposition of transition metal precursors onto the substrates. Subsequently, through the introduction of controlled heat treatment within an environment of sulfur vapor, these active precursors undergo a series of controlled chemical reactions, resulting in the formation of TMD layers. This approach is commonly employed to synthesize TMD layers with precise control over their composition and morphology. In the ‘double-vapor zone’ method, an inert gas is utilized to transport both the vapor states of transition metal precursors and chalcogens (sulfur or selenium or tellurium) in the tube. This gas flow facilitates deposition onto the target substrate, under controlled thermal treatment. This process, occurring within the realm of physics, enables controllable reactions and transformations, ultimately leading to the synthesis of desired materials with specific compositions and properties. In the ‘double vapor zone’ method, both transition metal precursors such as molybdenum oxide and tungsten oxide and chalcogen precursors such as sulfur, selenium or tellurium in their vapor states are carried by the inert gas flowing in the tube onto the target substrate under thermal treatment. Volatile precursors such as MoO₃ and S are co-evaporated and then reacted to form the desired TMD deposits (e.g., MoS₂) on the surface of specific targeted substrates. In a particular investigation, Kang et al. achieved the successful fabrication of single-layered films of MoS₂ and WS₂ on 4-inch silicon oxide wafers. To achieve this, they utilized Mo(CO)₆ and W(CO)₆ as transition metal precursors, while (C₂H₅)₂S was selected as the sulfur precursor.⁴³ A detailed description of typical CVD precursors and growth conditions utilized for TMDs is given in topical review articles.^42,44 In addition, several variations of the CVD method have been demonstrated. In summary, the CVD process has emerged as a highly effective and preferred technique for the large-scale synthesis of high-quality monolayers and multi-layers of various 2D TMDs. This method has proven in achieving precise control over the growth of TMD structures, enabling the production of extensive areas of TMD films with exceptional quality. In theory, the powder-based CVD growth of MoS₂ involves a sequence of thermal and chemical processes, which can be represented using the following chemical reaction equations:⁴⁵

2MoO₃ + xS → 2MoO_3−x + SO₂ (1)

2MoO_3−x + (7 − x)S → 2MoO₂ + (3 − x)SO₂ (2)

and combining eqn (1) and (2) gives the final equation,

2MoO₃ + 7S → 2MoS₂ + 3SO₂ (3)

In general, the CVD method presents a straightforward and readily accessible method for the synthesis of simple TMDs, primarily sulfides. However, it encounters limitations when it comes to producing more intricate 2D materials, such as FGT, primarily due to the unavailability of suitable precursors. An alternative approach, metal–organic chemical vapor deposition (MOCVD), has demonstrated successful results in the fabrication of polycrystalline MoS₂ on 4-inch SiO₂/Si substrates and the creation of FETs with an impressive 99% yield and satisfactory mobility around 30 V s cm⁻², as reported.⁴⁶ MOCVD offers a range of precursors with varying decomposition temperatures, allowing for the precise control of growth conditions including temperature and growth rate. Crucially, these precursors must prevent gas-phase reactions to maintain the crystalline quality and composition control. Among the most effective precursors are those containing carbon; however, carbon poses a significant challenge as it contaminates the resulting 2D layers. One potential solution to avoid carbon contamination in 2D sulfides is the use of a carbon-free precursor, such as H₂S. However, it is important to note that H₂S is an extremely toxic and highly flammable substance, which typically limited its use in academic research laboratories. An expandable approach for generating extensive monolayer and few-layer WSe₂ is cold-wall MOCVD, where the precursors utilized are W(CO)₆ and (CH₃)₂Se.⁴⁷ Despite these advancements, achieving uniform 2D TMD films with large domain sizes remains a persistent challenge. Additionally, the high toxicity of metal–organic compounds used in MOCVD processes presents a notable drawback that must be considered.

2.3. MBE: a novel approach to high-temperature TMD synthesis

Subliming molecular precursors using a molecular beam epitaxy (MBE) or an organic molecular vapor deposition (MVD) method (see Fig. 2c) onto a substrate offers the possibility of generating ultraclean interfaces, and the thickness of the active layer can be monitored exactly, e.g., by using a quartz crystal microbalance.^48–53 Several growth parameters can be fine-tuned to engineer the layer growth (e.g., substrate type, deposition time and sublimation temperature), offering the possibility to precisely control the film morphology.^54,55 MBE was successfully applied to grow a variety of TMDs; however, the main disadvantage of this technique is that it is relatively expensive, slow, and requires rather cumbersome equipment for maintaining ultra-high vacuum. However, the MVD method allows for the evaporation of a majority of organic molecules without requiring ultra-high vacuum conditions, as these molecules generally do not decompose before reaching their sublimation temperature. The ultraclean and highly controllable interfaces grown via sublimation are well suited for investigating various underlying physical processes concerning interfaces.^13,14

Further, MBE is a versatile method that simplifies precursor development for 2D materials, particularly complex 2D ferromagnetic materials like FGT.⁵⁶ It uses physical vapor deposition (PVD), and refractory metals such as Mo, W, Cr, Fe, and Co are vaporized.^46,57 However, magnetic metals can cause scattered evaporation beams. An alternative is using high-temperature (above 850 °C) effusion cells, but it increases thermal load and contamination risks. MBE plays a crucial role in the synthesis of CrTe₂ and Cr_xTe_y 2D-ferromagnetic materials since many of these compositions naturally adopt a quasi-2D structure, primarily because of significant self-intercalation of Cr.⁵⁸ It is also ideal for creating heterostructures with other 2D materials, essential for spintronic devices. The ultrahigh vacuum environment for MBE allows monitoring growth in situ using various surface characterization techniques, setting it apart from other methods like CVD. It is also attractive that MBE, being a UHV technique, allows for the in situ monitoring of the growth of 2D materials by a number of surface sensitive characterization techniques such as electron diffraction, low-energy electron diffraction (LEED), reflection high-energy electron diffraction (RHEED), scanning tunneling microscopy (STM), X-ray photoelectron spectroscopy (XPS), and angle-resolved photoemission spectroscopy (ARPES), which cannot be applied when any other growth methodology (e.g., CVD) is used.

2.4. Spin-coating

Spin-coating is a widely adopted method to form homogeneous films extending over large areas onto TMDs supported on a substrate.^59–63 The spin-coating process involves depositing a solution containing the desired molecules onto the center of a rotating substrate, which supports the TMDs (see Fig. 2d). Through a carefully optimized rotational speed, the solution is uniformly spread across the entire surface. Generally, a low boiling point solvent is employed, in some cases followed by an annealing step to eliminate the remaining solvent residues. In contrast to drop-casting and dip-coating techniques, spin-coating offers the advantage of producing uniform films whose thickness can be precisely adjusted by optimizing parameters such as the molecular concentration of the solution, the boiling point of the solvent, and the rotational speed.⁶⁴ Nevertheless, being a kinetically driven process, spin-coating does not give a high quality of structural order at the molecular level.

2.5. Drop-casting

Drop-casting is considered perhaps one of the most straightforward techniques for producing a 2D TMDO interface, and it has found application in several works. Generally, a drop of solution is applied onto the TMD crystal that has been exfoliated or CVD-synthesized on a selected substrate. Subsequently, the solvent is allowed to evaporate slowly.^65,66 For this approach, volatile solvents are often employed, and mild thermal annealing is sometimes used to facilitate solvent evaporation. The resulting films from drop-casting are usually non-uniform due to the influence of largely uncontrollable surface-dewetting and solvent evaporation dynamics, which lead to material accumulation at the droplet edges. However, the interface between a drop cast film and a TMD is not necessarily disordered, as the nanoscale molecular arrangement depends on the molecule and TMD involved. For instance, the drop-casting of long alkanes on 2D TMD layer results in highly ordered epitaxial films.⁴⁰ Unfortunately, not every molecule, but some small molecules, displays an adequately high solubility in common solvents (e.g., toluene, acetone, hexane, and 2-propanal). For example, perylene, pentacene, and metal phthalocyanine molecules are not soluble in common solvents. Hence, most small molecules are substituted by many groups to be soluble in the usually available solvents.

2.6. Dip-coating

Dip-coating serves as another widely employed and extremely practical method to fabricate TMDO interfaces.^21,65–69 TMDs supported on a substrate are immersed into a solution, where molecules interact with the TMD surface and impart the desired functionality. Specifically, dip-coating is typically associated with a scenario in which the substrate is immersed into the solution and quickly removed, but soaking refers to the situations in which the substrate remains immersed for some time. In both cases, low–boiling point solvents and gentle annealing are frequently used to achieve solvent-free molecular films. In contrast to the drop-casted molecular film, dip-coated layers are generally very thin, as often only molecules in direct contact with the sample surface become attached to it, while others remain in the solution. Therefore, dip-coating creates molecular monolayers, or even sub-monolayers. Interestingly, it is sometimes found that keeping the TMD immersed in a solution, and mixed with alkane or thiol for a longer time results in a considerable effect on its electronic properties. This phenomenon can be explained by considering that more molecules have time to interact with the ultra-flat TMD surface and impart their functionality.

2.7. Langmuir–Blodgett deposition method

The Langmuir–Blodgett (LB) technique is one of the important and widely employed methods for the fabrication of ultra-thin ordered 2D molecular structures.^70–73 Because the orientation, packing, and long-range structure of each individual layer can be finely adjusted with meticulous precision, it becomes possible to manipulate the thickness and uniformity of the films at the molecular scale. As an example, Ponomarenko et al. employed the LB technique to fabricate an oligothiophene derivative (D2-Und-4 T-Hex).⁷³ The formation mechanism of a uniform 2D monolayer on the air–water interface involves closely self-assembled molecules via intra-molecular π–π stacking interactions. To deposit a single layer via the LB method, the substrate must be placed at an air–water interface prior to cleaning and spreading the material, so that it only needs to be ‘pulled’ through the interface once. Before placing an SiO₂ substrate, it must be cleaned with acetone and isopropyl alcohol for a few minutes, dried with N₂, and then subjected to oxygen plasma for several minutes. This process endows the substrate with the hydrophilic nature needed to form MoS₂, which is also hydrophilic.^74,75 In general, amphiphilic molecules, which contain both hydrophilic and hydrophobic groups, are necessary as soluble materials to form TMDO hybrid materials. Chemical inertness to both organic materials and subphase water plays a crucial role, as it prevents the solvent from reacting with them. Before depositing the films, substrates must undergo hydrophobic or hydrophilic treatments, depending on the desired type of LB films’ preparation requirement. These characteristics could pose challenges to the progress of LB film technology in the optoelectronics field. Nevertheless, the LB technique remains an excellent approach for fabricating macroscopic MoS₂ devices, offering the advantage of enabling deposition across an entire wafer when required.⁷⁵

2.8. Inkjet printing

Various methods have been employed to deposit solution-processed TMDO hybrids onto desired substrates. Among these techniques, inkjet printing holds particular significance from both technological and industrial standpoints due to its capability to achieve low-cost production of large-area films. In all printing processes, including inkjet, screen, and roll-to-roll printing, the ink formulation plays a key pivotal role that fundamentally controls the physical properties of the deposited films. Molecules and polymers are commonly incorporated in TMD dispersion to make functional inks with the desired rheological properties, which have been used to fabricate rudimentary memory elements and photodetectors. Consequently, they present an opportunity to incorporate additional functionalities into CMOS (complementary metal-oxide-semiconductor) technology, following the more-than-Moore approach.^76–78

2.9. Other important parameters, and the effect of temperature, pressure, gas flow rate, substrate interface property, and organic moiety

2.9.1. Temperature. The growth behavior can undergo transition from kinetic control to thermodynamic control at different sublimation temperatures. Precise temperature control has opened up possibilities to fine-tune nucleation density, domain size, and thickness during synthesis. Elevated synthesis temperatures tend to favor the growth of TMD crystals with larger lateral dimensions and fewer defects, while lower temperatures can result in smaller flakes or nanoplatelets. Higher temperatures also enhance the incorporation of chalcogen atoms (S, Se, and Te) into the TMD lattice, potentially impacting the stoichiometry of the TMDs. For instance, Fangfang Cui reported controlled Cr₂S₃ growth with thicknesses ranging from 1.9 to over 10 nm at temperatures between 750 and 950 °C. The shape evolution of Cr₂S₃ was observed to shift from triangular to hexagonal.⁷⁹ Similarly, Cr₅Te₈ nanosheets grown on a mica substrate at temperatures ranging from 600 to 900 °C for a fixed synthesis time of 10 minutes exhibited shape changes from triangular to truncated-triangular to hexagonal.⁸⁰ Conversely, the thermodynamically controlled process is typically employed to produce high-quality crystals with stable atomic structures. However, these methods often require synthesis temperatures above 600 °C, and in many cases, they exceed 700 °C, such as in the synthesis of W and Cr-based TMDs (reference needed). Therefore, a significant challenge lies in developing general methods that enable the synthesis of ultrathin 2D materials using metal or metal oxide precursors at considerably lower growth temperatures, ideally below 700 °C. This is crucial for seamless integration with the conventional semiconductor industry.⁸¹ Recently, B. Qin and co-authors devised an innovative CVD approach that leverages BiOCl to effectively lower the volatilization temperature of various metal precursors. This groundbreaking method facilitates the growth of a wide array of 2D transition metal dichalcogenide (TMD) materials, encompassing both layered and non-layered variants, within a reduced temperature range spanning 280–500 °C.⁸¹ Remarkably, specific alkali metal halides such as NaCl and KCl have been demonstrated to markedly reduce the melting points of metal reactants, rendering them volatile at relatively modest temperatures.^82–84 Additionally, variations in growth temperature can exert diverse effects on the crystallization and growth of transition metal dichalcogenides (TMDs). For instance, a substantial increase in flake size for WS₂ when the growth temperature was elevated from 880 to 900 °C results in an increase in the size from ∼15 to ∼50 μm.⁸⁵ Moreover, Li et al., have provided insights into how alterations in growth temperature and duration influence the sizes, layer numbers, and shapes of WSe₂ flakes.⁸⁶

2.9.2. Pressure. The pressure inside a CVD chamber can vary widely, from a few atmospheres to very low levels, significantly affecting the gas flow behavior. A lower pressure leads to increased gas volume flow and velocity for the same molar flow, while reducing the precursor concentration according to the ideal gas equation (PV = nRT). This low precursor concentration and high mass feed velocity enhance the reaction control. Therefore, to achieve the growth of a continuous TMD film on a wafer scale, it is common to employ a low-pressure CVD method.^87,88 The partial pressure of a specific precursor profoundly influences the layer-by-layer process in TMD growth. For instance, in MoS₂ growth, lower Mo(CO)₆ pressure leads to nucleation at grain boundaries, while higher pressure results in random nucleation on the initial layer, yielding a mix of monolayer and multilayer products.⁸⁹
Substrate interface property. (i) Effect on surface structure: substrate properties, such as lattice mismatch and surface energy, can influence the nucleation, growth orientation, and stacking of TMD layers on the substrate. Epitaxial growth on specific substrates can result in ordered surface structures.⁹⁰ (ii) Effect on chemical composition: the interaction between TMDs and substrates can impact surface chemistry, potentially leading to surface reconstruction or modification.
Precursor delivery. Precursor delivery methods including gas-phase precursors or liquid-phase precursors in solutions can influence the growth mechanisms and surface kinetics, affecting the surface structures and layer stacking. The choice of precursor molecules and their delivery rates can directly impact the chemical composition of TMDs. Thus, controlling precursor delivery is critical for stoichiometric control.

2.9.3. The addition of an organic moiety modifies the intrinsic properties of TMDs. In the straightforward fabrication of electronic and optoelectronic devices, TMD materials are typically grown on SiO₂/Si substrates without the addition of any catalysts. On such SiO₂/Si substrates, various substances including PTAS, PTCDA, and rGO have been employed as nucleation and growth promoters for TMDs.⁹¹ Furthermore, the use of additives such as alkali metal halides (e.g., NaCl, KI, KCl, and NaBr), sodium cholate, and NaOH has become increasingly common to facilitate the growth of TMDs. These additives, also known as growth promoters, have the capacity to interact with high melting point precursors such as WO₃ and MoO₃, forming volatile intermediates.^84,92 Incorporating these growth promoters typically results in larger domain sizes for TMDs and wider growth windows. However, the precise mechanisms underlying the functions of these growth promoters require further investigation for a better understanding. In contrast, sapphire substrates present a unique situation. The specific lattice orientation of the c-plane sapphire, along with atomically smooth sapphire terraces on the surface, influences the orientation preference of TMD growth on such substrates. This preference is associated with the crystal symmetry and surface terraces of the substrates.⁸⁹ For instance, while the SiO₂ substrate allows for the growth of monolayer MoS₂ exceeding 70 μm in size, sapphire substrates enable submillimeter-sized monolayer MoS₂ growth. These techniques are applicable to a range of transition metal precursors including Cr, V, and W. The understanding and optimization of these parameters are crucial for tailoring the surface structures and chemical compositions of TMDs for specific applications such as electronics, optoelectronics, and catalysis. Therefore, it is essential to explore various synthesis conditions to attain the desired TMD properties, taking into account the interplay of these factors.

3. Energy band alignment TMDO

3.1. Band energy formation in TMDO hybrid heterostructures

In the realm of two-dimensional materials, a substantial number of 2D TMD layered compounds are present (see Fig. 1).² A library of molecules and key TMDs, with MoS₂ band energies as reference, are depicted in Fig. 3. In bulk, 2D-TMDs exhibit an indirect bandgap with many-body interactions that facilitate the features of their optoelectronic properties. However, many 2D-TMDs have a direct bandgap in the form of a monolayer, exhibiting a noticeable photoluminescence emission (see Fig. 5b and c).^119,120 Furthermore, much underlying physics can be found by considering the band-alignment, strain effects, doping levels, and manipulation of the surrounding active layer (or dielectric environment). These considerations provide ample scope for implementing heterostructure engineering based on TMD layers. The hybridization of d-orbitals of the transition metal atom (d_z², d_x²−y², d_xy) and the p orbitals of the chalcogen atom (p_x, p_y) in TMDs gives rise to the formation of the conduction band minimum, CBM (d_z²; p_x, p_y), and valence band maxima, VBM (d_x²−y²; p_x, p_y), respectively (see Fig. 5a).⁹³ These bands, CBM and VBM, play a significant role in band bending and charge transfer occurring across the heterointerface to the highest occupied molecular orbital (HOMO) level and lowest unoccupied molecular orbital (LUMO) level of the organic molecule within the TMDO heterostructure. Subsequently, the modulated light emission directly influences the 2D-TMD-based optoelectronic devices. It is interesting to note that the weak TMDO interactions enable the organic molecules to self-assemble according to the vdW weak intermolecular interactions. The vdW interactions between the TMD material and the deposited organic layer result in the molecular packing behavior within the first few overlayers. Importantly, thanks to their ultrahigh surface sensitivity, the optical and electrical characteristics of TMDs can be modified by a physisorbed ad-layer, without the need for strong chemical bonds.^13,14 A collection of molecules and some essential 2D-TMDs with their band energies are illustrated in the schematic in Fig. 3, where the relevant band energies of MoS₂ are taken as a reference. However, further theoretical investigations of these materials are essential to identify promising atomic heterostructures suited for specific device applications, depending on the field of usage.

Fig. 3 Band energy diagram of a few key TMDs, organic molecules, and polymers (abbreviations are mentioned in Table 1). The horizontal red (CBM) and black dotted line (VBM) refer to the MoS₂ energy levels.^{1,13,14,16,25,39,48,59,67,93–118}

3.2. Type of band alignments in TMDO hybrid structures and their impact on properties

Typically, in semiconductor heterostructures, the band alignment can vary based on the disparity in energy between the conduction and valence band extrema within the constituent layers (schematic representation in Fig. 3). The classification of heterojunctions into three distinct types, namely, type I (symmetric/normal gap), type II (staggered gap), or type III (broken gap), is determined by the relative positions of the CBM or LUMO, and the VBM or HOMO within the constituent layers of transition metal dichalcogenide/organic heterostructures. Fig. 4 visually presents the schematic representation of these different heterojunction types—type I, type II, and type III based on the band energy positions. These heterojunction types are distinguished based on the alignment of band energies, offering a range of opportunities for harnessing diverse functionalities within modern electronic and optoelectronic devices.¹²¹ In the context of heterostructures, the lowest-energy states for electrons and holes exist in the same layer in case of a type-I heterostructure. Conversely, for a type-II heterostructure, these lowest-lying states are segregated into distinct layers. This distinction in the spatial arrangement of electron and hole states in different heterostructure types carries important implications for their electronic and optical properties. Moreover, the band energy offsets of TMDO may result in a reduction of lower the CBM of one layer, potentially positioning it below the HOMO of the other layer (a similar effect may occur when considering the VBM of one layer and the LUMO of the other layer). This energy alignment leads to the formation of a type-III heterostructure.^122,123 Under type-I TMDO band alignment, photogenerated electrons and holes effectively undergo transfer from the active material characterized by a larger bandgap to the smaller bandgap. This process results in the spatial confinement of charge carriers in the layer with the lower bandgap and consequently enhances the carrier population, which may strongly influence the light emission properties of 2D-TMD.¹²² Type-I band alignments find extensive applications in optical devices, including light-emitting diodes (LEDs).^16,94 Currently, investigations on type-I TMDO heterostructures and their associated optoelectronic properties are somewhat limited. To achieve efficient charge transfer, it is essential to create a staggered (type-II) energy level alignment at the TMDO interface, where sufficiently large energy offsets exist between the frontier molecular orbital levels and the valence and conduction band edges of the TMD semiconductor. These pronounced energy offsets are crucial in promoting effective dissociation of excitons in TMDO systems (see Fig. 4).^16,124 It is useful for unipolar electronic device applications and photocatalysis. In other words, type-II staggered band alignment allows for the spatial separation of electrons and holes in the heterojunction in favor of various optoelectronic applications. In type-II heterojunctions, the structure confines electrons and holes into energy bands of different layers, hence leading to the implementation of a high-mobility transistor.^121,125,126 Within a type-II heterostructure, electrons accumulate in the layer with the lower CBM, whereas the holes accumulate in the layer with higher VBM. Nonetheless, energy transfer can still occur between the larger bandgap material to the smaller band gap material.¹²² Recent experimental studies have reported significant energy transfer phenomena in type-II heterojunctions based on TMDs.^127,128

Fig. 4 Illustration showing the generalization of band alignment: type-I (normal), type-II (staggered gap), and type-III (broken gap) semiconducting heterostructures. The band colors green and blue (lower panel) represent the VBM of TMDs and the HOMO of organic materials, respectively, while light green and sky-blue (upper panel) represent the CBM of TMDs and the LUMO of organic materials or vice versa, respectively.

Fig. 5 (a) d-Orbital filling and electronic character of various TMDs. Schematic illustration showing the progressive filling of d orbitals that are located within the bandgap of bonding (σ) and anti-bonding states (σ*) in TMDs. The filled and unfilled states are shaded with dark and light blue, respectively. (b) Calculated band structures of bulk to 1L MoS₂. The bold black arrows indicate the lowest energy transitions. (c) Microscopic image of MoS₂ flakes displaying different contrast corresponding to MoS₂ flakes of diverse thicknesses. (b) and (c) Reproduced from ref. 120 with permission from American Chemical Society, copyright 2010.

Type-III heterojunctions are also very useful for precisely tailoring the transition energy from the conduction band to the valence band. This is particularly crucial in tunneling FET, infrared lasers, and photodetectors.¹²⁹ The stronger charge transfer in type-III heterostructures enhances the magnetic proximity effect and the coulombic interaction between the molecular layer and the TMD to enhance the valley splitting. This implies more opportunities to find large valley splittings in type-III vdW heterostructures.^121,130

Table 1 Optical emission behavior of different TMDs and organic molecules/polymers

Organic molecule/polymer (energy level; type)	TMD	Device processing & type of band alignment	Optical properties (PL) (emission, λmax)	Ref.
VD = vacuum deposition, UHV = ultra-high vacuum, CVD = chemical vapor deposition, MBE = molecular beam epitaxy, PVT = physical vapor transport, DP = dip-coatinga For TMDO band energy alignment, see the schematic in Fig. 3.
	MoS2^a	VD & type-II	Quenched (λmax = 710 nm)	Homan et al., 2017^131
				N. A. Shen et al., 2017,^132
				Dong et al., 2017^133
				Ren et al., 2017,^134
				Jariwala et al., 2016^97
				Black et al., 2023^135
	MoSe2	VD & type-II	Enhanced (λmax = 788 nm)	L. Zhang et al. 2018,^16
	HfS2	TD & type-II	—	K. Jung et al., 2020^121
	PtSe2	UHV	—	Lin et al., 2017^136
	WS2	VD & type-II	Enhanced (λmax ≈ 740 nm)	Zhu et al., 2018^18
	WSe2	TD & type-II	Quenched (λmax ≈ 760 nm), ILX emission ∼1.71 eV	L. Ye et al., 2021^137
				Thompson et al., 2023^19
	MoS2	VD & Type II	Enhanced (λmax = 689 nm)	Obaidulla et al., 2019^14 Wang et al., 2018^138
				Habib et al. 2018^139
				Rijal et al., 2020^140
				Rijal et al., 2023^141
	MoSe2	TD & Type II	Quenched (λmax = 689 nm)	S. Park et al. 2021^20
				J. Gu et al., 2018^128
	WS2	VD & Type II	Quenched (λmax = 689 nm)	Z. Ye et al., 2014^142
				Liu et al., 2017^98
	MoS2	PVT & type-II	Enhanced (λmax ≈ 700 nm)	H. Zhao et al., 2019^100
	MoS2	VD & Type II	Quenched (λmax = 710 nm)	Obaidulla et al., 2019^14
	MoS2	OMVD	—	Mrkyvkova et al., 2019^143
	MoS2	VD & Type II	Quenched (λmax = 710 nm)	Yahya et al., 2020^101
	MoS2	PVT & type II	Quenched (λmax = 550 nm)	E. H. Cho et al., 2015^144
				Fucai Liu et al., 2015^1
				X. He et al., 2017^48
	MoS2	VD & type II	Quenched (λmax = 550 nm)	Park et al., 2017^52
	MoS2	VD & type II	Quenched (λmax = 550 nm)	Kong et al., 2022^13
	MoS2	VD & type II	Quenched (λmax = 550 nm); ILX (≈805 nm)	Kong et al., 2022^13
				M. C. Schwinn et al., 2022^145
				Wang et al., 2023^146
	MoS2	Drop casting, vapor deposition & type II	Enhanced (λmax = 670 nm)	G. Ghimire et al., 2018^147
				S Vélez et al., 2015^102
				S. H. Amsterdam et al., 2019^51
	MoS2	Immersing	—	Y. Yang et al., 2020,^69
	MoS2	Immersing	—	Kafle et al., 2016^148
				Y. Yang et al., 2020^69
				Y. Huang et al., 2018^105
	MoS2	Immersing	Enhanced (λmax = 660 nm)	Choi et al., 2016^106
	MoSe2	Immersing	Enhanced (λmax = 790 nm)
	WSe2	Immersing	Quenched (λmax = 750 nm)
	MoS2	Immersing	Enhanced (λmax = 660 nm)	Choi et al., 2016^106
	MoSe2	Immersing	Quenched (λmax = 790 nm)
	WSe2	Immersing	Quenched (λmax = 750 nm)
	ReS2	VD	--	P. V. Acuna et al. 2021^149
	MoS2	—		C. J. Benjamin et al., Nanoscale, 2018^150
	MoS2	VD & type II		N. S. Mutz et al., 2020^124
	MoS2	MBE	—	D. He et al., 2015^109
	MoTe2	VD & Type I	—	H. G. Shin et al., 2020^111
	MoS2	VD & Type II	—	H. G. Shin et al., 2020^111
	MoS2	Spin coating	—	Y. Zhang et al., 2015^151
	MoS2	Spin coating & type II	Quenching	Shastry et al., 2016^25
				Liu et al., 2019^152
	MoS2	Spin coating & type II	—	M. Sun et al., 2019^108
				Bijleveld et al., 2009^153
	WS2, WSe2 and MoSe2	Spin coating & type II	Enhanced & 1.55 eV	M. Tebyetekerwa et al., 2020^112
	WSe2	Spin-coating & type II	—	H. Qiu et al., 2019^59
	MoS2	VD		Osada et al., 2016^154
	MoSe2	VD	Quenched (1.65 ev)
	WS2			C. Zhao et al., 2021^155
	WSe2
	MoS2	VD & Type-I	Enhanced	Kwon et al., Adv. Sci. 2022,^30
Fluorinated fullerene (C60F48) (p-type)	WSe2	VD	—	Z. Song et al., 2017^50,156
ATTO 725 (p-type)	(1L/2L) WS2	Spin-coated	Eenhanced	Morales et al., 2023,^29
		Type-I
		Type-II	Quenched

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4. Strong optical response in TMDO

4.1. Novel phenomena in TMDO hybrid heterostructures

4.1.1. Interlayer exciton (ILX) emission: a recent advent in TMDO heterostructures. Interlayer excitons (ILXs) are a recently discovered class of excitons found within two-dimensional van der Waals heterostructures. In ILXs, the electrons and holes are spatially distributed across different sublayers of the heterostructure. The presence of ILXs can result in a significant improvement in the optoelectronic performance, giving rise to a range of novel phenomena and potential applications. The exploration of interlayer excitons is driven both by investigating the atypical electronic and optical properties of these systems and by their promising avenue for diverse device applications such as spin, valley, and optically related information processing and energy conversion.¹⁵⁷ Moreover, the spatial separation between electrons and holes leads to the establishment of a distinct electric dipole moment, thereby generating repulsive interactions between these dipoles. This unique combination offers a pathway to investigate intriguing many-body phenomena, including the Bose–Einstein condensation and superfluidity of excitons, and exciton transport through the implementation of patterned electrodes.^158–160

The commonly accepted mode of formation of the interlayer exciton is through the intralayer transfer process. Even though TMD layers in a TMD heterostructure often show moiré patterns because of the relative rotation between layers with similar lattice constants, layers in a TMDO heterostructure usually have lattice constants that are vastly different. For instance, the lattice constants of PTCDA (a = 3.78 Å, b = 19.30 Å, and c = 10.77 Å) are larger in comparison to the lattice constants of MoS₂ (a = b = 3.19 Å, and c = 17.65 Å).¹⁴ Importantly, it has also been reported that due to their static dipole moment, the ILX emission energy is tunable by means of out-of-plane electric field,^161,162 thus, establishing a signature of ILX emission that distinguishes it from other intralayer or defect-related states.¹⁵⁷ The optical properties of the TMDO hetero interface are predominantly influenced by the coulombic bound state of an electron and a hole, which are confined within separate layers. This emergence of complex and fascinating excitonic systems is due to the interplay between the geometrical arrangement, layer charge-polarization, and valley physics intrinsic to monolayer TMDs.¹⁴

In a recent study, Y. Kong and collaborators have made an important observation of a new PL emission peak occurring at a near-infrared wavelength of approximately 805 nm. This emission peak is attributed to the presence of interlayer charge transfer excitons, so-called interlayer excitons in type-II VOPc/MoS₂ (p–n junction) TMDO heterostructures (Fig. 6a), which is crucially dependent on the interface mid-gap states.¹³ A new room temperature–stable ILX formation mode via direct interlayer photo excitation in TMDO is discussed for the first time. This ILX emission in MoS₂/VOPc is further confirmed by the Hersam group.¹⁶³ They employed femtosecond transient absorption measurements (TA) to explore the mechanism of photoinduced charge transfer and the generation of interlayer excitons in MoS₂/VOPc-based type-II heterojunctions. Upon excitation of 514 nm laser light at the VOPc/MoS₂ heterojunction, electrons in both VOPc and MoS₂ undergo simultaneous transitions to their respective conduction bands. VOPc is an extremely narrow bandgap p-type semiconductor,¹⁶⁴ while MoS₂ is an n-type semiconductor.¹⁶⁵Fig. 6a shows the comparison of the individual emission peak at ∼665 nm and 880 nm of VOPc. The band energy level of the CBM of MoS₂ is located below the LUMO level of VOPc (Fig. 6c). Thus, the electrons transfer from the LUMO level of VOPc to the CBM level of MoS₂ to reduce their energy and accumulate there in large numbers, effectively reducing the PL of MoS₂ in the heterostructure. Moreover, the analysis of the density of states (DOS) reveals the presence of multiple mid-gap states, distinct from defect-related states. These mid-gap states are observed far from the Fermi level, yet they appear near the LUMO level of VOPc and are in proximity to the CBM of MoS₂, as depicted in Fig. 6c. The pre-occupied, photo-induced charge carriers can relax to these metastable, empty mid-gap states. In addition, the inrush electron could also transfer from the CBM of MoS₂ to the mid-gap states because their energy levels are energetically very close to each other. As a result, net electrons accumulate near the valence band maxima of MoS₂ in the heterostructure, as do net holes in the interface mid-gap states of VOPc. They form an interlayer electron–hole pair, i.e., ILX. As a result, a new sharp-transition (emission) peak with an energy of ∼1.54 eV (805 nm) in the NIR region appears.¹³ This type of observation is reported for the first time in TMDO.

Fig. 6 (a) PL emission spectra of the VOPc/MoS₂ heterostructure system. A striking feature is a NIR PL appearing around 805 nm (∼1.54 eV) shown in the yellow-shaded box and inset PL mapping. (b) Total density of state (TDOS) of VOPc/MoS₂ heterostructures. (c) Schematic representation of the energy band alignment of VOPc/MoS₂ (shaded area indicates the possible ILX formation). (d) PL spectra of 1L-WS₂, Tc thin film, and a 1L-WS₂/Tc. The new emission band at 1.7 eV signifies the formation of ILX. (e) Schematic of the formation of CT excitons and (f) type-II electronic band alignment of the 1L-WS₂/Tc heterostructure. (a)–(c) Reproduced from ref. 13 with permission from Royal Society of Chemistry, copyright 2022 and (d) reproduced from ref. 19 with permission from Royal Society of Chemistry, copyright 2023. (e) and (f) Reproduced from ref. 18 with permission from AAAS, copyright 2018.

In molecular physics, charge-transfer (CT) excitons emerge as a result of an electron and a hole occupying neighboring molecules. These excitons are most commonly observed in organic and molecular crystals, where they exhibit a static electric dipole moment.¹⁶⁶ The lifetime of a charge transfer exciton at the TMDO heterointerface has been found to be extremely short^131,167,168 However, due to the insufficient screening of the interfacial coulombic potential, the spatially separated electron–hole pairs do not exhibit as free quasi particles but rather form bound states with energies in the range of hundreds of MeV.¹⁰ This phenomenon gives rise to the formation of charge-transfer (CT) excitons, which exhibit similarities to interlayer excitons. The exact nature of the CT states that emerge at the interfaces of TMDO materials in TMD incorporated-molecular semiconductors remains unresolved. In recent research, detailed investigations were conducted to study the formation and transport of charge-transfer (CT) excitons in a type-II heterojunction involving 1L-WS₂ and tetracene (Tc) molecules.¹⁸ The PL spectrum shows a notable resonance at 2 eV, primarily from the bright intralayer KK exciton at higher temperatures (>200 K), as depicted in Fig. 6(d). Upon decreasing the temperature, a secondary peak at approximately 1.71 eV becomes evident, aligning with the anticipated position of the intralayer exciton (ILX).¹⁹ At lower temperatures, there is a greater occupancy of the energetically lower ILX states, resulting in a more pronounced ILX resonance within the PL spectra. Consequently, the temperature serves as an externally adjustable parameter for tuning the relative visibility of ILXs. The CBM of the 1L-WS₂ is at −3.4 eV, which is lower than the lowest unoccupied molecular orbital (LUMO) level of Tc molecules at −2.4 eV. This energy arrangement facilitates electron transfer from tetracene (Tc) to WS₂. However, the VBM of WS₂ is located at −5.8 eV, which is lower than the highest occupied molecular orbital (HOMO) level of Tc at −5.4 eV, enabling hole transfer from WS₂ to Tc. (Fig. 6f). The emission band detected at 1.7 eV is attributed to the presence of interlayer charge-transfer (CT) excitons. Further explanation is given that the quasi-flatness of the molecular dispersion means that ILXs are spread over a larger area of momentum space, reflecting its real space localisation. Furthermore, since the Brillouin zones of WS₂ and Tc layers are considerably different, they assume that all ILXs are bright as no momentum transfer is required. In this TMD/organic heterostructure, electrons are localized on the 1L-WS₂ layer, while holes are localized (or could carriers be almost immobile with heavy effective mass) on the Tc film.^19,169 The distinct energy band alignment between the two materials gives rise to the formation of these interlayer CT excitons. The investigation reveals further that the photoexcitation energy below the Tc bandgap at 2.1 eV exclusively excites WS₂ while suppressing the background emission from Tc at an energy similar to those of the CT excitons. The broad emission observed from the interlayer CT excitons indicates a distribution in the CT exciton energy levels. The charge-transfer (CT) excitons demonstrate significantly longer photoluminescence (PL) lifetimes compared to the singlet exciton of Tc (∼100 ps) or the A exciton of WS₂ (∼500 ps) (Fig. 6d). To further confirm this, photoluminescence excitation measurements were conducted, verifying that the excitation of the A exciton in WS₂ results in CT exciton emission.¹⁸ These findings present a promising pathway for manipulating the optoelectronic characteristics of van der Waals heterojunctions comprising 2D TMD materials and organic semiconductors. Further harnessing the structural deformability inherent in a two-dimensional molecular crystal is a way to introduce a periodic nanoscale potential, which has the capability to confine ILXs.¹⁴¹ The periodic modulation of molecular orbital energies serves as highly efficient trapping sites for IXs. Notably, when ILXs are generated within the PTCDI/MoS₂ system, a swift spatial concentration of electrons within the organic layer near the interface can be observed, underscoring the remarkable efficacy of these interfacial ILX trapping mechanisms.

4.1.2. Modulation of interlayer exciton emission: a recent versatile chemical treatment approach. Strongly bound ILXs in atomically thin MoS₂/WSe₂ heterostructures have recently reported with promising optoelectronic properties for spin-valleytronics and excitonic devices. It is highly desirable to probe and control ILXs for the development of spintronic applications. Recently, Ji and coworkers showed a versatile and novel chemical method for modulating the interlayer excitons in TMD heterostructures.¹⁷⁰ It is further shown for hybrid TMD/TMD/organic heterolayers that ILX could be preserved intact, quenched completely, or modulated.¹⁷⁰ The hybrid multi-heterointerfaces are considered to have different energy level alignments, including consecutive type-II/type-II (termed type-A), type-II/type-I (type-B), and discontinuous type-II/type-II (type-C) alignments. Monolayer MoS₂ and WSe₂ along with two sets of molecules, i.e., tetracyanoquinodimethane (TCNQ) and eosin Y (EY) are used to achieve various heterojunction alignments (see Fig. 7a–f).^171–173 The combinations of TMD and organic layers determine photoinduced charge transfer and dark-state doping, ultimately modulating ILXs. It is observed that the hybrid heterointerface of WSe₂/MoS₂/TCNQ (type-A) leads to quenching of ILX emission via photoinduced electron transfer to the organic layer (Fig. 7a), while WSe₂/MoS₂/EY (type-B) exhibits well-preserved ILXs with the electron transfer suppressed (Fig. 7d). Moderate changes in ILXs are observed with MoS₂/WSe₂/TCNQ (type-C) interfaces. The contrasting properties can be elucidated by analyzing the energy difference between WSe₂ and the organic layers. The energy gap between CBM of WSe₂ and LUMO of TCNQ is approximately 1 eV, which is greater than the energy gap between WSe₂ CBM and EY LUMO (approximately 0.4 eV).^171,174 With a larger energy gap, the TCNQ layer will induce more exciton dissociation and electron transfer at the WSe₂/TCNQ interface, resulting in a stronger quenching of intralayer WSe₂ PL and ILX emission. Further, surface potentials of the hybrid structures are examined by Kelvin probe force microscopy (KPFM), which reveals p-doping effects in the organic layers, regardless of photoillumination. It is found that ILXs are predominantly regulated by the photoinduced process rather than dark-state doping.

Fig. 7 (a) PL spectra of MoS₂/WSe₂ (red) and MoS₂/WSe₂/TCNQ (purple). Peak positions of MoS₂ exciton (blue), WSe₂ exciton (red), and ILX (black) of MoS₂/WSe₂ are marked with vertical dotted lines. While XWSe₂ emission is quenched completely, interlayer excitons are reduced moderately by the TCNQ layer. (b) Energy level diagram of MoS₂/WSe₂/TCNQ, which forms a discontinuous type-I/type-II alignment (Type-C). (c) WSe₂/MoS₂/TCNQ heterointerface, photoinduced electron transfer from MoS₂ to TCNQ (dotted line) resulting in the quenching of ILXs PL emission. Doping-induced charge transfer may also occur, as illustrated by the black solid line. (d) PL spectra of MoS₂/WSe₂ (red) and WSe₂/MoS₂/EY (green). ILX PL emission is decreased slightly with the EY layer. (e) Band energy alignment diagram. (f) WSe₂/MoS₂/EY hybrid structure; photoinduced charge transfer to EY is suppressed, and interlayer excitons generate a strong PL signature even with p-doping (solid line). (a)–(f) Reproduced from ref. 170 with permission from American Chemical Society, copyright 2020.

Hybrid TMDO heterostructures comprising MoS₂/B₃PyPB/WSe₂ or MoS₂/EY/WSe₂ demonstrate distinct and unique optoelectronic properties in comparison to MoS₂/WSe₂. The nature of the organic molecules significantly influences these optoelectronic characteristics.¹⁷⁵ Additionally, investigations have revealed that the introduction of large energy gap organic molecules such as 1,3-bis(3,5-dipyridin-3-ylphenyl)-benzene or B₃PyPB reduce the dielectric constant heterostructures. The organic layer also increases the distance between electron and hole charges in ILX. Consequently, the incorporation of organic layers results in a notable blueshift in the interlayer exciton (ILX) emission. Apart from the dielectric screening, the band alignment with the organic layer plays a significant role in altering the interlayer emission substantially. The inclusion of EY creates an energy state situated between MoS₂ and WSe₂, inducing a transition from tunneling to band-assisted transport. Consequently, the ILX emission experiences an even higher-energy blue shift. Additionally, the use of electron- or hole trapping layers, such as TCNQ or CoPc, respectively, can lead to the complete suppression of the interlayer emission.¹⁷⁵ Therefore, chemical modification of TMDs provides a versatile and effective approach toward reconfigurable, high-performance optoelectronic and excitonic devices using 2D heterostructured TMDs with modulated properties. These results provide critical insights into interlayer excitons at the TMD/TMD heterointerfaces and open up an adaptable approach for selectively tailoring them for optoelectronic applications. With the vast library of organic dyes, this approach could become a powerful platform to understand and regulate excitons in 2D materials. For example, photochromic molecules such as di-arylethenes can be used to switch the energy levels in response to external photoirradiation.⁵⁹ Thus, both interlayer and intralayer exciton behaviors could be controlled dynamically and remotely.

4.1.3. Enhancing photoinduced charge transfer: exploiting the strong optical response in TMDO. The TMDO heterojunctions are often considered solely in terms of their energy level alignment, i.e., the relative energies of the frontier molecular orbitals versus the TMD material conduction and valence band edges. In this model, a charge transfer between a TMD and organic layer occurs in the presence of an appropriate alignment between the relevant energy level in the TMD and molecular HOMO/LUMO levels. In particular, if the molecular LUMO lies at an energy below the conduction band in a TMD semiconductor (or the Fermi level in a TMD semimetal), it becomes energetically favorable for an electron in the TMD conduction band to occupy the molecular LUMO, resulting in electron transfer from the TMD to the molecule and hole doping of the TMD. Conversely, when the molecular HOMO resides above the valence band in a TMD semiconductor (or above the Fermi level in a TMD semimetal), electron transfer occurs from the molecule to the TMD, leading to electron doping, as illustrated in Fig. 3. To achieve effective charge transfer doping, p-type (n-type) molecular dopants are deliberately designed with LUMO (HOMO) levels significantly distant from (close to) the vacuum level. In either scenario, once the charge is transferred to the TMD, the dopant molecules remain as adsorbed species on the TMD surface. As a result of this carefully engineered TMDO heterostructure, charge transfer doping proves to be highly efficient. To assess the extent of charge transfer occurring between organic molecules and monolayer transition metal dichalcogenides (TMDs) in a vertical heterostructure, the line density of the charge difference along the z-axis (normal line to the TMD surface) can be calculated using the expression . Here, Δρ(x, y, z) represents the charge density difference of the adsorption complex concerning the noninteracting individual units, i.e., Δρ(x, y, z) = ρ_Org/TMD − ρ_Org − ρ_TMD.^176,177 In experimental investigations, photoelectron spectroscopy techniques (including photoemission and inverse photoemission) and scanning tunneling spectroscopy are employed to ascertain the energy positions of the frontier states and core levels within the organic and TMD layers. By analyzing the shifts in core or frontier levels in the energy diagram, one can calculate the alterations in the local electrostatic potential and probe the valuable insights into the type of doping and the degree of charge transfer occurring in the system. To comprehend the impact of adsorbing organic layers on the electronic and interfacial characteristics of TMDs, the plane-averaged charge density difference (CDD) is analyzed for the TMDO heterostructures.¹⁷⁶ To quantitatively analyze the charge transfer, the line density of the CDD along the z-axis (normal to the TMD surface), denoted as Δρ(z) is typically studied. In recent theoretical calculations involving 2D-TMD (MoS₂, MoSe₂, WS₂, WSe₂) and organic (Pentacene, PTCDA) heterostructures, diverse amounts of transferred charge have been observed, ranging from 0.0018 × 10⁻³e to 0.0109 × 10⁻³e. To quantify the charge transfer, one can typically study the line density of CDD along the z-axis (normal to the surface of TMD) (Δρ(z)). In recent theoretical calculations involving 2D-TMD (MoS₂, MoSe₂, WS₂, and WSe₂) and organic (Pentacene, PTCDA) heterostructures, diverse amounts of transferred charges have been observed, ranging from 0.0018 × 10⁻³e to 0.0109 × 10⁻³e.¹⁷⁶ Again, the charge transfer analysis reported that there is a net amount of 0.059 electrons per CuPc molecule and 0.052 electrons per TiOPc molecule transferred to the respective monovacancy in MoS₂ surfaces.¹⁰⁴ Further, density functional theory (DFT) studies reveal that pentacene adsorbed on 1T-type monolayer MoS₂ has a large degree of charge transfer ranging from 0.44 to 0.87 e per molecule, which, in turn, changes the Fermi energy level of MoS₂ by up to 1 eV.^132,178 The degree of charge transfer per molecule relies on the precise molecular adsorption geometry and the diverse phases of the substrate.¹³² Furthermore, the hybridization of 2D MoSe₂ with pentacene in type-I band alignments enables efficient and layer-dependent exciton pumping across the TMDO interfaces. Interestingly, the efficiency of interfacial exciton pumping surpasses that of the photoexcitation process in MoSe₂/pentacene by more than 86 times, despite the pentacene layer exhibiting much lower optical absorption compared to MoSe₂.¹⁶Fig. 8c depicts the electronic band alignment between pentacene and MoSe₂,^93,179 where the MoSe₂/pentacene forms a type-I heterojunction. In particular, the HOMO of pentacene lies below the VBM of 1L MoSe₂ and the LUMO of pentacene resides above the CBM of 1L MoSe₂. As illustrated in Fig. 7c, when carriers are photoexcited, they transfer from pentacene to MoSe₂ due to the significant built-in potential and subsequently recombine radiatively. This process leads to an increase in the PL intensity in the MoSe₂ layer and a reduction in the PL intensity in the pentacene layer, as shown in Fig. 8a. This substantial enhancement in pumping efficiency can be attributed to the elevated quantum yield in pentacene and the ultrafast energy transfer occurring within the TMDO interface. Additionally, the precise dielectric environment engineering provided by the organic counterparts significantly alters the binding of charged excitons in monolayer MoSe₂. Further, a strong PL enhancement and quenching behaviors were also reported in the PDI/MoS₂ heterostructures (here, PDI refers to perylene derivatives PTCDA or PTCDI-Ph) (Fig. 8d).¹⁴ There are several key reasons provided for this enhancement of the PL signal in the PTCDA/MoS₂ heterostructure: first, an increased carrier transition probability because of the enhancement of the DOS (due to significant hybridization of C-p_z and Mo-d_z² orbitals). Second, the PTCDA/MoS₂ heterostructure exhibits an increased exciton/trion ratio, and it effectively suppresses nonradiative emission.^14,139 Further, resonance energy transfer (RET) from a red-emitting WS₂ monolayer (1L-WS₂) to a layer of near-infrared (NIR)emitting organic dye molecule-based type I heterostructures is demonstrated.²⁹ It is found that the total PL yield of the heterostructures is up to a factor of eight higher than the PL yield of pristine 1L-WS₂. They have explained that the efficient RET alone is necessary but not sufficient for the desired PL enhancement. Additionally, a key requirement is maintaining the high PL yield of the dye when in direct contact with 1L-WS₂. This necessitates a type-I energy level alignment at the heterointerface. Photoelectron yield spectroscopy confirms this alignment. Although some level positions are uncertain, it is clear that the energy level alignment does not drive exciton dissociation or charge transfer between the dye and 1L-WS₂, thereby preserving a high PL yield in the hybrid structure. Again, a strong decrease in the PL of the dye is observed on 2L-WS₂ when it gives type II heterostructures. The significant PL quenching can be elucidated through the energy level alignment occurring at the interfaces of the 2L-WS₂/Atto hybrid system. The VBM of 2L-WS₂ shifts upward, creating a type-II alignment with the dye's HOMO at ≈5.4 eV. This alignment allows hole transfer from the dye's HOMO to the TMD's valence band, leading to exciton dissociation and reduced PL yield. S Park et al. recently reported on PTCDA/MoS₂ type-I heterostructures with enhanced PL emission.²⁰ Initial expectations pointed towards a staggered type-II energy level alignment and excited state interfacial charge transfer. However, through inverse and direct angle resolved photoelectron spectroscopy, they discovered a straddling type-I level alignment between PTCDA and ML-MoS₂ on insulating sapphire. PTCDA exhibited a wider energy gap. Further analysis through PL and sub-picosecond transient absorption revealed sub-picosecond resonance energy transfer, involving electron–hole pair (exciton) transfer from PTCDA to ML-MoS₂. This transfer significantly enhanced the PL yield from ML-MoS₂ in the heterostructure, consequently modulating the overall photoresponse. Again, the PL emission spectra of the DY1/MoS₂ heterostructure showed quenching of the DY1 emission but enhancement of the MoS₂ emission, indicating a strong electronic interaction between these two materials. The MoS₂ flakes coated with DY1 via evaporation displayed a 1.7-fold increase in PL emission compared to the untreated MoS₂ flakes. Minimal changes in the light-induced contact potential difference (CPD) at the DY1/MoS₂ interface eliminated the possibility of an interlayer charge transfer process, confirming the formation of a type-I heterojunction.³⁰ Again, it is found that while the excited-state electron transfer rate from TMDs (MoS₂ and WSe₂) to fullerenes (C₆₀) is relatively insensitive due to the band offset (more details discussed earlier), the electron transfer (ET) rate from TMDs to PTCDA is reduced by an order of magnitude when the band offset is quite large.¹⁴⁰ For the perylene crystal, the sensitivity of the electron transfer rate on the band offset is elucidated by the 1D nature of the electronic wave function, which limits the availability of states with the appropriate energy to accept the electron.¹⁴⁰ Recent research has demonstrated that using organic superacids like TFSI for MoS₂ surface passivation significantly reduces defect-induced nonradiative recombination, resulting in a near-unity PL efficiency.¹⁸⁰ Moreover, employing a p-doping procedure, HBr efficiently suppresses trap-state emissions and amplifies the emission of neutral excitons and trions in defective monolayers of MoSe₂. This results in a remarkable 30-fold increase in the overall PL intensity at room temperature.¹⁸¹ These defective states can also align with the orbital states of functionalized molecules, influencing carrier properties in TMD materials. Most recently, Liu's group has recently introduced a novel method involving pre-melting and re-solidification of chalcogen powders (S and Se). This technique ensures a consistent supply of chalcogens, facilitating the high quality and uniform growth of monolayer TMDs. After annealing the chalcogen powder in an inert environment, the precursor's surface area decreases significantly, which is beneficial for the subsequent growth of 2D TMDs. This method has been successfully extended to the growth of other high-quality 2D TMDs, including MoS₂, WSe₂, and MoSe₂, demonstrating its versatility and effectiveness.^182,183

Fig. 8 (a) PL emission spectra of 1L MoSe₂ and WL-PEN type-I heterostructure. (b) Calculated density of states of the MoSe₂/pentacene heterostructure. (c) Type-I band energy alignment of a MoSe₂/pentacene heterostructure. (d) PL spectra and PL maps of the PTCDA/MoS₂ type-II hybrid structure (orange), PTCDA (red), and MoS₂ monolayer (blue). (e) DFT study of PTCDA/MoS₂ heterostructures: TDOS and PDOS of the PTCDA/MoS₂ heterostructure, PTCDA and MoS₂. (f) Type-II band alignment for PTCDA/MoS₂ schematically showing the photoexcited carriers’ transfer within the hybrid heterostructure interface. (a) and (c) Reproduced from ref. 16 with permission from Wiley-VCH, copyright 2018. (b) Reproduced from ref. 176 with permission from Wiley-VCH, Copyright 2020. (d)–(f) Reproduced from ref. 14 with permission from Wiley-VCH, copyright 2019.

4.1.4. Energy transfer. Nonradiative Förster resonance energy transfer (FRET) encompasses the efficient transfer of an excited electron–hole pair from an emitter (donor) to an absorbing medium (acceptor), facilitating the directed flow of energy from the donor to the acceptor.^128,140,184 The electron–hole coupling occurs through near-field interactions, wherein no photons are emitted and can be understood to the quenching of donor dipoles in the presence of a lossy medium. Several key factors influence the efficiency of nonradiative energy transfer, including the spectral overlap between the donor's emission and the acceptor's absorption, the spatial separation between the donor and the acceptor, the dimensionality of the materials involved, and the relative alignment of the donor and acceptor dipole moments.^128,185 An efficient FRET was observed between excitons in monolayer MoSe₂ and few layers thick PTCDA.¹²⁸ The exciton transition dipoles are arranged horizontally, facilitating efficient energy transfer between these distinct materials with different lattice constants. This energy transfer phenomenon is observed experimentally through time-resolved and steady-state PL and PL excitation spectroscopy. Time-resolved measurements indicate a decrease in the lifetime of the donor PTCDA, while steady-state emission experiments show a concurrent decrease in the emission intensity of the donor (PTCDA) and an increase in the emission intensity of the acceptor (MoSe₂). The PL excitation spectra show a spectral dependence of the energy transfer process, with the highest efficiency occurring at the donor's absorption peak. Fourier space imaging is employed to ascertain the planar dipole orientation. Notably, efficient energy transfer is expected in the presence of a low mobility organic layer.

The Fourier plane images of the emission intensity reveal a horizontal alignment of the excited state dipoles for both MoSe₂ and PTCDA. As a result of efficient energy transfer across the interface, the emission intensity of the donor (PTCDA) decreases, while that of the acceptor (MoSe₂) increases. Moreover, the energy transfer is accompanied by a shortened fluorescence decay time of the donor, considering a new channel for nonradiative recombination from the donor's perspective. When a 10-nm-thick spacer (Al₂O₃) was introduced and a new heterostructure of PTCDA/Al₂O₃/MoSe₂ formed, the energy transfer process was blocked; thus, no discernible change in lifetime was observed in the new heterostructure. The exciton lifetime of MoSe₂ is considerably shorter than that of PTCDA. However, due to Förster resonance energy transfer (FRET), the excited state of MoSe₂ continues to be populated, leading to a prolongation of the effective PL lifetime of MoSe₂ in the heterostructure. PL excitation experiments have revealed that the emission intensity of MoSe₂ reaches its maximum at the absorption peak of PTCDA in the heterostructure, providing direct evidence of efficient energy transfer from PTCDA to MoSe₂.¹²⁸

Finally, recently, the energy transfer on a prototypical heterostructure consisting of tetracene (Tc) and a monolayer WSe₂ was studied using ultrafast spectroscopy.¹³⁷ Upon photoexcitation in type-II WSe₂/Tc heterostructure, an ultrafast (∼3.4 ps) singlet exciton energy transfer from tetracene to WSe₂ occurs, followed by a 16.5 ps hole transfer from WSe₂ to tetracene, resulting in long-lived (∼565 ps) charge separation. Notably, this charge separation takes place prior to the slow singlet fission (SF) process (>20 ps) in tetracene. However, this study presents various challenges and emphasizes the need for deliberate design to enhance singlet fission in 2D optoelectronic devices.

4.2. Doping

It is well known that defective states in TMDs, which are formed within the bandgap near to the VBM or near to the CBM, can play a significant role for carrier mobilities and quasiparticles such as excitons and trions.^186–188 Furthermore, these types of states are often associated with dangling bonds, which may serve as active sites for the chemical or physical adsorption of different species.^189–192 In particular, the interplay between the localized defective states and organic molecular orbitals may result in new optical excitations and modulation of electron–hole recombination processes.^193–195 A recent study has showed that the passivation of the MoS₂ surface with organic superacids greatly suppressed the defect-mediated nonradiative recombination and led to near-unity photoluminescence.¹⁸⁰ In addition, resonance of these defective states with the orbital states of functionalized molecules or atoms was found to enable the rectification of carrier content and polarity in the host TMD materials. However, the combined impact of doping molecules and sulfur vacancies on the charge transfer and electronic properties of monolayer MoS₂ remains largely unexplored. The interactions between molecules and semiconducting 2D materials have been studied before. For example, the control of charge carriers and PL emission by hybrid 2D-TMDs sheets such as MoS₂^39,40 and WS₂^66,196 with chemical functional groups was investigated. Many prototype devices, demonstrating the robust molecule-MoS₂ interaction, were engineered to realize highly effective gas sensors and detectors. Physisorbed or chemisorbed molecular groups were shown to lead to efficient layer-by-layer exfoliation and significant modification of the surface properties of MoS₂ layers.¹⁹⁷ In addition, because of the atomically thin nature, the accumulated or depleted charges transferred between the molecules and MoS₂ were demonstrated to result in a remarkably high sheet charge density.^21,197 Therefore, a proper chemical functionalization treatment of TMDs by an organic layer results in modulated photoluminescence^198–201 and enhanced TMDO constituent stability.¹⁸¹ The most commonly used molecular non-covalent p-dopants are tetracyanoquinodimethane (TCNQ) and tetracyanoethylene (TCNE).^202,203 These materials are strong electron acceptors with a solid-state electron affinity of 5.2 eV, and the planarity of 2,3,5,6-tetrafluoro-7,7,8,8-tetracyanoquinodimethane (F₄TCNQ) enables efficient blending with many hosts.

Moreover, doping is attained by exposing the crucial TMD surface to solutions containing a pentamethylrhodocene dimer as the reductant n-dopant and “Magic Blue” [N(C₆H₄-p-Br)₃]SbCl₆ as the oxidant (p-dopant).²⁰⁴ The current–voltage characteristics of FETs reveal that irrespective of their initial transport behavior all key TMDs can serve as p- or n-channel devices upon appropriate doping. In Fig. 9a, the arrangement of dopant molecules at TMDO interfaces under low coverage conditions is shown, where the molecules lie flat on the surface. Various donor and acceptor molecules are depicted, and their respective energy levels (HOMO/LUMO) are indicated with respect to MoS₂ (Fig. 3). The optoelectronic properties of 2D TMDs (MoS₂, MoSe₂, and WSe₂) were modulated by incorporating two metal-centered phthalocyanine (MPc) molecules, namely NiPc and MgPc.¹⁰⁶ The investigation further demonstrates the selective and reversible engineering of PL emission by facilitating energetically favorable electron transfer from photoexcited TMDs to MPcs.

Fig. 9 (a) Band alignment of MoS₂, MoSe₂, WS₂, and WSe₂ along with effective ionization energy (IE) of the n-dopant and electron affinity (EA) of the p-dopant. (b) and (c) PL emission of MoS₂ after n- and p-doping treatment. (d) Band alignment of MoS₂, MoSe₂, and WSe₂ along with reduction potentials of NiPc and MgPc and (e), (f) corresponding PL spectra. (g) Schematic of functionalization scheme ₂H to 1T phases. (h) HRTEM image of an atomically thin-phase boundary (indicated by the arrows) between the 1T and 2H phases in 1L MoS₂. (i) PL spectra obtained from 1L MoS₂ (2H phase), from the metallic 1T phase and from the functionalized 1T phase. A strong PL peak at ∼1.8 eV is observed for the 2H phase, consistent with single-layer 2H MoS₂, whereas PL is quenched after conversion to the metallic 1T phase. Inset: A typical CVD-grown flake of 1L MoS₂. (j) Modulation of photoluminescence peak intensity with the increasing amount of functionalization. (b) and (c) Reproduced from ref. 204 with permission from Wiley-VCH, copyright 2018. (d)–(f) Reproduced from ref. 106 with permission from American Chemical Society, Copyright 2015. (g), (i) and (j) Reproduced from ref. 199 with permission from Springer Nature, Copyright 2014. (h) Reproduced from ref. 276 with permission from Nature Publishing Group, copyright 2014.

4.3. Defect repairing

In general, commercially available 2D-TMDs may contain defects in their lattice structure, including atom vacancies within the basal and/or edge planes, as well as substitutional defects. Specifically, the presence of chalcogenides (S, Se, and Te) in TMDs contributes to their low sublimation temperature (<200 °C), resulting in their volatility and the occurrence of defects. These chalcogen vacancies create energy band levels within the gap states, which function as charge scatterers and traps. As a result, this phenomenon leads to the quenching of PL, decreased mobility, and enhanced hysteresis properties in TMDs.^186,205,206 Significantly, TMDs obtained through mechanical exfoliation generally exhibit a lower density of inherent defects than the sheets produced via chemical vapor deposition (CVD), chemical, or liquid-phase exfoliation methods. The presence of a high density of defects in the latter limits their optoelectronic applications.⁸⁹ To address this, the concept of ‘defect healing’ or ‘repairing’ becomes crucial in order to enhance the electronic performance of TMDs, making them suitable for large-area fabrication and practical industrial applications. In 2D-TMDs, the surface is often considered inert and challenging to functionalize covalently due to dangling-bond free surface. However, chalcogen (S, Se, and Te) vacancy sites possess reactive dangling bonds that can serve as anchoring sites for specific molecular species.^6,207 As a result, a promising defect repairing strategy involves exploring molecular functionalization at these defect sites. This approach offers potential for enhancing the optical and electronic properties of defective 2D-TMDs.

As an example of a defect healing approach, thiol (-SH) molecules were applied to functionalize sulfur vacancies in TMDs, supported on substrates or in solutions. In this functionalization process, the sulfur S atom in the thiol group (-SH) forms a covalent bond with a sulfur vacancy site, effectively repairing the defect while simultaneously functionalizing the TMD. In another research work, Nguyen et al. employed thiol (R–S–H) chemistry to functionalize the surface of MoS₂, which resulted in the passivation of defects, hence modulating the electronic properties.²⁰⁸ The functionalization of MoS₂ with thiols induced alterations in its chemical properties, offering potential applications in the design of efficient sensors and catalysts. Two categories of thiols were investigated: (i) thiols with aromatic rings possessing distinct electron-withdrawing capabilities, such as p-mercaptophenol and thiophenol and (ii) alkyl thiols with varying chain lengths, including 1-propanethiol, 1-nonanethiol, and 1-dodecanethiol.²⁰⁹ Thiol functionalization has demonstrated the capability to enhance the electronic properties of MoS₂ devices by repairing intrinsic defects or controlled vacancies introduced via ion irradiation. Additionally, annealing processes were found to cleave the S–C bonds, releasing the thiol-grafted molecule and restoring the integrity of the MoS₂ lattice. The incorporation of thiol-based dopant groups not only facilitated defect healing but also introduced simultaneous doping effects. Thus, grafting the functional groups may result in the passivation of defects and reduction of the charge scattering sites. Thiol functionalization strategies have recently demonstrated that they effectively advance the optoelectronic performance of TMDs such as WSe₂.²¹⁰ This effective approach has proven successful not only for transition metal disulfides but also for TMDs with Se and Te terminations.

Defects in TMDs fall into two main categories: chalcogen vacancies and transition metal vacancies. Chalcogen vacancies, where chalcogen atoms (typically S/Se/Te) are missing from the crystal lattice, are more common. These defects can be unintentionally generated by adjusting growth parameters such as temperature, gas flow, precursor ratios, and substrate. Temperature variations and thermal annealing can also induce structural changes in chalcogen atomic defects within 2D TMDs. Understanding this evolution is crucial to prevent surface/interface degradation and holds significance for future device applications. For instance, annealing monolayer MoTe₂ at moderate temperatures leads to the formation of Te vacancies and pinholes, with their extent increasing as the temperature rises. Further, in TMDs, chalcogen vacancies (S, Se, and Te) can be effectively filled with isoelectronic oxygen atoms and chloride ions.^68,211,212 Specifically, laser irradiation in air has been reported to greatly enhance the optoelectronic properties of defective WSe₂. This process facilitates the binding of oxygen atoms at Se vacancies.²¹² For instance, laser irradiation can generate Te vacancies in MoTe₂, leading to a phase transformation from 2H to 1T.²¹³ Additionally, plasma treatment, whether with oxygen plasma or argon plasma, proves effective in creating chalcogen atomic defects in 2D TMDs.²¹⁴ Tailoring chalcogen atomic defects to manipulate properties and enhance device performance can be achieved by adjusting synthesis parameters to introduce stoichiometry deviations during crystal growth and by applying post-growth treatments to selectively create these defects. For instance, Pranjal Kumar Gogoi demonstrated oxygen passivation in MoS₂, altering exciton and trionic properties.²¹⁵ It is also highlighted that the possibility of substitutional oxygen atoms is more prevalent in the crystal structure of TMDs. Therefore, nanoscale characterization techniques were utilized to investigate the defect origins in MoSe₂ and WS₂. These oxygen substitutional atoms can serve as anchoring sites for further functionalization, similar to what has been previously discussed for SAMs on the top surface of TMDs.⁶⁴

The defect healing approach for chalcogens other than sulfur (S), such as selenium (Se) and tellurium (Te), is indeed a valuable aspect of this research field. An air-stable p-doping of WSe₂ has been successfully demonstrated through the chemisorption of NO₂ at 150 °C. This innovative process led to a remarkable reduction in contact resistance with the Pd metal by five orders of magnitude, resulting in a degenerate doping concentration of 1.6 × 10¹⁹ cm⁻³. Their work provides a promising pathway for achieving air-stable p-doping in WSe₂, with broad potential applications.⁶⁸ In an exciting development, researchers reported the creation of an optically switchable multilevel FET by interfacing WSe₂ with a bicomponent donor–acceptor-acceptor (DAE) blend. This blend consisted of two DAE molecules with specific energy levels engineered to trap either electrons or holes in WSe₂. The resulting device demonstrated the ability to modulate both electron and hole transport simultaneously through remote light stimuli. This approach holds promise for the development of optically switchable ambipolar FETs and is applicable to various 2D materials.²¹⁶ Moreover, a novel method has been introduced to rectify Te vacancies within defective MoTe₂ by filling them with reactive oxygen via a low-pressure annealing process. This technique resulted in significant improvements in FET characteristics, including greatly enhanced current density control and a record-high hole mobility of 77 cm² V⁻¹ s⁻¹.²¹⁷ A one-step synthesis method was reported for producing atomically homogeneous Fe-doped VSe₂. This material exhibited enhanced catalytic activity, particularly in the hydrogen evolution reaction, making it promising for various catalytic applications.²¹⁸ Defect-induced magnetism was observed in single-layer PdSe₂, primarily attributed to Se vacancies that create in-gap defect states. This discovery transforms PdSe₂ into a 2D magnet and highlights the common occurrence of defects, especially Se vacancies, during exfoliation. The findings position PdSe₂ as an ideal platform for the study of defect-induced phenomena and offer insights for experimental control of defect patterns.²¹⁹

4.4. Phase engineering with covalent functionalization

Chemical modifications such as increasing their solubility in common solvents offer additional opportunities to modulate the properties of 2D TMDs. Although theoretical studies on the functionalization of TMDO have shown promising progress, the experimental covalent functionalization still poses challenges.^{199,209,220–223} In past, functionalization has been generally achieved via reductive alkylation by introducing a halide reagent to n-doped carbon nanotubes and graphene.^224–226 In recent studies, a novel approach to functionalize 2D TMDs involves ligands interacting with unsaturated metal atoms located at the edges or defects on the basal plane. This method opens up a promising avenue for enhancing the properties of these materials.¹⁹⁹ Further, the chemical exfoliation can result in defective sheets, and hence, some of the functionalizing molecules may also bind to defect sites. The organohalide solution-functionalization reactions of the chemically exfoliated 2D monolayer were highly efficient because of the functional group of TMD (e.g., MoS₂, WS₂, and MoSe₂). However, the functionalization reactions are usually considered efficient if the ratio of the functional group per 2D TMD surface area is greater than 25%, which is confirmed from the XPS study.²⁰⁹ In a recent study, covalent functionalization of the basal plane of TMDs has attracted significant attention. As discussed in the preceding section, the strategies employed for defect functionalization also led to the establishment of covalent bonds. This versatility opens up new possibilities for tailoring the properties of TMDs via covalent modifications. In this regard, there are diverse approaches that facilitate covalent bonding without necessarily requiring the presence of defects.

The key distinction between the two approaches lies in their impact on binding molecules to the surface. In the defect healing process, the number of molecules that can bind is limited by the presence of defects. Conversely, in the direct covalent grafting scheme to the TMDs’ basal plane, a larger quantity of molecules can be immobilized. It is further noted that due to the strong inertness of TMDs, aggressive chemical reactions were required to achieve significant covalent functionalization. The covalent functionalization of semiconducting TMDs has primarily been observed in chemically exfoliated compounds obtained through Li intercalation in solutions. This process results in electron transfer from Li atoms, leading to negatively charged mono- and few-layered TMD flakes. A versatile, general and simple method has been successfully employed for the covalent functionalization of 2D TMD nanosheets, including MoS₂, WS₂, and MoSe₂, without the need for defect engineering. Instead, the functionalization reaction involves electron transfer between the electron-rich metallic 1T phase and an organohalide reactant, leading to the covalent attachment of functional groups (such as 2-iodoacetamide or iodomethane) to the chalcogen atoms of the TMD, as illustrated in Fig. 9g and h.¹⁹⁹ The covalent attachment of functional groups to the TMD nanosheets results in significant alterations in their optoelectronic properties. Notably, the metallic 1T phase is transformed into a semiconducting state, leading to pronounced and tunable photoluminescence as well as gate modulation effects in field-effect transistors, as depicted in Fig. 9i and j. Voiry and coworkers reported the formation of C–S bonds on the basal plane of exfoliated MoS₂ through a reduction negative charge using an organohalide reactant. The chemically exfoliated sheets, which were initially negatively charged, underwent the reaction leading to the creation of C–S bonds. This process resulted in a functionalization degree of approximately 20%, determined as the ratio of functional groups to MoS₂ units.^{199,227–229} A similar methodology applied to 2D materials, diazonium salt chemistry, was utilized to covalently functionalize the chemically exfoliated TMDs in solutions. Recently, a chemical approach involving aryldiazonium was adopted to achieve surface functionalization of MoS₂.²³⁰

4.5. Beyond band energy alignment in TMDO: unveiling the influence of molecular orbital interactions, molecular orientation, and thickness variation

When two distinct materials come into contact to form a heterojunction, their electronic states do not remain unchanged or static. Even in 2D vdW interfaces, where the bonding is weak, interactions between orbitals can lead to alterations in the alignment of band energy levels.¹² These perturbations involve specific effects on the orbitals, such as hybridization, as well as more general effects related to charge transfer and electrostatic interactions. In this context, hybridization refers to the emergence of new or mixed electronic states in the hybrid TMDO system, where the electron density originates from both the organic molecule and the TMD material. If this hybridization leads to an apparent alteration in the distribution of electron density, it gives rise to ground-state charge transfer, effectively facilitating an electron transfer from one component to the other. Furthermore, the presence of dipoles and charges within molecules can induce changes in the electronic structure of the TMD material, analogous to the effect of a gate voltage in a FET transistor. As a result, the Fermi level of the TMD material may be shifted either upwards or downwards.¹² Typically, widely used organic semiconductors including planar phthalocyanine molecules, acenes, and perylene derivatives possess neutral charges and exhibit negligible or small net dipole moments. In such instances, the dynamics of charge carriers is primarily influenced by hybridization and ground-state charge transfer phenomena. As an example, in PTCDA/MoS₂ heterojunctions, the process of hybridization combines the S-p_z and Mo-d_z² orbitals near the CBM (Conduction Band Minimum) with the conjugated C-p_z orbitals (as depicted in Fig. 6b).¹³⁹ This hybridization process leads to a reduction in the bandgap of MoS₂ effectively and an enhancement in the density of states (DOS) near the band edges. This effect further causes a significant increase in the intensity of the PL signal, accompanied by a redshift. Furthermore, a charge transfer phenomenon occurs, which is due to the rearrangement of electron density within the hybrid orbitals. The contrasting optoelectronic behaviors of PTCDA and other non-planar perylene derivatives such as PTCDI-Ph become more apparent when compared. The planar geometry of PTCDA facilitates a higher overlap of molecular π-orbitals with the out-of-plane orbitals in the MoS₂ conduction band, leading to more pronounced effects in hybridization and charge transfer. In contrast, PTCDI-Ph's bent structure reduces this overlap, resulting in a less significant impact on the electronic properties of MoS₂.¹⁴ Consequently, there is a decrease in the net charge transfer between PTCDI-Ph and MoS₂, and the surface binding energy obtained from density functional theory (DFT) is 30% lower for PTCDI-Ph as compared to PTCDA.^14,139 In another example of the FePc/MoS₂ heterostructure, angle-resolved photoelectron spectroscopy (ARPES) has revealed the hybridization between FePc and MoS₂ in the vicinity of the K point.²³¹

Further, the orientation of organic molecules, film morphology, and thickness are crucial; however, they are often neglected factors in TMDO heterojunctions (see Fig. 10a).^12,14,53,232 It is widely acknowledged that interactions between molecules and substrates have a substantial impact on the adsorption of individual molecules onto solid surfaces, playing a pivotal role in directing the orientation of these molecules.²³³ Recently, Wang et al. have demonstrated the crystallographic structures of diverse 2D atomic crystals via the strategic utilization of an oleamide molecule as a probe, harnessing orientation-dependent interactions with the substrate.²³⁴ The process of oleamide assembly onto 2D atomic crystals comprises two simple steps: initially, applying a diluted oleamide solution in chloroform onto the 2D materials via spin-coating, followed by annealing the samples at a moderate temperature. Following this process, dense nanoribbons were observed on various hexagonal 2D atomic crystals, including mechanically exfoliated MoS₂, WSe₂, ReS₂ and other 2D materials (Fig. 10b–d). The strikingly different orientations of oleamide nanoribbons on various 2D atomic crystals, each possessing distinct crystallographic structures, strongly indicate that the alignment of these nanoribbons is probably influenced by the atomic arrangements of the underlying substrates.

Fig. 10 Self-assembly of oleamide on different 2D atomic crystals. (a) Schematic of the oleamide self-assembly on 2D atomic crystals along specific orientations. Hydrogen bonds are indicated with blue dashed lines. (b)–(d) Images of oleamide nanoribbons on MoS₂, WSe₂ and ReS₂ respectively. (a)–(d) Reproduced from ref. 234 with permission from Nature Publishing group, copyright 2017.

These nanoribbons exhibited alignment along three high-symmetry directions with rotations of approximately 120 degrees and covered a substantial area. This investigative approach has led to an intriguing revelation: the periodic atomic arrangement in 2D materials directs oleamide molecules to self-assemble into quasi-one-dimensional nanoribbons, exhibiting distinct alignments that serve as precise indicators of the underlying lattice orientations. Using oleamide molecules as probes, they have achieved the successful identification of crystallographic orientations for approximately 12 distinct 2D materials, all the while preserving their intrinsic properties without degradation.²³⁴ As a result, vdW materials offer a distinctive platform for fine-tuning these properties, allowing the growth of high-quality organic phases that may not be observed on other substrates.²³⁵ The orientation of a molecule on a surface plays a crucial role in controlling its interactions with the surface. For example, the significance of orientation has been demonstrated in studies involving H₂Pc and CuPc films on MoS₂. The two phthalocyanine molecules H₂Pc and CuPc exhibit different orientations on MoS₂. CuPc lies in a π-face-on configuration, whereas H₂Pc adopts an approximately edge-on orientation on MoS₂.^{12,51,232,236} The variation in molecular orientation has significantly impacts on the recombination times upon photoexcitation. In the CuPc/MoS₂ interface, excitons more readily dissociate, leading to the formation of long-lived free carriers. Several other large π-conjugated molecules such as pentacene and perylenes have been shown face-on orientation to 2D TMD materials.¹⁴³ The face-on orientation of large π-conjugated molecules such as pentacene and perylenes on 2D TMD materials contrasts with the edge-on orientations observed on conventional surfaces such as SiO₂ and sapphire. This orientation/alignment frequently differs from the edge-on orientations observed on more conventional surfaces such as SiO₂ and sapphire.^232,237 This distinct molecular orientation has led to the emergence of van der Waals materials, particularly hexagonal boron nitride (hBN), as favorable substrates for facilitating the epitaxial growth of functional organic molecules.¹²

The thickness of 2D TMDs is a crucial factor influencing the band energy level alignment, charge transfer phenomena, and prominence of energy transfer observed at TMDO interfaces.^18,131,238 It is demonstrated that the thickness dependence in WS₂/Tc heterojunctions results in distinct changes when transitioning from monolayer (1L) to bilayer (2L) WS₂. Specifically, with the increase in WS₂ thickness, the valence band minimum rises, leading to the blocks of hole transfer to the tetracene layer.¹⁸ With the increase in WS₂ layer numbers, the interface undergoes a transition from Type-II to Type-I, leading to energy transfer during selective tetracene excitation. In the case of a type-I 1L-WS₂/tetracene junction, excitonic energy transfer takes place. To understand this process, dynamics in type I heterostructures constructed with 2L-WS₂ and tetracene was examined. In type I heterostructures, the hole transfer from WS₂ is inhibited, allowing only excitonic energy transfer.¹⁸ The group further observed a correlation between the thickness and the energy transfer speed, with the transfer time increasing from 44 ± 5 ps for 2L-WS₂/Tc to 84.8 ± 4.4 ps for 7L-WS₂/Tc heterostructures. Even when the energy level alignment remains favorable for charge transfer, variations in the thickness of transition metal dichalcogenides (TMDs) can still impact the photophysics of heterojunctions. The photophysics of heterojunctions can still be influenced by variations in the thickness of TMD, even when the energy level alignment for charge transfer remains favorable. Kafle et al. also addressed the thickness dependence of the TMD layer with phthalocyanine molecules, using MoS₂/ZnPc interfaces.⁵³

It is worth to mention that up until now, remarkable progress has been achieved with the study on 2D magnetic TMDs such as CrSe₂, CrTe₂, FeSe₂, FeTe₂ and VeTe₂,^239–242 and novel magnetic properties at the hybrid TMDO interface in the 2D limit remain virtually unexplored for diverse optoelectronic applications. Magnetic molecules, for example, TbPc₂ with 2D TMD could give novel properties in the future with diverse applications. In this regard, MOKE techniques could characterize the ultra-thin magnetic layers with thicknesses in the sub-nanometer range.²⁴³

5. Advancement in field-effect device applications

5.1. Field-effect transistors (FETs)

Over the last decade, the distinctive electronic, chemical, and physical properties of two-dimensional transition metal dichalcogenides have attracted significant attention. Although in 2D family, pristine graphene has ultrahigh mobility around 10⁴ cm² V⁻¹ s⁻¹,^244–246 zero bandgap and poor switching characteristics, which make it difficult to be used as a channel material. Even though pristine graphene belongs to the 2D family and has extremely high mobility, its zero bandgap and poor switching properties make it challenging to use as a channel material. Therefore, due to their advantageous optoelectronic characteristics and large band gaps of 1–3 eV,^247,248 TMDs have gained new research interests, revealing them as a distinct class of materials with distinctive technological potential.^247,248 As vdW heterostructures advance technologically from their humble beginnings, more research groups have access to more advanced tools and materials. In particular, the vdW assembly protocols need to be scaled up for any industrial application.²⁴⁹ Consequently, the study of optoelectronic devices incorporating the synergistic properties of TMDs and molecules has emerged as a prominently researched area in present times.

In this section, we will further discuss on the device architecture and novel properties of TMDO heterostructure-based FETs (Table 2). This investigation highlights the vast potential of hybrid heterostructures in creating novel devices with distinctive functionalities. In both cases, these devices exploit the semiconducting properties of specific organic molecules, while the TMDO interface plays a pivotal role in determining the device performance. As we will highlight, the controllable manipulation of energy bands at the TMDO heterointerface is beneficial to the devices. This modification is achievable through the application of an electric field controlled by a gate terminal.

Table 2 Summary of FET properties in 2D TMDs incorporated with organic materials via heterostructure fabrication or doping

Hetero junction: TMDO (type-I, II, III)	Mobility [in cm^2 V^−1 s^−1]	I on/off ratio	V GS (in V)	Contact electrodes	Fabrication method TMDO	Ref.
MoS2/TiOPc (type-II)	—	>10^7	−50 to 50	Ag/Au	CVD/MBE	Park et al., 2017^52
MoS2/rubrene (type-II)	μ e = 1.45 × 10^−2	4 × 10^3	−30 to 30	Al/Cr/Au	Exfoliated/(PVT)	Park et al., 2018^250
	μ h = 1.64 × 10^−2
	μ e = 0.36	10^3	−30 to 30		CVD/PVT	He et al., 2017^48
	μ h = 1.27
	μ e = 1–10	10^3	0 to 30	Cr/Au	Exfoliated flakes/PVT	Liu et al., 2015^1
	μ h = 0.1
2L-MoS2/pentacene (type-II)	μ e = 1.7	10^3	−80 to 60	Au	Exfoliated/TVD	Deep et al., 2016^97
	μ h = 0.004
	μ e = 0.01	10^6	50	Au	CVD/TVD	Ren et al., 2017^134
	μ h = 0.14
MoS2/tetracene (type-II)	μ e = 2.92 × 10^−5	—	−80 to 80	Au/Ti	CVD/VD	Park et al., 2018^251
	μ h = 2.48 × 10^−5
MoS2/P3HT (type-II)	μ h = 7.50 × 10^−1	—	0 to −20	ITO	Exfoliated/drop casted	Surya et al., 2017^252
MoS2/C8-BTBT (type-II)	—	10^5	20 to −40	Au	Exfoliated/MBE	He et al., 2015^109
MoS2/VOPc (type-II)	μ e ∼ 47	∼10^5	−15 to 80	Au	Exfoliated/VD	Andleeb et al., 2023^253
MoS2/CuPc (type-II)	—	10^4	−30 to 30	Ti/Au	PDMS dry transfer/TVD	S. Vélez et al., 2015^102
MoS2/CuPc (type-II)	μ e = 1.67	10^6	−40 to 40	Ti	Exfoliated/TVD	J. Pak et al., 2015,^254
1L MoS2/CuPc (type-I)	—	10^2	−60 to 60	Cr/Au	CVD/drop cast	Ganesh et al., 2018^147
MoS2/perylenediimide; (type-II)	μ e = 1	10^2	0 to 2	Ti/Au	Exfoliated/drop casting	Aday J. Molina-Mendoza et al., 2016^65
MoS2/tetraphenyl porphyrin (type-II)	μ e = 3 × 10^−3	10^3	0 to 2	Ti/Au	Exfoliated/drop casting	Aday J. Molina-Mendoza et al., 2016^65
MoS2/ZnPc (type II)	μ e = 5.4 × 10^−3	10^5	−10 to 40	Cr/Au	CVD/solution immersion	Y. Huang et al., 2018^105
MoS2/Alq3 NPs (−)	—	—	—	—	Drop-casting	G. Ghimire et al., 2017^255

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MX2:dopant (n/p dopant)	Mobility [cm^2 V^−1 s^−1]	I on/off ratio	V GS (V)	Contact and doping	Fabrication method	Ref.
MoS2:F4-TCNQ (p-dopant)	—	10^7	−50 to 20	Cr/Au	Exfoliated/CVD	J. Wang et al., 2018^114
			−20 to 50
MoS2:benzyl viologen (BV) (n-dopant)	μ e = 24.7	10^5	−2 to 2	Ni/Au	Exfoliated/solution immersion	D. Kiriya et al., 2014,^117
MoS2:butanethiol	μ e = 8.0	4 × 10^7	−80 to 80	Au	Exfoliated/vapor treatment	S. Bertolazzi et al., 2017,^206
PdSe2:F4TCNQ (p-dopant)	μ h = 12	10^2	−35 to 35	Ti/Au	Exfoliated/drop cast	W. L. Chow et al., 2017^256
MoS2:TTF (n-dopant)	—	—	—		Soaking	S. Dey et al., 2013^257
MoS2:p-toluene sulfonic acid (PTSA) (n-dopant)	μ e = 61.4	—	−45 to 25	Cr/Au	Soaking	S. Andleeb ewt al., 2015^258
MoS2:polyethyleneimine (PEI) (n-dopant)	μ h = 12.3	∼10^4	—	Ti/Au	Soaking	Y. Du et al., 2013^259
MoS2:MEA (n-dopant)	—	10^3	−40 to 40	Ti/Au	Soaking	D. M. Sim et al., 2015^260
MoS2:PPh3 (n-dopant)	μ e = 229	10^6	−50 to 50	Ti/Au	Spin-coating	K. Heo et al., 2018^261
MoS2:OTS (p-dopant)	μ h = 4.8	10^3	−40 to 0	Ti/Au	Exfoliated/soaked	D. Kang et al., 2015^262
MoS2:APTES (n-dopant)	μ e = 142.2
WSe2:OTS (p-dopant)	μ h = 168.9	10^3	10 to −30	Pt/Au	Exfoliated/soaked	D. Kang et al., 2015^262
WSe2:APTES (n-dopant)	μ h = 8.98
1L WS2:F4TCNQ (p-dopant)	μ e = 9.5	—	−30 to 35	Ni/Au	Exfoliated/drop cast	N. Peimyoo et al., 2014^66
MoS2:2-mercaptoethylamine (n-dopant)	—	10^3	−40 to 40	Ti/Au	Exfoliated/soaked	D. M. Sim et al., 2015^260
MoS2:1H,1H,2H,2H-perfluorodecanethiol (p-dopant)	—	—
MoS2:hexadecanethiol (HS(CH2)n1CH3) (n-dopant)	μ e = 35.2	10^4	−40 to 40	Ti	Exfoliated/dip coating	K. Cho et al., 2015^190
MoS2:(2-Fc-DMBI)2 (n-dopant)
MoS2:2-Fc-DMBI-H (n-dopant)	—	10^3	−50 to 200	Ti/Au	CVD/dip coating	A. Tarasov et al., 2015^263
MoS2:oleylamine (n-dopant)	μ e = 25	2.9 × 10^0	−10 to 10	Au	Exfoliated/dip coating	C. J. Lockhart de la Rosa et al., 2016^264
MoS2:OTS (p-dopant)	μ e = 3.47 ± 0.36	10^3	−60 to 60	Al	Exfoliated/dip coating	Y. Li et al., 2013^40
MoS2:APTMS (n-dopant)	μ e = 1.82 ± 0.45	10^4
MoS2:FOTS (n-type)	μ e = 0.087	10^0	−40 to 40	Ti/Au	CVD/vapor	S. Najmaei et al., 2014^22
MoS2:amine (–NH2),	μ e = 3.61 ± 0.37	10^4
MoS2:methyl (–CH3),	μ e = 2.31 ± 0.30	10^2
MoS2:fluoro (–CF3),	μ e = 2.17 ± 0.40	10^4
and MoS2:thiol (–SH)	μ e = 13.0	10^6
WS2:N2 (p-dopant)	μ e = 0.53	10^2	−20 to 20	Ti/Au	Magnetron sputtering/vapor exposed	B. Tang et al., 2018^265
MoS2:O2 (n-dopant)	μ e = 6.5	10^5		Cr/Au	Exfoliated/vapor exposed	Long Qi et al., 2016^266
WSe2:Cs2CO3	μ e = 3.9–27	∼10^7		Ti/Au	Exfoliated/evaporated	Lei et al., 2017^267
MoS2:4ATP (n-dopant)	—	10^5		Ti/Au	Exfoliated/dip coating	Im et al., 2021^268
MoS2:4NTP (p-doant)

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5.2. Unipolar FET

Doping. FETs based on monolayer or few-layer MoS₂ have demonstrated an outstanding current on/off (I_on/off) ratio of approximately 10⁸ (where I_on/off is the ratio of the on-state current and off-state current) and room-temperature mobility of > 200 cm² V⁻¹ s⁻¹, and the layered WS₂ has also been reported to exhibit I_on/off ∼ 10⁵ at room temperature modulation and bipolar behavior.^269,270 The I_on/off ratio is the figure of merit for having high performance and low leakage for the FETs. Typically, more gate control leads to more I_on/off. Due to the donor states associated with sulfur vacancies, MoS₂-based FETs are typically in an on-state without applying any gate voltage. Such device therefore requires additional negative bias to shut down, which inevitably results in power consumption even when the device is not in use. In addition, the sulfur vacancy states are always associated with poor subthreshold swing (SS) and low reliability.¹¹⁴

In order to solve this drawback, a new strategy to tune the MoS₂ FETs’ electrical behavior via TMDO heterostructure fabrication is proposed. Precise doping of 2D-TMD semiconductors can provide a robust approach to manipulate their electrical and optical properties, yielding significant improvements in device performance. Utilizing organic molecular material F₄TCNQ with strong electron-withdrawing ability is employed, whose LUMO is located at around E_LUMO = −5.3 eV.^271,272 Charge transfer TMDO system is formed by organic vacuum deposition of F₄TCNQ onto the surface of MoS₂, at which the electron transfer takes place from the MoS₂ layer to F₄TCNQ. Fig. 11(c) shows the transfer curves of linear region at V_ds = 2 V, for FETs with/without F₄TCNQ's coating. FETs of the bare MoS₂ layer show onset voltages around V_on = −30 V, while after depositing F₄TCNQ, a positive threshold voltage shift of the transfer curves takes place, and the positions of V_on are more densely distributed within the range from −10 to 0 V. This should be attributed to the charge transfer occurring at the TMDO interfaces, during which the mobile electrons distributed in the MoS₂ layer are captured by the deposited F₄TCNQ molecules and become highly localized.^186,273 The formation of the TMDO interfaces does not induce any degradation of FET mobility.²⁷⁴ Hence, with the depletion of residual electrons from sulfur vacancies, the onset voltages (V_on) of MoS₂ FETs are modulated from minus dozens of volts to around zero, together with more ideal subthreshold slope (SS) values and with no degradation of the field-effect mobility. In a recent study, Wang et al. reported on the effective functionalization, explored how the axial ligands have impact on the electronic properties of MoS₂. The back-gate, top-contact FETs with monolayer MoS₂ as a channel material were fabricated, as portrayed in Fig. 11(d).⁶⁹ After the formation of self-assembled metal phthalocyanine (MPcs), a large shift of the transfer curve (I_ds–V_g) to the positive direction is observed (Fig. 11e and f). This observed evidence can be attributed to the formation of p–n junctions.^150,254 By further coordinating the FR Pys, shifts in two opposite directions are observed: for 3,5-FP and 3-FP, the I_ds–V_gs curves shift to positive voltages while for 2-FP and 4-AP, the curves shift to the negative direction.⁶⁹ For MoS₂/ZnPc, ligand functionalization enables the modulation of the carrier density where C_ox is the capacitance per unit area, ε_ox is the dielectric constant, ΔV_th is the change of threshold voltage, and t_ox is the thickness of SiO₂. In MoS₂, the modulation ranges from −3.0 × 10¹² cm⁻² in 3-FP to 1.8 × 10¹² cm⁻² in 4-AP, while for MoS₂/CoPc, the modulation ranges from −3.5 × 10¹² cm⁻² in 3-FP to 5.32 × 10¹¹ cm⁻² in 4-AP, with respect to the reference case of MoS₂/MPc. The magnitude of the changes is surprisingly high, being comparable to MoS₂ doping with ZnPc and CoPc, which amounts to 1.8 × 10¹² and 2.1 × 10¹² cm⁻², respectively. Considering that the major charge carriers in MoS₂ are electrons, it is concluded that 3,5-FP and 3-FP decrease the electron density in MoS₂/MPc, inducing p-doping, whereas 2-FP and 4-AP increase the electron density, inducing n-doping.⁶⁹

Fig. 11 (a) Transfer characteristics (I_SDvs. V_G) of a pristine MoS₂ FET before treatment and after consecutive n-doping with RhCp*Cp with three different concentrations/exposure times. (b) I_SDvs. V_G of a pristine MoS₂ FET and after p-doping with MB for increasing exposure times. (c)–(e) I_SDvs. V_G of the remaining three MX₂ FETs (MoSe₂, WS₂, and WSe₂) before and after doping with n- and p-dopants. Treating the transistor with n-dopants shifts the threshold voltage (V_th) to more negative values, indicating n-doping (arrow). (f) Effects of different n- and p-dopants on V_th. (g) Schematic of monolayer MX₂ transistor (M = Mo, W; X = S, Se). (a)–(g) Reproduced from ref. 204 with permission from Wiley-VCH, copyright 2018.

The implementation of ion beam irradiation has been recognized as an potential method to engineer chemically active defects in 2D TMD materials.²⁰⁶ In this context, Bertolazzi et al. employed the argon-ion bombardment method to create sulfur vacancies in monolayer MoS₂.²⁰⁶ However, a deep understanding of the effects of generated defects on the functional characteristics of MoS₂ is still needed. A systematic investigation of the correlation between critical electronic device parameters and the density of sulfur vacancies is reported through the fabrication and characterization of a 1L MoS₂ FET that has been exposed to a variable fluence of low-energy argon ions. Exposing both pristine and ion-irradiated FETs to vapors of butanethiol—a short linear thiolated molecule can further improve or restore their electrical properties.^{190,206,208,275} This solvent-free chemical treatment, investigated strictly in an inert atmosphere, eliminates secondary healing effects generated by oxygen or oxygen-containing molecules. The outcome provides valuable insights for designing optoelectronic devices based on MoS₂ with a controlled density of sulfur vacancies. Furthermore, these results pave the way for introducing specific molecular functionalities using effective thiol chemistry approaches.

Another alternative method has been reported for the substitution of p-type doping in 1L to few-layer WS₂ by introducing highly reactive nitrogen atoms.²⁶⁵ It is demonstrated that the introduction of nitrogen atoms in atomically thin WS₂ induces minimal lattice distortion owing to its low kinetic energy. The separation of the A_1g and ¹E_2g Raman peaks remains the same as that for the pristine sample, which indicates that no obvious lattice distortion was induced during the doping process. However, a noticeable peak broadening in the ¹E_2g phonon is found from the N-WS₂ sample, which is attributed to symmetry weakening of the ¹E_2g phonon mode caused by the nitrogen-induced strain in WS₂. The transfer characteristics of WS₂ FETs clearly exhibit an n-channel-to-p-channel conversion with nitrogen incorporation. Tang et al. further explored the defect formation energy and the origin of p-type conduction by first-principles calculations.¹⁹⁸ In nitrogen-doped WS₂, defect state emerges near the Fermi level, resulting in a shallow acceptor level located above the valence band maximum. By employing this doping strategy, it becomes possible to achieve substitutional p-type doping in intrinsically n-type TMDs. The process allows for the precise control of dopants, presenting an avenue for creating complementary metal-oxide-semiconductor FETs and optoelectronic devices on 2D TMDs. This approach overcomes one of the main limitations of TMDs, namely, their n-type unipolar conduction.²⁶⁵ Kappera et al. introduced a novel phase engineering approach to achieve low-resistance contacts and enhance the performance of FETs based on ultrathin MoS₂.²⁷⁶ They have further established a methodology to fabricate metallic 1T phase patterns on ultrathin semiconducting MoS₂ nanosheets and used them as electrodes. The low contact resistances of 200–300 Ω μm at zero gate bias are attributed to the atomically sharp interface between the phases and the similarity in work function of the 1T phase and the conduction band energy relative to the vacuum level of the 2H phase (approximately 4.2 eV) (Fig. 12h–i). Both top-gated and bottom-gated devices employing 1T phase electrodes exhibit superior performance when compared to devices with direct metal contacts to the semiconducting 2H phase. The measured properties in air indicate mobility values of approximately ∼50 cm² V⁻¹ s⁻¹, subthreshold swing (SS) values of 90–100 mV decade⁻¹, and current on/off ratios of >10⁷.²⁷⁶ Additionally, the performance of devices with 1T electrodes is remarkably reproducible and remains independent of the type of contact metal pads used.

Fig. 12 (a) Schematic of the MoS₂ FET with F₄TCNQ-coated nanoparticles. (b) Energy band diagram at the vdW interface between MoS₂ and F₄TCNQ. (c) I_DSvs. V_GS of MoS₂ FETs with and without F₄TCNQ. (d) Device schematic of MoS₂ functionalized with MPc/fPy. (e) Comparison of transfer curves of MoS₂/ZnPc heterojunction coordinated with ligands. Comparison of I_DSvs. V_GS of MoS₂/CoPc heterojunction coordinated with ligands. (f) Comparison of transfer curves of MoS₂/CoPc heterojunction coordinated with ligands. (g) Schematic of 1T phase electrodes and Au electrode on the 2H phase. (h) and (i) Transfer characteristics with 1T and 2H contacts, respectively. (a)–(c) Reproduced from ref. 114 with permission from Wiley-VCH, copyright 2018. (d)–(f) Reproduced from ref. 69 with permission from Wiley-VCH, copyright 2020. (h)–(i) Reproduced from ref. 276 with permission from Nature Publishing Group, copyright 2014.

5.3. Anti-ambipolar FET

An atomically thin TMD partially screened electric fields when applied perpendicularly to the monolayer, leading to the development of ‘gate-tunable’ TMDO p/n heterojunction diodes for potential use in high-speed communications circuits, for example, pentacene/MoS₂ and CuPc/MoS₂.^97,102 Furthermore, applied voltages can be employed to control extended defects like grain boundaries in monolayer TMDs.

As a result, this leads to the production of gate-tunable memristor devices, showing great promise as fundamental building blocks for ‘non-volatile’ computer memory and neuromorphic (brain-like) circuit architectures.²⁷⁷ Due to the atomically thin nature of MoS₂, the junction displays gate-tunability with an asymmetric anti-ambipolar transfer characteristic.¹³³ Through the utilization of spatial photocurrent mapping and surface potential profiles, along with finite element simulations, it has been confirmed that there is band bending occurring in both the lateral and vertical directions in type-II band alignment. The type-II alignment induces a built-in potential, resulting in the rectification of I_DS–V_DS (output) characteristics at different V_G values. In the pentacene/MoS₂ heterojunction, V_DS represents the bias applied to the pentacene electrode (electrode 3 in Fig. 13a), where V_DS > 0 corresponds to the forward bias, while the MoS₂ electrode is grounded. The heterojunction exhibits a transition from nearly insulating behavior at extreme V_GS values (V_GS = 60 V and V_GS = −80 V) to highly rectifying output characteristics at intermediate values. The transfer plots depicted in Fig. 13c provide additional evidence of the gate-tunability of the current through the p–n heterojunction. Notably, the transfer curve of the heterojunction (green) exhibits an anti-ambipolar response, which is also observed in inorganic p–n junctions.^278–280 In contrast to earlier devices featuring symmetric transfer characteristics, the pentacene/MoS₂ heterojunction demonstrates asymmetric transfer characteristics with different transconductances on either side of the I_DS peak. The observed asymmetry in the anti-ambipolar transfer characteristic presents promising opportunities in emerging circuit applications. It has the potential to be effectively utilized for simultaneous phase and amplitude shift keying in wireless telecommunication technologies.^97,281 Velez and their group presented experimental findings on a CuPc/MoS₂ hybrid p–n junction, revealing its gate-modulate diode characteristics and photovoltaic effects.¹⁰² In Fig. 13f, the transfer curves of the CuPc/MoS₂ device are displayed at different V_ds values. The transfer curve exhibits a characteristic shape that remains consistent for all V_ds values, demonstrating an asymmetry originating from the small subthreshold swing.²⁶⁹ The p–n junction devices based on monolayer MoS₂ yield similar curves. Remarkably, efficient gate control over the heterostructure is consistently achieved. This capability is directly associated with the use of thin MoS₂ flakes, whether monolayers or bilayers, which facilitates the extension of the electric field from the gate into the CuPc layer.

Fig. 13 (a) Schematic of the pentacene/MoS₂ p–n heterojunction, and optical and AFM micrographs of a representative device. Electrical properties of the pentacene/MoS₂ anti-ambipolar p–n heterojunctions. (b) Gate-dependent I_DS–V_DS output characteristics of a representative device. (c) I_DSvs. V_GS of the pentacene (V_DS = 10 V) and MoS₂ (V_DS = 1 V) FETs as well as the junction (V_DS = 10 V). The junction transfer curve shows an asymmetric anti-ambipolar characteristic. (d) Simulated (dashed lines) and experimental (solid lines) transfer characteristics of pentacene and MoS₂ FETs. (e) p/n junction CuPc/MoS₂ device. (f) Transfer curves (I_dsvs. V_g) measured in CuPc/MoS₂ FET at V_ds = 16 V, in the MoS₂ bilayer of the same device at V_ds = 5 V and in a CuPc transistor produced on the same wafer at V_ds = 16 V. (g) I_dsvs. V_g of the CuPc/MoS₂ device measured at different V_ds. (a)–(d) Reproduced from ref. 97 with permission from American Chemical Society, copyright 2022. (e)–(g) Reproduced from ref. 102 with permission from Royal Society of Chemistry, copyright 2015.

5.4. Phototransistors

The operation of a phototransistor relies on the photogating effect, where a gate bias is applied to modulate current. This modulation process resembles that of a conventional FET structure. However, phototransistors differ in their ability to effectively detect electrical signals resulting from alterations in source-drain current or shifts in the transfer curve due to gate bias changes when exposed to light.^105,250,282 Evaluating the performance of phototransistors in light sensing applications involves considering several crucial parameters. These parameters encompass photoresponsivity (R), photosensitivity (P), detectivity (D), and external quantum efficiency (EQE), each playing pivotal roles in assessing their effectiveness. For a more in-depth examination of these parameters and their significance, several articles can be referred.^283,284 In particular, photoresponsivity (R) stands out as a key metric for gauging phototransistor performance, and its determination relies on the analysis of transfer characteristics. The formula for calculating R is provided by eqn (4):where P_light represents the overall incident optical power. Another significant factor in characterizing the detector's performance is photodetectivity (D*), which is computed using eqn (5):where A is the channel area and q is the absolute value of the electron charge.²³ Summarizing eqn (2) and eqn (3) gives the EQE and it is often expressed by the device responsivity (R; in units of A Watts⁻¹), which is the ratio of the photocurrent (I_light) to the incident light intensity (P_light)where q is the elementary charge and hv is the photon energy.

The visible range of the direct bandgap offers effective light absorption and luminescence capabilities, which allow a MoS₂-like TMD gap semiconductor to be used as a channel material for phototransistors. However, in atomically thin semiconductors, the total photocarrier generation efficacy is limited by the level of optical absorption. When organic semiconductor is combined with layered TMD semiconductor, the monolayer TMD typically led to a higher optical absorption than any TMD heterostructures.^17,23,144 TMD-based phototransistors can suffer from persistent photoconductance, failing to turn off when the light is removed due to charge trapping at the substrate/TMD interface and inherent trap states in thin materials. To address this, molecular systems forming type II alignments with TMDs have been employed. These systems quickly dissociate photoexcited charges through the built-in potential at the type II interface, enabling a swift photoresponse.^254,267,285

TMDO hybrids outperform pristine TMDs by promoting spontaneous charge transfer across layers, enhancing electrical currents. Additionally, carefully chosen organic systems can absorb a broad spectrum of light, enabling TMDO phototransistors to operate over a wide range of wavelengths. The combination of 2D-TMDs with organic semiconducting crystals in heterostructures offers exciting possibilities for the fabrication of diverse hybrid devices such as phototransistors.^282,285,286 It has been recently reported that thiol-based modification enhances electrical and optical performance in multilayer MoS₂. Two types of thiol-terminated organic molecules, namely, 4-amino thiophenol (4ATP) and 4-nitro thiophenol (4NTP) are used for electron-donating and electron-withdrawing effects, respectively.²⁶⁸ These molecules chemically bond with MoS₂, with 4ATP boosting the current and carrier concentration and 4NTP reducing the film conductance. In Fig. 14b, the 4ATP-MoS₂ transistor exhibits a higher drain current (I_DS) than that of the pristine case, and the threshold voltage (V_TH) shifts negatively from −19.6 V to −49.9 V after 24 hours. Further, Fig. 14c shows time-resolved photoswitching I_DS curves for pristine and 4ATP-MoS₂ transistors (V_GS = 0 V). The photoswitching characteristics are suitable, with increased I_ph due to enhanced I₀ after functionalization. Additionally, they modulate MoS₂ transistor photoresponsiveness, with 4ATP enhancing photoresponsivity (≈669 A W⁻¹) and 4NTP improving photoswitching.²⁶⁸ Thiol-based functionalization effectively tailors multilayer MoS₂ electrical and optical properties. Besides functionalization, CuPc molecules boosted the MoS₂ performance in CuPc/MoS₂ hybrid structures, increasing the photoresponsivity and photodetectivity tenfold. The hybrid achieved an EQE of 12%, over twice that of pristine MoS₂, demonstrating superior performance in the visible spectrum. Molecular systems enhance the quantum efficiency and speed up photoresponse in TMD phototransistors.²⁵⁴

Fig. 14 Phototransistor: (a) device structure (inset: details of electron transfer at the interface). (b) Transfer curves. (c) Time-resolved photoswitching I_DS curves. Photo-switchable superlattice: (d) schematic representation of the molecular dipoles (black arrows) are randomly oriented before irradiation, yet well aligned after UV irradiation. (e) STM images of the merocyanine (MC) assembly on MoS₂. (f) I_DSvs. V_GS curve: trace 1 (black) for clean MoS₂, trace 2 (red) for MoS₂ covered by the SP layer, trace 3 (blue) for MoS₂/MC after UV irradiation over the whole flake. (g) Trace 4 (green) for the pristine MoS₂/SP trace can be recovered upon irradiation of the whole surface with green light. (a)–(c) Reproduced from ref. 268 with permission from Wiley-VCH, copyright 2021. (d)–(g) Reproduced from ref. 62 with permission from Nature Publishing Group, copyright 2018.

A recent investigation has demonstrated the development of highly sensitive light detection phototransistors based on a fully ordered organic SAM incorporated into hybrid organic–inorganic van der Waals (vdW) heterojunctions.²⁸⁷ Incorporating a SAM as the bottom layer offers many advantages: (i) the SAM allows precise control over the long-range ordered molecular film, achieving a single molecular thickness,²⁸⁸ effectively minimizing charge scattering and recombination as compared to multilayer organic films;²⁸⁹ (ii) the processing method involving adsorption from a solution is straightforward;²² and (iii) BTBT-SAM enables efficient two-dimensional lateral π–π interactions for effective charge transfer.²⁸⁸ The heterojunctions, formed by monolayer MoS₂ onto organic BTBT SAM, are studied by Raman and PL spectroscopy along with Kelvin probe force microscopy. Remarkably, this heterojunction transistor displays an exceptional photoresponsivity reaching 475 A W⁻¹ and an enhanced external quantum efficiency of 1.45 × 10⁵%. Additionally, it exhibits an exceedingly low dark photocurrent in the pA range.²⁸⁷ It is further demonstrated that the hybridization of SAMs with TMD materials holds great promise for creating a wide range of optoelectronic devices with distinctive properties. Phototransistors offer numerous design possibilities for stacking and integrating 2D TMD atomic-molecular layers by manipulating the SAM structures. This novel heterostructure shows exciting potential for advanced optoelectronic devices.

5.5. Photo-switchable superlattices

Molecular switches facilitate the development of multifunctional devices, wherein an electrical output can be influenced by external stimuli. However, proving the operational mechanism of these devices often poses challenges, as the molecular switching events are indirectly confirmed through electrical characterization without direct real-space molecular-scale visualization. In a recent study, Gobbi and co-workers⁶² demonstrated multi-responsive devices that rely on the collective switching action of photochromic molecules self-assembled on the surface of TMDs. They unveiled the mechanism by introducing a method to optically modulate the electrical properties of versatile hybrid devices, utilizing photochromic SP and MoS₂ superlattices. They achieved this by creating supramolecular assemblies of photochromic spyropyrans (SP) on the surfaces of monolayer MoS₂, enabling the investigation of light-induced changes in physicochemical properties on a macroscopic scale (see Fig. 14e). After inducing SP/merocyanine (MC) isomerization with UV light, they observed a significant reduction in the work function of the underlying materials. Conversely, exposure to visible light triggered MC/SP isomerization, nearly fully restoring the work function of the material. This demonstrated their ability to exercise optical control over the local charge carrier density in high-performance transistors through a comprehensive study that integrated sub-nanometer-resolved STM investigation and characterization of mesoscopic devices (as shown in Fig. 14e). Their STM imaging findings aligned well with the results obtained from molecular dynamics simulations and density functional theory calculations. They successfully showcased the ability to adjust the threshold voltage in MoS₂ transistors by exposing them to UV-visible light at the heterojunctions between inorganic surfaces and 2D supramolecular networks (as depicted in Fig. 14f and g).⁶² This work's notable achievement lies in the direct visualization of structural changes at the molecular level accompanying photoisomerization. By precisely manipulating the electrical characteristics of 2D vdW material devices through ultra-high spatial resolution molecular switching, they have created opportunities for tailored adjustments in the electrical output of organic–inorganic heterojunction devices, offering promising prospects for multifunctional sensors and optoelectronic applications.

A similar level of control over switching events and molecular assembly has been previously demonstrated in dynamic molecular crystals. These dynamic molecular crystals refer to self-assembled crystalline structures, usually three-dimensional, where macroscopic structural changes originate from collective molecular events.^290–292 However, bulk dynamic molecular crystals are generally poor electrical conductors and unsuitable for macroscopic device operation. This limitation arises due to the chemical structure of molecular switches, which is not optimized for efficient charge transport. In this regard, by adopting this approach, it becomes feasible to incorporate a functional supramolecular assembly, representing the quasi 2D limit of molecular dynamic crystals, into high-performance devices capable of exhibiting switchable electric outputs.

Moreover, future optical, electrical, and quantum information applications revolve around newly emerging quantum phenomena in electronically coupled two-dimensional heterostructures. Such processes could be controlled by tailoring the electronic band gaps in coupled heterostructures, which is of great research interest. vdW heterostructures made of atomically thin materials of TMDs have emerged as another promising platform for studying correlated electronic phenomena with highly tunable parameters.²⁹³ These layered materials can be assembled into non-covalently bonded heterostructures without the lattice matching results electronic interaction is maximization. Such a superposition of the TMD layer often gives rise to an interference effect known as the moiré pattern.^294,295 Such regions in the moiré pattern appear dark in the STM images due to the reduced density of states near the Fermi level. This pattern defines a superlattice, which has a periodicity larger than that of either of the individual lattices. This in-plane superlattice induces local variations in the heterostructure's electronic structures, leading to the formation of minibands.^296,297 This effectively enhances the effective mass of the charge carriers and makes correlation effects more prominent in such vdW heterostructures.²⁹⁸ In this regard, 2D TMD represents an ideal platform to study the interplay between the molecular assembly on surfaces and electrical transport in devices. On the exposed surface of 2DMs, well-defined molecular groups can be arranged at predetermined spatial locations with atomic precision by tailoring of molecular architectures.⁶²

Single-layer porphyrin polymerization. The development of superlattices with periodic TMD material components and defined functions is possible with continuing structure and interface engineering. All these qualities might be used in the future to create advanced optoelectronic, neuromorphic, and other multifunctional devices that are directly interfaced with soft robotics or biological systems for communication, computation, prosthetics, bio-sensing, and personal healthcare. Further, Zhong et al.²⁹⁹ synthesized monolayer 2DPs on a wafer scale and integrated them into hybrid heterostructures with precise monolayer control. They developed a versatile method called laminar assembly polymerization (LAP), compatible with various molecular building blocks and two primary polymerization chemistries (coordination and covalent bonding). LAP allows large-area synthesis, operates under ambient conditions, and is compatible with established patterning and integration techniques. This method enabled the creation of superlattices combining monolayer 2DPs with 2D atomic crystals. Their design approach used porphyrin building blocks with two variable sites: one at the center of the porphyrin ring (M = 2H, Fe(III), or Pt(II)), affecting optical properties, and the other on the phenyl groups (R = –COOH or –NH₂), controlling monomer-monomer bonds. They cross-linked the monomers using either coordination bonds with copper paddle wheel structures in the presence of Cu²⁺ ions or covalent bonds via the Schiff base reaction with terephthalaldehyde (TPA). This resulted in coordination 2DPs (also known as monolayer metal–organic frameworks, MOFs) and covalent 2DPs (also known as monolayer covalent organic frameworks, COFs). The linkage chemistry was confirmed by Fourier-transform infrared (FTIR) spectroscopy.²⁹⁹

5.6. A recent advent: multifunctional synaptic FET

Biological synapses are integral components of the nervous system, serving as essential elements for computing and memory functions in the human brain. In artificial neuromorphic systems, synaptic devices play a pivotal role as fundamental building blocks for constructing neuromorphic networks (see schematic Fig. 15a–c). Therefore, the advancement of novel artificial synapses holds paramount importance in the field of neuromorphic computing.^300,301 In a nervous network, the synaptic weight defined as the connection strength between presynaptic neuron and postsynaptic neuron can be modulated by action potentials, which have been well discussed in a recent review.^302–304 Synaptic plasticity, a fundamental neurobiological phenomenon, is crucial for memory and learning processes in the brain, enabling the implementation of intricate brain functions. It encompasses two key components: short-term plasticity (STP) and long-term plasticity (LTP). These mechanisms play vital roles in shaping the efficiency and adaptability of neural circuits.³⁰⁵ STP stored in the hippocampus is characterized by a temporal change in synaptic connections, which rapidly decays to its original state after the removal of an external spike. STP is essential for efficient information processing in neural circuits.

Fig. 15 TMDO hybrid heterojunctions constituting a multifunctional artificial synaptic transistor. (a) Biological synaptic diagram: electrical or optical spikes serve as a presynaptic input spike and a source–drain channel current as stimulated postsynaptic current. (b) Simplified illustration of a biological synapse. (c) TMDO hybrid heterojunction constituting a multifunctional artificial synaptic transistor.

When a potential pulse is applied to the presynaptic membrane, it triggers the release of neurotransmitters, resulting in a change in the postsynaptic current. Excitatory postsynaptic current (EPSC) occurs when excitatory neurotransmitters are released and results in the strengthening of the synaptic weight. However, inhibitory postsynaptic current (IPSC) occurs when prohibiting neurotransmitters are released, and results in the weakening of synaptic weight. Paired-pulse facilitation (PPF) is one of the typical STP behaviors, which describes the fact that an enhanced postsynaptic response is obtained by the consecutively applied pair of presynaptic spikes. In contrast, LTP represents a longstanding transformation in synaptic connections, playing a crucial role in memory and learning processes. Through repetitive rehearsal, STP can be converted into LTP.^306,307 Further, simulating synaptic short-term plasticity (STP) and long-term potentiation (LTP) behaviors using unique artificial devices, like inhibitory post-synaptic current (IPSC), excitatory post-synaptic current (EPSC), paired-pulse depression (PPD), paired pulse facilitation (PPF), and spiking-rate-dependent plasticity (SRDP), holds great promise. These devices play a crucial role in advancing the development of neuromorphic computation, enabling the study and implementation of complex synaptic responses and facilitating memory and learning processes akin to those observed in biological systems. In addition, LTP in synaptic devices provides a physiological substrate for learning and memory, whereas STP supports a variety of computations. It is worth noting that organic bioelectronics have gained intense interest in neural applications due to their low-cost processing, availability of a virtually infinite number of molecules with multifunctional optoelectronic properties, tenability and mechanical flexibility. Furthermore, because of their unique internal and interfacial structures, and electrical and optical properties,^308,309 2D-TMDs such as MoS₂ and WSe₂ have been reported as potential candidates for complex neuromorphic applications. Recently, Wang and coworkers have demonstrated the first multifunctional artificial neural synapse transistor fully based on 2D-TMDs and organic layers, that is, MoS₂/PTCDA hybrid heterostructures with neuromorphic functions of STP and LTP including IPSC, EPSC, PPD, PPF, and SRDP, which can mimic some crucial behaviors of biological neurons.³¹⁰

By carefully designing the energy band structure, the system achieves both robust electrical and optical modulation modes. The gate-tunable band alignment of such a type-II heterojunction allows the carriers to be both electrically and optically injected into MoS₂, separated and captured at the heterojunction interface and gradually released to simulate the basic functions of biological synapses. They have reported the synaptic inhibition and excitation simultaneously by efficient modulation of the gate voltage in a single device. It exhibits a threshold voltage (V_th) drift window of 6 V in a static direct-current transfer curve. Notably, it exhibits a minimum inhibition of 3% and a maximum facilitation ratio of 500% by increasing the number of gate electrical pulses. Additionally, the device's response to varying frequency voltage signals demonstrates dynamic filtering capabilities, making it an efficient means for information transmission within an artificial neural network during electrical modulation. In the realm of optical modulation, this device demonstrates remarkable flexibility in tuning short-term potentiation (STP) and long-term potentiation (LTP). The long-term synaptic weight change (ΔW/W) achieved in this mode is notably improved, reaching up to 60, surpassing the results reported in other works. In the optical modulation mode, a very good flexible tunability of STP and LTP is reported and the long-term synaptic weight change (ΔW/W) is greatly improved, reaching up to 60, compared to the results reported in other works.^310–313 Moreover, upon comparing the optimal performance parameters of MoS₂/PTCDA devices, they exhibit remarkable superiority when compared to other artificial synaptic devices based on 2D materials, whether they are organic or inorganic in nature.^{122,271,313–327} The acceptable power consumption of 10 pJ is 90 times smaller than the conventional CMOS and close to the human brain.^316,328 Therefore, the multifunctional artificial synaptic transistor based on the TMDO hybrid heterojunction is a promising candidate for constructing a neuromorphic system. It paves a whole new and efficient path for building neuromorphic system hardware elements.

They observed synaptic behavior by applying a presynaptic V_cg voltage pulse while maintaining V_bg constant (Fig. 16a and 2b). Prior to the V_cg pulse, I_ds of the MoS₂/PTCDA synaptic transistor remained stable over time, corresponding to stage I. During the 25 ms V_cg pulse, I_ds underwent modulation via electrostatic doping. Depending on the pulse (positive or negative), I_ds was either suppressed (Fig. 16a) or enhanced (Fig. 16b), corresponding to stages II and IV, respectively. Importantly, after the pulse cessation, I_ds did not return monotonically to the reference level. Instead, following a relatively positive (negative) V_cg voltage pulse, it settled at a lower (higher) level relative to the reference, resembling postsynaptic current behavior with a slow relaxation phenomenon. This corresponds to stages III and V, manifesting as synaptic IPSC and EPSC behaviors, respectively. As shown in Fig. 16c, the positive V_cg pulses (−12 V, 25 ms) applied at Δt = 25 ms lead to a further suppression of the postsynaptic current magnitude, resulting in A₂ < A₁, as indicated in the gray portion. Conversely, Fig. 16d shows that negative V_cg pulses (−20 V, 25 ms) enhance the postsynaptic current magnitude (A₂ > A₁), corresponding to synaptic PPD and PPF behaviors, respectively. Further, the synaptic transistors based on MoS₂/PTCDA hybrid heterojunctions can also operate in the optical modulation mode due to the strong light absorption (served as input gate signal) capability of both MoS₂ and PTCDA. A characteristic EPSC response is observed when the MoS₂/PTCDA transistor is subjected to a pair of laser pulses. The laser pulses have a pulse width of 400 ms with an interval of 100 ms, and V_ds is set at 0.1 V (the green area represents laser pulse irradiation). This results in a higher excitation current amplitude (A₂ > A₁), as shown in Fig. 16e. The inset depicts the phenomenon that the optically modulated PPF ratio (A₂/A₁) decreases as a function of Δt. It shows a decrease in the PPF ratio with increasing Δt, eventually stabilizing around 100%. The maximum value of the PPF ratio is achieved at the minimum time interval Δt of 25 ms, reaching a value of 147%. The artificial synaptic behavior induced by the transfer and exchange of internal charge carriers in the heterojunction under electrical and optical modulation indicating a stronger synaptic plasticity (STP and LTP) is promising. The observed phenomenon exhibits striking similarity to the process of neurotransmitter release observed in biological synapses.

Fig. 16 Artificial inhibitory and excitatory synaptic behaviors based on the MoS₂/PTCDA hybrid heterojunction in the electrical modulation mode. (a) I_DS of the MoS₂/PTCDA synaptic transistor exhibits time-dependent behavior (bottom panel) in response to a relatively positive V_cg pulse (top panel). The response can be classified into three distinct stages (I, II, and III). (b) Upon the application of a relatively negative V_cg pulse, a series of EPSC behaviors can be divided by distinct stages labeled as I, IV, and V. (c) IPSC behavior when stimulated by a pair of relatively positive V_cg pulses, with pulse intervals of 25 ms. (d) EPSC behavior when subjected to a pair of V_cg pulses, with a pulse interval of 25 ms. This phenomenon enables an optical transition from STP to LTP, facilitating the long-term modulation of synaptic weight in the device. (e) EPSC of the MoS₂/PTCDA transistor is triggered by a pair of laser pulses. Pulse width = 400 ms, interval = 100 ms, and V_ds = 0.1 V; the green area represents the irradiation of the laser pulses. The inset shows the PPF index A₂/A₁ (black hollow square) as a function of laser pulse. (f) Statistical comparison of the synaptic weight changes and PPF ratio with other reported works. (a)–(f) Reproduced from ref. 138 with permission from Wiley-VCH, copyright 2018.

In neuroscience and computer science, synaptic weight represents the strength or amplitude of the connection between two nodes, elucidating the extent to the stimulation effect of a specific neuron on other neurons within biological systems.³²⁹ Changes in synaptic weight are recognized as synaptic plasticity;^330,331 however, the underlying mechanism responsible for long-term alterations remains uncertain. The long-term synaptic weight change is defined as ΔW/W = (I − I₀)/I₀, where I and I₀ are the steady-state base postsynaptic currents before and after laser pulses, respectively. With the increase in the number of laser pulses, synaptic weight changes can be effectively increased, and long-term synaptic weight changes are able to further strengthen via varying V_bg, which is consistent with the transition from STP to LTP. In other words, enhancing the synaptic connections is consistent with the long-term memory effects of the biological nervous system. Remarkably, a significant long-term synaptic weight change of up to 60 has been achieved.¹³⁸Fig. 16f exhibits the statistics and comparison of synaptic weight changes and PPF ratios with various heterostructures. Additionally, comparing the optimal performance parameters of novel MoS₂/PTCDA heterostructure devices reported superiority to that of the other 2D-based artificial synaptic devices. Therefore, it is possible to fabricate new, novel, and promising synapse element-based TMDO hybrid heterostructure transistors. Again, TMDO-based devices with superior and flexible versatility are used as artificial neural network chip units, which is of great significance for constructing artificial neural network systems. Further, the device has a robust dual electrical and optical modulation mode of operation, making it highly desirable for neural network applications, thereby overcoming challenges related to implementing these functionalities within a single device. In summary, based on the powerful dual mode of operation, the system operation of synaptic STP and LTP in the device could be achieved. Therefore, TMDO demonstrates that it has tremendous potential for application in neuromorphic computation.

6. Photodetector

For photodetection, the layered 2D-TMD-based photodetectors have been demonstrated with very high responsivity and fast photoresponse.^332–334 Beyond the structures of MoS₂ FET alone, it is reported that through stacking of other materials, MoS₂-based heterostructures can exhibit a higher photoresponsivity.^{23,254,335,336} The development of TMDO heterostructures is required to lower the recombination rate of carriers and to reduce the metal–semiconductor contact resistance, which should lead to an improved performance of TMD-based electronic and optoelectronic devices such as photodetectors. The challenging issue for MoS₂ devices exists due to their persistent photoconductance (PPC),³³³ which has been debatably ascribed to the minority carrier trapping by surface absorbents, inherent trap states in MoS₂, and the neighboring dielectrics.^337,338 Although such minority carrier trapping can improve the gain mechanism in detectors by lengthening the apparent carrier lifetime, a dramatic sacrifice in the response speed occurs in devices since the PPC effect often lasts from minutes to hours. Attempts to improve the response dynamics, e.g., using surface passivation and field-effect methods that modulate the density of trap states and their occupation, however, often leads to the loss of sensitivity due to the adverse suppression of trap-induced gain mechanisms.¹⁰⁵

In present, the ultrafast photoresponse dynamics was achieved in monolayer MoS₂ detectors by the surface assembly of ZnPc organic molecules.¹⁰⁵ It is found that the assembly of ZnPc molecules tends to compensate the intrinsic electron doping in MoS₂ by withdrawing electrons from MoS₂. Such charge transfer contributes to a spontaneous reverse separation of electron–hole pairs under light illumination, driving holes to ZnPc molecules. This is found to favorably suppress the rather slow minority carrier trapping to the inherent trap states in MoS₂ and the substrates, providing an accelerated response speed (<8 ms) compared to that of bare MoS₂ (20–40 s). The fast photoresponse dynamics here is further demonstrated to endure tailorable engineering of the device responsivity (0.92–7.84 and further to 880 A W⁻¹) when using external gate modulation via Al₂O₃ passivation, manifesting the special promise of vdW-coupled molecules on 2D TMDs for high-performance photodetection.

The field-effect transfer curves for the same MoS₂ device before and after varied ZnPc treatments (10, 40 min) were first measured, as shown in Fig. 17a–d. The measurements were conducted both in the darkness and under illumination (3.64 mW cm⁻²) using a V_ds value of 1.6 V. As displayed in Fig. 17b, the measured transfer curves for all devices manifest the n-type conduction with electrons as the majority carriers. Increasing the ZnPc treatment leads to a decrease in the source–drain current (I_ds) and a positive shift of the threshold voltage (V_th), which are consistent with the electron compensation effect of ZnPc molecules to MoS₂. After 40 min of ZnPc treatment, one witnesses the emergence of bipolar characteristics (p-type at V_g < 20 V, n-type at V_g > 5 V) in MoS₂ under dark conditions. However, under light illumination, the samples again display n-type conduction with considerably negatively shifted threshold voltages of less than −15 V. This phenomenon suggests the intense n-type photodoping effect in MoS₂ by the preferential trapping of photogenerated holes compared to electrons.^262,333,338Fig. 17c shows the dark and photo current–voltage characteristics for both the bare and ZnPc-treated MoS₂ devices. A transition of current–voltage characteristic from nonlinear to linear behavior is found when the devices are switched from darkness to illumination. Such a change is related to the reduced Schottky barrier heights (SBHs) at the contacts by the significantly increased electron concentration in MoS₂ due to the pronounced photodoping.^21,105,339

Fig. 17 (a) Schematic of a FET device with a flat-lying ZnPc molecule on 1L MoS₂. (b) I_DS–V_G characteristics for a MoS₂ device after varied ZnPc treatments in the darkness and under illumination (532 nm, 3.64 mW cm⁻²). (c) Dark and photo current–voltage characteristics for both the bare and ZnPc-treated MoS₂ devices. (d) Transfer curves and transient photoresponse characteristic of the device before and after Al₂O₃ passivation. (e) Schematic of the CuPc/MoS₂ device illuminated by a laser. The molecular structures of the CuPc are also indicated. (f) I_DS–V_G curves on the semilogarithmic scale of the MoS₂ devices without (labeled with “w/o CuPc”) and with CuPc layers of various thicknesses (1, 2, 3, 4, 5, and 10 nm) measured at a fixed V_DS = 0.1 V. The inset is a linear plot of I_DS–V_G curves for a pristine MoS₂ device without a CuPc layer and a 10 nm-thick CuPc/MoS₂ hybrid device. Arrows indicate shifts in threshold voltages and currents. (g) Photocurrent of pristine MoS₂ and CuPc/MoS₂ photodetectors measured at V_G = −40 V and V_DS = 0.1 V as a function of the laser intensity. (h) I_DS–V_G curves of pristine MoS₂ and CuPc/MoS₂ photodetectors measured in the darkness and under illumination conditions at different wavelengths and a fixed laser power (40 mW) (adapted from ref. 187). (a)–(d) Reproduced from ref. 105 with permission from American Chemical Society, copyright 2018. (d)–(g) Reproduced from ref. 254 with permission from Royal Society of Chemistry, copyright 2015.

Furthermore, an additional study demonstrates favorable photo-responsive characteristics of MoS₂ FET by incorporating a p-type organic copper phthalocyanine (CuPc) layer on the MoS₂ surface (see Fig. 17e–h).²⁵⁴ Specifically, the CuPc/MoS₂ hybrid devices demonstrated superior photodetection capabilities when compared to pristine MoS₂ FETs, owing to enhanced electron–hole pair separation at the interface. The stacking of CuPc layers on MoS₂ FETs allows for the control of the threshold voltage through a charge transfer phenomenon occurring at the interface. Fig. 17f shows the photoresponsive data of both the pristine MoS₂ FET and CuPc/MoS₂ FETs measured at V_GS = −40 V, while V_DS = 0.1 V is constant under two different conditions: in the darkness and under laser illumination. The experiment involved illuminating the devices with light of different wavelengths (405, 520, 658, and 780 nm) while maintaining a constant power of 40 mW. The results indicated that the photoinduced current exhibited an increment as the incident light's wavelength decreased. This observation suggests that incident light with energies higher than the band gap of CuPc and MoS₂ facilitates a more efficient generation of electron–hole pairs within the material. Moreover, by optimizing the CuPc thickness to approximately 2 nm on the MoS₂ surface, the photodetector demonstrated superior performance. It exhibited a photoresponsivity of approximately 1.98 A W⁻¹, a detectivity of around 6.11 × 10¹⁰ Jones, and an external quantum efficiency of about 12.57%. This study proposes that the vertical hybrid structure, fabricating 2D-TMDs like MoS₂ with organic materials, holds great promise for the development of efficient photodetecting devices and optoelectronic circuits.²⁵⁴

It is essential to mention that while graphene has been investigated for phototransistors in photosensitive devices, it exhibits limitations in photoresponsivity (∼6 mA W⁻¹) owing to its low light absorption coefficient and fast photoinduced carrier recombination rate.³⁴⁰ On the contrary, photodetectors utilizing MoS₂ have demonstrated outstanding photo-response characteristics.²⁵⁴ For instance, devices incorporating single- and multilayer MoS₂ films have shown photoresponsivities reaching 880 A W⁻¹ and approximately 0.1 A W⁻¹, respectively.^163,334 Due to their relatively large light capture cross-sections, organic semiconductors combined with 2D-TMD materials show great potential in enhancing the performance of photodetection and photovoltaic performances.

7. Challenges and future perspectives

By giving increasing importance to TMDO hybrid heterostructures both in fundamental science and for promising applications, these hybrid materials will advance rapidly and have transformative impacts in the area of materials science. While recently there has been notable progress in the fabrication, transport properties, and applications of optoelectronic devices, there remains a need for further improvements in their performance to make them suitable for real-life applications. There are many opportunities in TMDO systems; however several limitations and related challenges remain unaddressed.

i. From the viewpoint of synthesis and fabrication: the stability of monolayer 2D TMD films is comparatively lower than that of their bulk counterparts, and the selection of available 2D-TMD crystals is somewhat limited. To address this, organic molecules, which offer a wide range of possibilities, fill the gap. Furthermore, organic molecular arrangements on TMDs could become well controlled on the atomic scale to build atomically thin TMDO interfaces with controlled molecular configurations. This opens up avenues to investigate the impact of molecular aggregations on light–matter interactions within the hybrid heterostructure system. The successful realization of high-performance organic electronics heavily relies on the use of high-quality organic semiconductors on TMD substrates with minimal disorder. However, achieving such high quality is challenging due to the absence of an epitaxial relation and the presence of disorder, making the growth of defect-free organic crystals with few grain boundaries on conventional insulating substrates particularly difficult. Large-area integration of 2D-TMD materials with organic materials of other dimensionality is likely to have a sizable impact on semiconductor technology. Although reliable heterostructures have been demonstrated from micrometer-scale mechanically exfoliated samples, efforts to synthesize TMD materials reproducibly on the wafer scale are still in their infancy. An atomically thin (1ML) TMD is highly sensitive, therefore transferring it from one substrate to another post synthesis is prone to contamination, strain effects, ripples and voids. However, intensive research on TMDOs is just at the beginning stage. Difficulties in producing scaled-up TMDO heterostructures hamper their development: the mechanical exfoliation and transfer methods are inefficient and not suitable for actual production. More work on in situ growth is urgently needed in a single experimental setup. For the growth of TMDO heterostructures, new approaches must be developed to enable the rational and scalable integration of 2D-TMDs with organic semiconductors.

ii. From the viewpoint of hybrid heterostructure formation—limitation and issues: at the TMDO interface, novel properties with an amazing variety of physical and chemical phenomena can be realized by the elegant design of vertically stacked heterostructures. This interfacial engineering of van der Waals–coupled 2D-TMD layered materials has been rapidly expanding since the van der Waals–coupled layers can virtually be achieved by stacking any kind of 2D materials (which are virtually infinite) together, without any need for epitaxial relation, and can host a great platform for studying a wide range of physical phenomena for multifunctional real-life applications. However, the industrial applications of these coupled layers still have a way to go, because the large-scale and controlled synthesis of coupled 2D-TMD materials of predesigned stacking is still of great challenge. Higher mobility and flexible low-power electronics can be realized by combining the scalable high-quality production methods with the developed device geometry. From the viewpoint of band alignment, at type-I interfaces, the emission properties are mostly enhanced; however, they are scarcely reported. It needs to be explored more in future because type-I TMDO heterostructures are highly desirable for synaptic devices for neuromorphic applications. Further theoretical studies could help us by providing new insights into the interfacial properties. Though type-I and type-II heterostructures have been explored with remarkable properties, type-III band energy alignment is still unexplored and it may give very unusual new optoelectronic properties for multifunctional devices. Further, achieving a high-quality interface between TMDs and organic materials is essential for optimal device performance, but it can be difficult to control and characterize. The synthesis methods for TMDs and organic materials can be both compatible and challenging, as it may require different conditions and equipment. The long-term reliability and durability of TMDO-based devices, particularly under harsh or real-world conditions, are a significant concern.

iii. From the viewpoint of device applications: the challenges are still related to doping methods. In the case of molecular chemisorption/physisorption, the doping effect tends to be gradually reduced over time due to the weak interaction between the adsorbent and TMDs, whereas structural damage (or even etching) by energetic ions remains a major concern, particularly for monolayers, with plasma implantation. Hence, there is a strong demand for developing a methodology to achieve efficient, air-stable, and controllable doping in wafer-scale 2D-TMDs without causing any damage. Additionally, efforts are directed towards reducing contact resistance to enhance the performance of FETs. One major obstacle in using TMDs in semiconductor or optoelectronic platforms encounters a significant challenge posed by the presence of intrinsic defects at their surfaces. These defects arise from volatile chalcogens, leading to trapping states and undesirable doping effects. Furthermore, the electronic or optical properties of TMDs can be perturbed by surface adsorbates introduced during fabrication processes or exposure to ambient air. The presence of these surface defects primarily leads to a degradation in the I_on/I_off ratio in FETs or a reduction in luminescence quantum yields.

Moreover, one of the most fascinating electronic features exhibited by these p–n junctions is the anti-ambipolar phenomenon, which is highly promising for applications in radio communication circuits as frequency multipliers. In comparison to conventional multipliers, the anti-ambipolar-based device showcases significantly enhanced efficiency, surpassing 90%, with the advantage of utilizing just a single device. However, a major challenge faced by TMDO-based multipliers is their limited small signal gain, which arises from the poor electrical performances of the TMDO p–n junction. This discrepancy in electrical properties can be attributed to the mismatch between the 2D-TMDs and the polycrystalline organic semiconductors. In this regard, the high mobility (1–10 cm² V⁻¹ s⁻¹) of organic single crystals, commensurate with the mobility of 2D-TMDs, is crucial. The ultrashort channel configuration in TMDO holds significant potential for high-frequency operations. However, challenges such as interfacial defects, energy level pinning, and dipolar disorder within these structures are currently restricting their operational speed. Again, the atomically thin nature of 2D-TMD materials imposes limitations on light–matter interactions. To overcome this challenge, the integration of organic materials with high optical absorption becomes essential to enhance charge and energy transfer, thereby improving the overall performance of optoelectronic devices. Notably, several photodetectors based on 2D-TMD heterostructures have been reported in recent years, showcasing promising applications in the field of optoelectronics. However, there remain several challenges, leaving ample room for further advancements. To achieve high-performance photodetectors, it is essential to employ well-considered device designs and carefully selected materials for various components of the photodetectors. The prevalent photodetectors employing 2D TMDO heterostructures are typically p–n diodes, which exhibit rapid responses. However, these p–n diodes display low responsivities, primarily due to the low charge carrier concentration. The selection of p-type materials with high mobility leads to the consideration of organic materials as viable candidates. Extensive studies on p-type organic semiconductor layers have made them particularly attractive for this purpose. These organic layers can be conveniently deposited onto 2D-TMD films using spin-coating or vacuum deposition systems, offering precise control over their thickness. This combination of organic materials and 2D-TMD films presents a promising avenue for achieving efficient and controllable photodetector devices. Moreover, the flexible photodetectors utilizing TMDOs hold potential for diverse applications in areas such as military communications, remote sensing, biological imaging, and other fields, making them worthy of serious consideration.

8. Conclusion

In conclusion, this review provided a comprehensive exploration of the exciting developments in 2D TMD/organic-enabled interlayer exciton emission and its modulation across a broad spectrum of 2D-TMD materials and organic layers, with a keen focus on the pivotal role of band energy alignment. We have examined the potential of 2D-TMDO material properties for synaptic devices, delving into their operational principles and showcasing the latest device demonstrations, including field-effect transistors, detectors, and phototransistors. Beginning with an overview of the distinctive characteristics exhibited by different TMDs and organic semiconductors, we have explored the emerging applications unlocked by TMDO hybrid heterostructures. We have also shed light on the physical and chemical synthesis routes and fabrication processes, addressing some recent experimental techniques and suggesting improvements for the fabrication of TMDO hybrid heterostructures. The intricate realm of band energy alignment schemes has been unveiled, illustrating how various TMDs align with a library of organic molecules and polymers, encompassing various band alignment types and elucidating the band bending and charge transfer processes within TMDO heterojunctions from both band structure and molecular orbital theory perspectives. A significant breakthrough, the room-temperature interlayer exciton emission in TMDO hybrid heterojunctions, has been spotlighted. Despite the often-challenging lack of lattice matching in TMDO systems, we presented the latest research advances in emission modulation, exploring phenomena such as photoinduced charge transfer, energy transfer, doping, defect healing, and phase engineering. We provided an in-depth overview of recent progress in 2D-TMDO material–based optoelectronic synaptic devices designed for neuromorphic applications. Our discussion has spanned various 2D-TMDs and organic materials, assessing their compatibility for synaptic devices, revisiting operating principles, and highlighting recent device demonstrations. Additionally, we have closely examined the advancement of field-effect transistors, anti-ambipolar transistors, phototransistors, and the intriguing concept of photo-switchable superlattices. We have shone a spotlight on one of the practical applications – photodetectors based on TMDO hybrid heterostructures, providing valuable insights. Furthermore, we pinpointed the current challenges and proposed potential solutions for TMDO detectors. We have drawn this review to a close by outlining the existing challenges that span from intricate synthesis processes to practical device applications. We presented an outlook on the continuously evolving field of emerging TMDO heterostructures, highlighting the promising horizons awaiting exploration in this fascinating realm of 2D materials science.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors acknowledge support from the Center of Excellence for Advanced Materials and Sensing Devices, ERDF Grant No. KK.01.1.1.01.0001. A. S. acknowledges support from the Austrian Science Fund (FWF) under grant no. I4323-N36. We extend our appreciation to Mr Dayal Das, Mr Md. Nur Hasan, and Dr Arun Kumar Maurya from S. N. Bose National Centre for Basic Sciences (SNBNCBS), India for their valuable suggestions.

References

1. F. Liu, W. L. Chow, X. He, P. Hu, S. Zheng, X. Wang, J. Zhou, Q. Fu, W. Fu, P. Yu, Q. Zeng, H. J. Fan, B. K. Tay, C. Kloc and Z. Liu, Adv. Funct. Mater., 2015, 25, 5865.
2. M. Chhowalla, H. S. Shin, G. Eda, L.-J. Li, K. P. Loh and H. Zhang, Nat. Chem., 2013, 5, 263.
3. A. V. K. A. J. Tominaga, Two-Dimensional Transition-Metal Dichalcogenides, Springer International Publishing, Switzerland, 2016.
4. A. Kuc and T. Heine, Chem. Soc. Rev., 2015, 44, 2603.
5. L. M. Xie, Nanoscale, 2015, 7, 18392.
6. S. Bertolazzi, M. Gobbi, Y. Zhao, C. Backes and P. Samorì, Chem. Soc. Rev., 2018, 47, 6845.
7. H. Wang, Z. Li, D. Li, X. Xu, P. Chen, L. Pi, X. Zhou and T. Zhai, Adv. Funct. Mater., 2021, 31, 2106105.
8. Z. Wang, R. Li, C. Su and K. P. Loh, Smart Mat, 2020, 1, e1013.
9. W.-M. Zhao, L. Zhu, Z. Nie, Q.-Y. Li, Q.-W. Wang, L.-G. Dou, J.-G. Hu, L. Xian, S. Meng and S.-C. Li, Nat. Mater., 2022, 21, 284.
10. N. R. Wilson, P. V. Nguyen, K. Seyler, P. Rivera, A. J. Marsden, Z. P. L. Laker, G. C. Constantinescu, V. Kandyba, A. Barinov, N. D. M. Hine, X. Xu and D. H. Cobden, Sci. Adv., 2017, 3, e1601832.
11. G.-B. Liu, W.-Y. Shan, Y. Yao, W. Yao and D. Xiao, Phys. Rev. B: Condens. Matter Mater. Phys., 2013, 88, 085433.
12. S. H. Amsterdam, T. J. Marks and M. C. Hersam, J. Phys. Chem. Lett., 2021, 12, 4543.
13. Y. Kong, S. M. Obaidulla, M. R. Habib, Z. Wang, R. Wang, Y. Khan, H. Zhu, M. Xu and D. Yang, Mater. Horiz, 2022, 9, 1253.
14. S. M. Obaidulla, M. R. Habib, Y. Khan, Y. Kong, T. Liang and M. Xu, Adv. Mater. Interfaces, 2020, 7, 1901197.
15. F. Zhang, Y. Ma, Y. Chi, H. Yu, Y. Li, T. Jiang, X. Wei and J. Shi, Sci. Rep., 2018, 8, 8208.
16. L. Zhang, A. Sharma, Y. Zhu, Y. Zhang, B. Wang, M. Dong, H. T. Nguyen, Z. Wang, B. Wen, Y. Cao, B. Liu, X. Sun, J. Yang, Z. Li, A. Kar, Y. Shi, D. Macdonald, Z. Yu, X. Wang and Y. Lu, Adv. Mater., 2018, 30, 1803986.
17. D. Jariwala, T. J. Marks and M. C. Hersam, Nat. Mater., 2017, 16, 170.
18. T. Zhu, L. Yuan, Y. Zhao, M. Zhou, Y. Wan, J. Mei and L. Huang, Sci. Adv., 2018, 4, eaao3104.
19. J. J. P. Thompson, V. Lumsargis, M. Feierabend, Q. Zhao, K. Wang, L. Dou, L. Huang and E. Malic, Nanoscale, 2023, 15, 1730.
20. S. Park, N. Mutz, S. A. Kovalenko, T. Schultz, D. Shin, A. Aljarb, L.-J. Li, V. Tung, P. Amsalem, E. J. W. List-Kratochvil, J. Stähler, X. Xu, S. Blumstengel and N. Koch, Adv. Sci., 2021, 8, 2100215.
21. D.-H. Kang, M.-S. Kim, J. Shim, J. Jeon, H.-Y. Park, W.-S. Jung, H.-Y. Yu, C.-H. Pang, S. Lee and J.-H. Park, Adv. Funct. Mater., 2015, 25, 4219.
22. S. Najmaei, X. Zou, D. Er, J. Li, Z. Jin, W. Gao, Q. Zhang, S. Park, L. Ge, S. Lei, J. Kono, V. B. Shenoy, B. I. Yakobson, A. George, P. M. Ajayan and J. Lou, Nano Lett., 2014, 14, 1354.
23. S. H. Yu, Y. Lee, S. K. Jang, J. Kang, J. Jeon, C. Lee, J. Y. Lee, H. Kim, E. Hwang, S. Lee and J. H. Cho, ACS Nano, 2014, 8, 8285.
24. Y. J. Jang and J.-H. Kim, Chem. – Asian J., 2022, 17, e202200265.
25. T. A. Shastry, I. Balla, H. Bergeron, S. H. Amsterdam, T. J. Marks and M. C. Hersam, ACS Nano, 2016, 10, 10573.
26. K. S. Lee, Y. J. Park, J. Shim, C.-H. Lee, G.-H. Lim, H. Y. Kim, J. W. Choi, C.-L. Lee, Y. Jin, K. Yu, H.-S. Chung, B. Angadi, S.-I. Na and D. I. Son, J. Mater. Chem. A, 2019, 7, 15356.
27. Q. Luo, G. Feng, Y. Song, E. Zhang, J. Yuan, D. Fa, Q. Sun, S. Lei and W. Hu, Nano Res., 2022, 15, 8428.
28. X. Wang, P. Wang, J. Wang, W. Hu, X. Zhou, N. Guo, H. Huang, S. Sun, H. Shen, T. Lin, M. Tang, L. Liao, A. Jiang, J. Sun, X. Meng, X. Chen, W. Lu and J. Chu, Adv. Mater., 2015, 27, 6575.
29. N. Zorn Morales, N. Severin, J. P. Rabe, S. Kirstein, E. List-Kratochvil and S. Blumstengel, Adv. Opt. Mater., 2023, 2301057.
30. S. Kwon, D. Y. Jeong, C. Hong, S. Oh, J. Song, S. H. Choi, K. K. Kim, S. Yoon, T. Choi, K.-J. Yee, J.-H. Kim, Y. You and D.-W. Kim, Adv. Sci., 2022, 9, 2201875.
31. P. Rivera, J. R. Schaibley, A. M. Jones, J. S. Ross, S. Wu, G. Aivazian, P. Klement, K. Seyler, G. Clark, N. J. Ghimire, J. Yan, D. G. Mandrus, W. Yao and X. Xu, Nat. Commun., 2015, 6, 6242.
32. Y. Jiang, S. Chen, W. Zheng, B. Zheng and A. Pan, Light: Sci. Appl., 2021, 10, 72.
33. M. Troue, J. Figueiredo, L. Sigl, C. Paspalides, M. Katzer, T. Taniguchi, K. Watanabe, M. Selig, A. Knorr, U. Wurstbauer and A. W. Holleitner, Phys. Rev. Lett., 2023, 131, 036902.
34. A. Ripin, R. Peng, X. Zhang, S. Chakravarthi, M. He, X. Xu, K.-M. Fu, T. Cao and M. Li, Nat. Nanotechnol., 2023, 18, 1020.
35. K. F. Mak and J. Shan, Nat. Nanotechnol., 2018, 13, 974.
36. K. Ulman and S. Y. Quek, Nano Lett., 2021, 21, 8888.
37. X. Xiong, J. Kang, Q. Hu, C. Gu, T. Gao, X. Li and Y. Wu, Adv. Funct. Mater., 2020, 30, 1909645.
38. J. Khan, R. T. M. Ahmad, J. Tan, R. Zhang, U. Khan and B. Liu, Smart Mater., 2023, 4, e1156.
39. J. Li, J. Wierzbowski, Ö. Ceylan, J. Klein, F. Nisic, T. L. Anh, F. Meggendorfer, C.-A. Palma, C. Dragonetti, J. V. Barth, J. J. Finley and E. Margapoti, Appl. Phys. Lett., 2014, 105, 241116.
40. Y. Li, C.-Y. Xu, P. Hu and L. Zhen, ACS Nano, 2013, 7, 7795.
41. A. Castellanos-Gomez, M. Buscema, R. Molenaar, V. Singh, L. Janssen, H. S. J. van der Zant and G. A. Steele, 2D Mater., 2014, 1, 011002.
42. H. F. Liu, S. L. Wong and D. Z. Chi, Chem. Vap. Deposition, 2015, 21, 241.
43. K. Kang, S. Xie, L. Huang, Y. Han, P. Y. Huang, K. F. Mak, C.-J. Kim, D. Muller and J. Park, Nature, 2015, 520, 656.
44. R. Dong and I. Kuljanishvili, J. Vac. Sci. Technol. B, 2017, 35, 030803.
45. R. Guan, J. Duan, A. Yuan, Z. Wang, S. Yang, L. Han, B. Zhang, D. Li and B. Luo, CrystEngComm, 2021, 23, 146.
46. A. Dimoulas, Adv. Mater. Interfaces, 2022, 9, 2201469.
47. S. M. Eichfeld, L. Hossain, Y.-C. Lin, A. F. Piasecki, B. Kupp, A. G. Birdwell, R. A. Burke, N. Lu, X. Peng, J. Li, A. Azcatl, S. McDonnell, R. M. Wallace, M. J. Kim, T. S. Mayer, J. M. Redwing and J. A. Robinson, ACS Nano, 2015, 9, 2080.
48. X. He, W. Chow, F. Liu, B. Tay and Z. Liu, Small, 2017, 13, 1602558.
49. R. Chua, J. Zhou, X. Yu, W. Yu, J. Gou, R. Zhu, L. Zhang, M. Liu, M. B. H. Breese, W. Chen, K. P. Loh, Y. P. Feng, M. Yang, Y. L. Huang and A. T. S. Wee, Adv. Mater., 2021, 33, 2103360.
50. Z. Song, T. Schultz, Z. Ding, B. Lei, C. Han, P. Amsalem, T. Lin, D. Chi, S. L. Wong, Y. J. Zheng, M.-Y. Li, L.-J. Li, W. Chen, N. Koch, Y. L. Huang and A. T. S. Wee, ACS Nano, 2017, 11, 9128.
51. S. H. Amsterdam, T. K. Stanev, Q. Zhou, A. J. T. Lou, H. Bergeron, P. Darancet, M. C. Hersam, N. P. Stern and T. J. Marks, ACS Nano, 2019, 13, 4183.
52. J. H. Park, A. Sanne, Y. Guo, M. Amani, K. Zhang, H. C. P. Movva, J. A. Robinson, A. Javey, J. Robertson, S. K. Banerjee and A. C. Kummel, Sci. Adv., 2017, 3, e1701661.
53. T. R. Kafle, B. Kattel, P. Yao, P. Zereshki, H. Zhao and W.-L. Chan, J. Am. Chem. Soc., 2019, 141, 11328.
54. K. Lasek, P. M. Coelho, P. Gargiani, M. Valvidares, K. Mohseni, H. L. Meyerheim, I. Kostanovskiy, K. Zberecki and M. Batzill, Appl. Phys. Rev., 2022, 9, 011409.
55. M. Liu, Y. L. Huang, J. Gou, Q. Liang, R. Chua, Arramel, S. Duan, L. Zhang, L. L. Cai, X. Yu, D. Zhong, W. Zhang and A. T. S. Wee, J. Phys. Chem. Lett., 2021, 12, 7752.
56. H. Wang, Y. Liu, P. Wu, W. Hou, Y. Jiang, X. Li, C. Pandey, D. Chen, Q. Yang, H. Wang, D. Wei, N. Lei, W. Kang, L. Wen, T. Nie, W. Zhao and K. L. Wang, ACS Nano, 2020, 14, 10045.
57. Y. Ou, W. Yanez, R. Xiao, M. Stanley, S. Ghosh, B. Zheng, W. Jiang, Y.-S. Huang, T. Pillsbury, A. Richardella, C. Liu, T. Low, V. H. Crespi, K. A. Mkhoyan and N. Samarth, Nat. Commun., 2022, 13, 2972.
58. K. Lasek, P. M. Coelho, K. Zberecki, Y. Xin, S. K. Kolekar, J. Li and M. Batzill, ACS Nano, 2020, 14, 8473.
59. H. Qiu, Y. Zhao, Z. Liu, M. Herder, S. Hecht and P. Samorì, Adv. Mater., 2019, 31, 1903402.
60. M. Gobbi, S. Bonacchi, J. X. Lian, Y. Liu, X.-Y. Wang, M.-A. Stoeckel, M. A. Squillaci, G. D’Avino, A. Narita, K. Müllen, X. Feng, Y. Olivier, D. Beljonne, P. Samorì and E. Orgiu, Nat. Commun., 2017, 8, 14767.
61. Y. Zhao, S. Bertolazzi and P. Samorì, ACS Nano, 2019, 13, 4814.
62. M. Gobbi, S. Bonacchi, J. X. Lian, A. Vercouter, S. Bertolazzi, B. Zyska, M. Timpel, R. Tatti, Y. Olivier, S. Hecht, M. V. Nardi, D. Beljonne, E. Orgiu and P. Samorì, Nat. Commun., 2018, 9, 2661.
63. W. Su, L. Jin, X. Qu, D. Huo and L. Yang, Phys. Chem. Chem. Phys., 2016, 18, 14001.
64. Y. Zhao, M. Gobbi, L. E. Hueso and P. Samorì, Chem. Rev., 2022, 122, 50.
65. A. J. Molina-Mendoza, L. Vaquero-Garzon, S. Leret, L. de Juan-Fernández, E. M. Pérez and A. Castellanos-Gomez, Chem. Commun., 2016, 52, 14365.
66. N. Peimyoo, W. Yang, J. Shang, X. Shen, Y. Wang and T. Yu, ACS Nano, 2014, 8, 11320.
67. H. G. Ji, P. Solís-Fernández, D. Yoshimura, M. Maruyama, T. Endo, Y. Miyata, S. Okada and H. Ago, Adv. Mater., 2019, 31, 1903613.
68. P. Zhao, D. Kiriya, A. Azcatl, C. Zhang, M. Tosun, Y.-S. Liu, M. Hettick, J. S. Kang, S. McDonnell, S. Kc, J. Guo, K. Cho, R. M. Wallace and A. Javey, ACS Nano, 2014, 8, 10808.
69. Y. Wang, S. M. Gali, A. Slassi, D. Beljonne and P. Samorì, Adv. Funct. Mater., 2020, 30, 2002846.
70. J. Qian, S. Jiang, S. Li, X. Wang, Y. Shi and Y. Li, Adv. Mater. Technol., 2019, 4, 1800182.
71. L. Wang, C. Wang, X. Yu, L. Zheng, X. Zhang and W. Hu, Sci. China: Mater., 2020, 63, 122.
72. L. Jiang, H. Dong, Q. Meng, H. Li, M. He, Z. Wei, Y. He and W. Hu, Adv. Mater., 2011, 23, 2059.
73. A. S. Sizov, D. S. Anisimov, E. V. Agina, O. V. Borshchev, A. V. Bakirov, M. A. Shcherbina, S. Grigorian, V. V. Bruevich, S. N. Chvalun, D. Y. Paraschuk and S. A. Ponomarenko, Langmuir, 2014, 30, 15327.
74. S. Behura and V. Berry, ACS Nano, 2015, 9, 2227.
75. J. Roberts, Langmuir-Blodgett Deposition of 2D Materials for Unique Identification, Springer International Publishing, Cham, 2017.
76. G. Hu, J. Kang, L. W. T. Ng, X. Zhu, R. C. T. Howe, C. G. Jones, M. C. Hersam and T. Hasan, Chem. Soc. Rev., 2018, 47, 3265.
77. A. G. Kelly, T. Hallam, C. Backes, A. Harvey, A. S. Esmaeily, I. Godwin, J. Coelho, V. Nicolosi, J. Lauth, A. Kulkarni, S. Kinge, L. D. A. Siebbeles, G. S. Duesberg and J. N. Coleman, Science, 2017, 356, 69.
78. D. McManus, S. Vranic, F. Withers, V. Sanchez-Romaguera, M. Macucci, H. Yang, R. Sorrentino, K. Parvez, S.-K. Son, G. Iannaccone, K. Kostarelos, G. Fiori and C. Casiraghi, Nat. Nanotechnol., 2017, 12, 343.
79. F. Cui, X. Zhao, J. Xu, B. Tang, Q. Shang, J. Shi, Y. Huan, J. Liao, Q. Chen, Y. Hou, Q. Zhang, S. J. Pennycook and Y. Zhang, Adv. Mater., 2020, 32, 1905896.
80. C. Chen, X. Chen, C. Wu, X. Wang, Y. Ping, X. Wei, X. Zhou, J. Lu, L. Zhu, J. Zhou, T. Zhai, J. Han and H. Xu, Adv. Mater., 2022, 34, 2107512.
81. B. Qin, M. Z. Saeed, Q. Li, M. Zhu, Y. Feng, Z. Zhou, J. Fang, M. Hossain, Z. Zhang, Y. Zhou, Y. Huangfu, R. Song, J. Tang, B. Li, J. Liu, D. Wang, K. He, H. Zhang, R. Wu, B. Zhao, J. Li, L. Liao, Z. Wei, B. Li, X. Duan and X. Duan, Nat. Commun., 2023, 14, 304.
82. J. Zhou, J. Lin, X. Huang, Y. Zhou, Y. Chen, J. Xia, H. Wang, Y. Xie, H. Yu, J. Lei, D. Wu, F. Liu, Q. Fu, Q. Zeng, C.-H. Hsu, C. Yang, L. Lu, T. Yu, Z. Shen, H. Lin, B. I. Yakobson, Q. Liu, K. Suenaga, G. Liu and Z. Liu, Nature, 2018, 556, 355.
83. X. Hu, P. Huang, B. Jin, X. Zhang, H. Li, X. Zhou and T. Zhai, J. Am. Chem. Soc., 2018, 140, 12909.
84. S. Li, S. Wang, D.-M. Tang, W. Zhao, H. Xu, L. Chu, Y. Bando, D. Golberg and G. Eda, Appl. Mater. Today, 2015, 1, 60.
85. Y. Zhang, Y. Zhang, Q. Ji, J. Ju, H. Yuan, J. Shi, T. Gao, D. Ma, M. Liu, Y. Chen, X. Song, H. Y. Hwang, Y. Cui and Z. Liu, ACS Nano, 2013, 7, 8963.
86. H. Li, Y. Li, A. Aljarb, Y. Shi and L.-J. Li, Chem. Rev., 2018, 118, 6134.
87. A. L. Elías, N. Perea-López, A. Castro-Beltrán, A. Berkdemir, R. Lv, S. Feng, A. D. Long, T. Hayashi, Y. A. Kim, M. Endo, H. R. Gutiérrez, N. R. Pradhan, L. Balicas, T. E. Mallouk, F. López-Urías, H. Terrones and M. Terrones, ACS Nano, 2013, 7, 5235.
88. Y. Yu, C. Li, Y. Liu, L. Su, Y. Zhang and L. Cao, Sci. Rep., 2013, 3, 1866.
89. Z. Cai, B. Liu, X. Zou and H.-M. Cheng, Chem. Rev., 2018, 118, 6091.
90. S. Zhou, R. Wang, J. Han, D. Wang, H. Li, L. Gan and T. Zhai, Adv. Funct. Mater., 2019, 29, 1805880.
91. Y.-H. Lee, L. Yu, H. Wang, W. Fang, X. Ling, Y. Shi, C.-T. Lin, J.-K. Huang, M.-T. Chang, C.-S. Chang, M. Dresselhaus, T. Palacios, L.-J. Li and J. Kong, Nano Lett., 2013, 13, 1852.
92. H. Kim, D. Ovchinnikov, D. Deiana, D. Unuchek and A. Kis, Nano Lett., 2017, 17, 5056.
93. C. Gong, H. Zhang, W. Wang, L. Colombo, R. M. Wallace and K. Cho, Appl. Phys. Lett., 2013, 103, 053513.
94. Y. Sun, Z. Zhou, Z. Huang, J. Wu, L. Zhou, Y. Cheng, J. Liu, C. Zhu, M. Yu, P. Yu, W. Zhu, Y. Liu, J. Zhou, B. Liu, H. Xie, Y. Cao, H. Li, X. Wang, K. Liu, X. Wang, J. Wang, L. Wang and W. Huang, Adv. Mater., 2019, 31, 1806562.
95. L. A. Burton, T. J. Whittles, D. Hesp, W. M. Linhart, J. M. Skelton, B. Hou, R. F. Webster, G. O'Dowd, C. Reece, D. Cherns, D. J. Fermin, T. D. Veal, V. R. Dhanak and A. Walsh, J. Mater. Chem. A, 2016, 4, 1312.
96. J. Wang, R. Jia, Q. Huang, C. Pan, J. Zhu, H. Wang, C. Chen, Y. Zhang, Y. Yang, H. Song, F. Miao and R. Huang, Sci. Rep., 2018, 8, 17755.
97. D. Jariwala, S. L. Howell, K.-S. Chen, J. Kang, V. K. Sangwan, S. A. Filippone, R. Turrisi, T. J. Marks, L. J. Lauhon and M. C. Hersam, Nano Lett., 2016, 16, 497.
98. X. Liu, J. Gu, K. Ding, D. Fan, X. Hu, Y.-W. Tseng, Y.-H. Lee, V. Menon and S. R. Forrest, Nano Lett., 2017, 17, 3176.
99. K. Kobashi, R. Hayakawa, T. Chikyow and Y. Wakayama, Adv. Electron. Mater., 2017, 3, 1700106.
100. H. Zhao, Y. Zhao, Y. Song, M. Zhou, W. Lv, L. Tao, Y. Feng, B. Song, Y. Ma, J. Zhang, J. Xiao, Y. Wang, D.-H. Lien, M. Amani, H. Kim, X. Chen, Z. Wu, Z. Ni, P. Wang, Y. Shi, H. Ma, X. Zhang, J.-B. Xu, A. Troisi, A. Javey and X. Wang, Nat. Commun., 2019, 10, 5589.
101. Y. Khan, S. M. Obaidulla, M. Rezwan Habib, Y. Kong and M. Xu, Appl. Surf. Sci., 2020, 530, 147213.
102. S. Vélez, D. Ciudad, J. Island, M. Buscema, O. Txoperena, S. Parui, G. A. Steele, F. Casanova, H. S. J. van der Zant, A. Castellanos-Gomez and L. E. Hueso, Nanoscale, 2015, 7, 15442.
103. Z. Liu, X. Zhang, Y. Zhang and J. Jiang, Spectrochim. Acta, Part A, 2007, 67, 1232.
104. P. Choudhury, L. Ravavarapu, R. Dekle and S. Chowdhury, J. Phys. Chem. C, 2017, 121, 2959.
105. Y. Huang, F. Zhuge, J. Hou, L. Lv, P. Luo, N. Zhou, L. Gan and T. Zhai, ACS Nano, 2018, 12, 4062.
106. J. Choi, H. Zhang and J. H. Choi, ACS Nano, 2016, 10, 1671.
107. Y. Garcia-Basabe, G. G. Parra, M. B. Barioni, C. D. Mendoza, F. C. Vicentin and D. G. Larrudé, Phys. Chem. Chem. Phys., 2019, 21, 23521.
108. M. Sun, P. Yang, D. Xie, Y. Sun, J. Xu, T. Ren and Y. Zhang, Adv. Electron. Mater., 2019, 5, 1800580.
109. D. He, Y. Pan, H. Nan, S. Gu, Z. Yang, B. Wu, X. Luo, B. Xu, Y. Zhang, Y. Li, Z. Ni, B. Wang, J. Zhu, Y. Chai, Y. Shi and X. Wang, Appl. Phys. Lett., 2015, 107, 183103.
110. K. Osada, M. Tanaka, S. Ohno and T. Suzuki, Jpn. J. Appl. Phys., 2016, 55, 065201.
111. H. G. Shin, D. Kang, Y. Jeong, K. Kim, Y. Cho, J. Park, S. Hong, Y. Yi and S. Im, ACS Nano, 2020, 14, 15646.
112. M. Tebyetekerwa, Y. Cheng, J. Zhang, W. Li, H. Li, G. P. Neupane, B. Wang, T. N. Truong, C. Xiao, M. M. Al-Jassim, Z. Yin, Y. Lu, D. Macdonald and H. T. Nguyen, ACS Nano, 2020, 14, 7444.
113. J. I. Beltrán, F. Flores, J. I. Martínez and J. Ortega, J. Phys. Chem. C, 2013, 117, 3888.
114. J. Wang, Z. Ji, G. Yang, X. Chuai, F. Liu, Z. Zhou, C. Lu, W. Wei, X. Shi, J. Niu, L. Wang, H. Wang, J. Chen, N. Lu, C. Jiang, L. Li and M. Liu, Adv. Funct. Mater., 2018, 28, 1806244.
115. S. Mouri, Y. Miyauchi and K. Matsuda, Nano Lett., 2013, 13, 5944.
116. H. Liu, M. Cui, C. Dang, W. Wen, X. Wang and L. Xie, ACS Appl. Mater. Interfaces, 2019, 11, 34424.
117. D. Kiriya, M. Tosun, P. Zhao, J. S. Kang and A. Javey, J. Am. Chem. Soc., 2014, 136, 7853.
118. Y. Cho, J. H. Park, M. Kim, Y. Jeong, S. Yu, J. Y. Lim, Y. Yi and S. Im, Nano Lett., 2019, 19, 2456.
119. K. F. Mak, C. Lee, J. Hone, J. Shan and T. F. Heinz, Phys. Rev. Lett., 2010, 105, 136805.
120. A. Splendiani, L. Sun, Y. Zhang, T. Li, J. Kim, C.-Y. Chim, G. Galli and F. Wang, Nano Lett., 2010, 10, 1271.
121. K.-S. Jung, K. Heo, M.-J. Kim, M. Andreev, S. Seo, J.-O. Kim, J.-H. Lim, K.-H. Kim, S. Kim, K. S. Kim, G. Y. Yeom, J. H. Cho and J.-H. Park, Adv. Sci., 2020, 7, 2000991.
122. L. Sun, Y. Zhang, G. Hwang, J. Jiang, D. Kim, Y. A. Eshete, R. Zhao and H. Yang, Nano Lett., 2018, 18, 3229.
123. V. O. Özçelik, J. G. Azadani, C. Yang, S. J. Koester and T. Low, Phys. Rev. B, 2016, 94, 035125.
124. N. Mutz, S. Park, T. Schultz, S. Sadofev, S. Dalgleish, L. Reissig, N. Koch, E. J. W. List-Kratochvil and S. Blumstengel, J. Phys. Chem. C, 2020, 124, 2837.
125. A. F. J. Levi and T. H. Chiu, Appl. Phys. Lett., 1987, 51, 984.
126. J. D. Werking, C. R. Bolognesi, L. D. Chang, C. Nguyen, E. L. Hu and H. Kroemer, IEEE Electron. Dev. Lett., 1992, 13, 164.
127. W. Xu, D. Kozawa, Y. Zhou, Y. Wang, Y. Sheng, T. Jiang, M. S. Strano and J. H. Warner, Small, 2020, 16, 1905985.
128. J. Gu, X. Liu, E.-C. Lin, Y.-H. Lee, S. R. Forrest and V. M. Menon, ACS Photonics, 2018, 5, 100.
129. R. Yan, S. Fathipour, Y. Han, B. Song, S. Xiao, M. Li, N. Ma, V. Protasenko, D. A. Muller, D. Jena and H. G. Xing, Nano Lett., 2015, 15, 5791.
130. C. Lei, Y. Ma, X. Xu, T. Zhang, B. Huang and Y. Dai, J. Phys. Chem. C, 2019, 123, 23089.
131. S. Bettis Homan, V. K. Sangwan, I. Balla, H. Bergeron, E. A. Weiss and M. C. Hersam, Nano Lett., 2017, 17, 164.
132. N. Shen and G. Tao, Adv. Mater. Interfaces, 2017, 4, 1601083.
133. J. Dong, F. Liu, F. Wang, J. Wang, M. Li, Y. Wen, L. Wang, G. Wang, J. He and C. Jiang, Nanoscale, 2017, 9, 7519.
134. Q. Ren, Q. Xu, H. Xia, X. Luo, F. Zhao, L. Sun, Y. Li, W. Lv, L. Du, Y. Peng and Z. Zhao, Org. Electron., 2017, 51, 142.
135. E. Black, P. Kratzer and J. M. Morbec, Phys. Chem. Chem. Phys., 2023 10.1039/D3CP01895D.
136. X. Lin, J. C. Lu, Y. Shao, Y. Y. Zhang, X. Wu, J. B. Pan, L. Gao, S. Y. Zhu, K. Qian, Y. F. Zhang, D. L. Bao, L. F. Li, Y. Q. Wang, Z. L. Liu, J. T. Sun, T. Lei, C. Liu, J. O. Wang, K. Ibrahim, D. N. Leonard, W. Zhou, H. M. Guo, Y. L. Wang, S. X. Du, S. T. Pantelides and H. J. Gao, Nat. Mater., 2017, 16, 717.
137. L. Ye, Y. Liu, Q. Zhou, W. Tao, Y. Li, Z. Wang and H. Zhu, J. Phys. Chem. Lett., 2021, 12, 8440.
138. S. Wang, C. Chen, Z. Yu, Y. He, X. Chen, Q. Wan, Y. Shi, D. W. Zhang, H. Zhou, X. Wang and P. Zhou, Adv. Mater., 2019, 31, 1806227.
139. M. R. Habib, H. Li, Y. Kong, T. Liang, S. M. Obaidulla, S. Xie, S. Wang, X. Ma, H. Su and M. Xu, Nanoscale, 2018, 10, 16107.
140. K. Rijal, F. Rudayni, T. R. Kafle and W.-L. Chan, J. Phys. Chem. Lett., 2020, 11, 7495.
141. K. Rijal, S. Amos, P. Valencia-Acuna, F. Rudayni, N. Fuller, H. Zhao, H. Peelaers and W.-L. Chan, ACS Nano, 2023, 17, 7775.
142. Z. Ye, T. Cao, K. O’Brien, H. Zhu, X. Yin, Y. Wang, S. G. Louie and X. Zhang, Nature, 2014, 513, 214.
143. N. Mrkyvkova, M. Hodas, J. Hagara, P. Nadazdy, Y. Halahovets, M. Bodik, K. Tokar, J. W. Chai, S. J. Wang, D. Z. Chi, A. Chumakov, O. Konovalov, A. Hinderhofer, M. Jergel, E. Majkova, P. Siffalovic and F. Schreiber, Appl. Phys. Lett., 2019, 114, 251906.
144. E. H. Cho, W. G. Song, C. J. Park, J. Kim, S. Kim and J. Joo, Nano Res., 2015, 8, 790.
145. M. C. Schwinn, S. Rafiq, C. Lee, M. P. Bland, T. W. Song, V. K. Sangwan, M. C. Hersam and L. X. Chen, J. Chem. Phys., 2022, 157, 184701.
146. Z. Wang, C. Sun, X. Xu, Y. Liu, Z. Chen, Y. M. Yang and H. Zhu, J. Am. Chem. Soc., 2023, 145, 11227.
147. G. Ghimire, S. Adhikari, S. G. Jo, H. Kim, J. Jiang, J. Joo and J. Kim, J. Phys. Chem. C, 2018, 122, 6794.
148. T. R. Kafle, B. Kattel, S. D. Lane, T. Wang, H. Zhao and W.-L. Chan, ACS Nano, 2017, 11, 10184.
149. P. Valencia-Acuna, T. R. Kafle, P. Zereshki, H. Peelaers, W.-L. Chan and H. Zhao, Appl. Phys. Lett., 2021, 118, 153104.
150. C. J. Benjamin, S. Zhang and Z. Chen, Nanoscale, 2018, 10, 5148.
151. Y. Zhang, S. Liu, W. Liu, T. Liang, X. Yang, M. Xu and H. Chen, Phys. Chem. Chem. Phys., 2015, 17, 27565.
152. X.-Y. Liu, W.-K. Chen, W.-H. Fang and G. Cui, J. Phys. Chem. Lett., 2019, 10, 2949.
153. J. C. Bijleveld, A. P. Zoombelt, S. G. J. Mathijssen, M. M. Wienk, M. Turbiez, D. M. de Leeuw and R. A. J. Janssen, J. Am. Chem. Soc., 2009, 131, 16616.
154. A. A. Abdelaziz, Y. Teramoto and H.-H. Kim, J. Phys. D: Appl. Phys., 2021, 55, 065201.
155. C. Zhao, W. Tao, Z. Chen, H. Zhou, C. Zhang, J. Lin and H. Zhu, J. Phys. Chem. Lett., 2021, 12, 3691.
156. Z. Song, Q. Wang, M.-Y. Li, L.-J. Li, Y. J. Zheng, Z. Wang, T. Lin, D. Chi, Z. Ding, Y. L. Huang and A. Thye Shen Wee, Phys. Rev. B, 2018, 97, 134102.
157. O. Karni, E. Barré, S. C. Lau, R. Gillen, E. Y. Ma, B. Kim, K. Watanabe, T. Taniguchi, J. Maultzsch, K. Barmak, R. H. Page and T. F. Heinz, Phys. Rev. Lett., 2019, 123, 247402.
158. P. Rivera, H. Yu, K. L. Seyler, N. P. Wilson, W. Yao and X. Xu, Nat. Nanotechnol., 2018, 13, 1004.
159. M. Combescot, R. Combescot and F. Dubin, Rep. Progress Phys., 2017, 80, 066501.
160. A. A. High, E. E. Novitskaya, L. V. Butov, M. Hanson and A. C. Gossard, Science, 2008, 321, 229.
161. A. Ciarrocchi, D. Unuchek, A. Avsar, K. Watanabe, T. Taniguchi and A. Kis, Nat. Photonics, 2019, 13, 131.
162. L. A. Jauregui, A. Y. Joe, K. Pistunova, D. S. Wild, A. A. High, Y. Zhou, G. Scuri, K. D. Greve, A. Sushko, C.-H. Yu, T. Taniguchi, K. Watanabe, D. J. Needleman, M. D. Lukin, H. Park and P. Kim, Science, 2019, 366, 870.
163. W. Choi, M. Y. Cho, A. Konar, J. H. Lee, G.-B. Cha, S. C. Hong, S. Kim, J. Kim, D. Jena, J. Joo and S. Kim, Adv. Mater., 2012, 24, 5832.
164. S. M. Obaidulla, S. Singh, Y. N. Mohapatra and P. K. Giri, J. Phys. D: Appl. Phys, 2017, 51, 015110.
165. J. Ryou, Y.-S. Kim, S. Kc and K. Cho, Sci. Rep., 2016, 6, 29184.
166. J. D. Wright, Molecular Crystals, Cambridge University Press, 1995.
167. X. Hong, J. Kim, S.-F. Shi, Y. Zhang, C. Jin, Y. Sun, S. Tongay, J. Wu, Y. Zhang and F. Wang, Nat. Nanotechnol., 2014, 9, 682.
168. H. Fang, C. Battaglia, C. Carraro, S. Nemsak, B. Ozdol, J. S. Kang, H. A. Bechtel, S. B. Desai, F. Kronast, A. A. Unal, G. Conti, C. Conlon, G. K. Palsson, M. C. Martin, A. M. Minor, C. S. Fadley, E. Yablonovitch, R. Maboudian and A. Javey, Proc. Natl. Acad. Sci. U. S. A., 2014, 111, 6198.
169. J. Hagel, S. Brem, C. Linderälv, P. Erhart and E. Malic, Phys. Rev. Res., 2021, 3, 043217.
170. J. Ji, C. M. Delehey, D. N. Houpt, M. K. Heighway, T. Lee and J. H. Choi, Nano Lett., 2020, 20, 2500.
171. R. Precht, R. Hausbrand and W. Jaegermann, Phys. Chem. Chem. Phys., 2015, 17, 6588.
172. G. Mondio, F. Neri, G. Curró, S. Patané and G. Compagnini, J. Mater. Res., 1993, 8, 2627.
173. J. Ji and J. H. Choi, Adv. Mater. Interfaces, 2019, 6, 1900637.
174. R. Sharma, A. Kamal and R. K. Mahajan, RSC Adv., 2016, 6, 71692.
175. J. Ji and J. H. Choi, ACS Appl. Electron. Mater., 2021, 3, 3052.
176. M. R. Habib, W. Wang, A. Khan, Y. Khan, S. M. Obaidulla, X. Pi and M. Xu, Adv. Theory Simul., 2020, 3, 2000045.
177. L. Kou, Y. Ma, S. C. Smith and C. Chen, J. Phys. Chem. Lett., 2015, 6, 1509.
178. T. Tian and C.-J. Shih, Ind. Eng. Chem. Res., 2017, 56, 10552.
179. H. Yoshida, K. Yamada, J. Y. Tsutsumi and N. Sato, Phys. Rev. B: Condens. Matter Mater. Phys., 2015, 92, 075145.
180. M. Amani, D.-H. Lien, D. Kiriya, J. Xiao, A. Azcatl, J. Noh, S. R. Madhvapathy, R. Addou, K. C. Santosh, M. Dubey, K. Cho, R. M. Wallace, S.-C. Lee, J.-H. He, J. W. Ager, X. Zhang, E. Yablonovitch and A. Javey, Science, 2015, 350, 1065.
181. H.-V. Han, A.-Y. Lu, L.-S. Lu, J.-K. Huang, H. Li, C.-L. Hsu, Y.-C. Lin, M.-H. Chiu, K. Suenaga, C.-W. Chu, H.-C. Kuo, W.-H. Chang, L.-J. Li and Y. Shi, ACS Nano, 2016, 10, 1454.
182. Q. Wu, H. Nong, R. Zheng, R. Zhang, J. Wang, L. Yang and B. Liu, Angew. Chem., Int. Ed., 2023, 62, e202301501.
183. L. Tang, J. Tan, H. Nong, B. Liu and H.-M. Cheng, Acc. Mater. Res., 2021, 2, 36.
184. J. J. P. Thompson, M. Gerhard, G. Witte and E. Malic, npj 2D Mater. Appl., 2023, 7, 69.
185. J. A. Schuller, S. Karaveli, T. Schiros, K. He, S. Yang, I. Kymissis, J. Shan and R. Zia, Nat. Nanotechnol., 2013, 8, 271.
186. H. Qiu, T. Xu, Z. Wang, W. Ren, H. Nan, Z. Ni, Q. Chen, S. Yuan, F. Miao, F. Song, G. Long, Y. Shi, L. Sun, J. Wang and X. Wang, Nat. Commun., 2013, 4, 2642.
187. H. Nan, Z. Wang, W. Wang, Z. Liang, Y. Lu, Q. Chen, D. He, P. Tan, F. Miao, X. Wang, J. Wang and Z. Ni, ACS Nano, 2014, 8, 5738.
188. P. D. Cunningham, K. M. McCreary, A. T. Hanbicki, M. Currie, B. T. Jonker and L. M. Hayden, J. Phys. Chem. C, 2016, 120, 5819.
189. S. S. Chou, M. De, J. Kim, S. Byun, C. Dykstra, J. Yu, J. Huang and V. P. Dravid, J. Am. Chem. Soc., 2013, 135, 4584.
190. K. Cho, M. Min, T.-Y. Kim, H. Jeong, J. Pak, J.-K. Kim, J. Jang, S. J. Yun, Y. H. Lee, W.-K. Hong and T. Lee, ACS Nano, 2015, 9, 8044.
191. Y. Liu, P. Stradins and S.-H. Wei, Angew. Chem., Int. Ed., 2016, 55, 965.
192. Z. Li, Y. Xiao, Y. Gong, Z. Wang, Y. Kang, S. Zu, P. M. Ajayan, P. Nordlander and Z. Fang, ACS Nano, 2015, 9, 10158.
193. M. Erementchouk, M. A. Khan and M. N. Leuenberger, Phys. Rev. B: Condens. Matter Mater. Phys., 2015, 92, 121401.
194. P. K. Chow, R. B. Jacobs-Gedrim, J. Gao, T.-M. Lu, B. Yu, H. Terrones and N. Koratkar, ACS Nano, 2015, 9, 1520.
195. B. Chen, H. Sahin, A. Suslu, L. Ding, M. I. Bertoni, F. M. Peeters and S. Tongay, ACS Nano, 2015, 9, 5326.
196. V. Vega-Mayoral, C. Backes, D. Hanlon, U. Khan, Z. Gholamvand, M. O'Brien, G. S. Duesberg, C. Gadermaier and J. N. Coleman, Adv. Funct. Mater., 2016, 26, 1028.
197. G. Guan, S. Zhang, S. Liu, Y. Cai, M. Low, C. P. Teng, I. Y. Phang, Y. Cheng, K. L. Duei, B. M. Srinivasan, Y. Zheng, Y.-W. Zhang and M.-Y. Han, J. Am. Chem. Soc., 2015, 137, 6152.
198. Y. Yu, Y. Yu, C. Xu, Y.-Q. Cai, L. Su, Y. Zhang, Y.-W. Zhang, K. Gundogdu and L. Cao, Adv. Funct. Mater., 2016, 26, 4733.
199. D. Voiry, A. Goswami, R. Kappera, C. D. C. C. E. Silva, D. Kaplan, T. Fujita, M. Chen, T. Asefa and M. Chhowalla, Nat. Chem., 2015, 7, 45.
200. M. Amani, P. Taheri, R. Addou, G. H. Ahn, D. Kiriya, D.-H. Lien, J. W. Ager, R. M. Wallace and A. Javey, Nano Lett., 2016, 16, 2786.
201. K. P. Dhakal, D. L. Duong, J. Lee, H. Nam, M. Kim, M. Kan, Y. H. Lee and J. Kim, Nanoscale, 2014, 6, 13028.
202. Y. Cai, H. Zhou, G. Zhang and Y.-W. Zhang, Chem. Mater., 2016, 28, 8611.
203. W. Chen, S. Chen, D. C. Qi, X. Y. Gao and A. T. S. Wee, J. Am. Chem. Soc., 2007, 129, 10418.
204. S. Zhang, H. M. Hill, K. Moudgil, C. A. Richter, A. R. Hight Walker, S. Barlow, S. R. Marder, C. A. Hacker and S. J. Pookpanratana, Adv. Mater., 2018, 30, 1802991.
205. Q. Zhao, W. Wang, F. Carrascoso-Plana, W. Jie, T. Wang, A. Castellanos-Gomez and R. Frisenda, Mater. Horiz., 2020, 7, 252.
206. S. Bertolazzi, S. Bonacchi, G. Nan, A. Pershin, D. Beljonne and P. Samorì, Adv. Mater., 2017, 29, 1606760.
207. K. Greben, S. Arora, M. G. Harats and K. I. Bolotin, Nano Lett., 2020, 20, 2544.
208. E. P. Nguyen, B. J. Carey, J. Z. Ou, J. van Embden, E. D. Gaspera, A. F. Chrimes, M. J. S. Spencer, S. Zhuiykov, K. Kalantar-zadeh and T. Daeneke, Adv. Mater., 2015, 27, 6225.
209. J. Azadmanjiri, P. Kumar, V. K. Srivastava and Z. Sofer, ACS Appl. Nano Mater., 2020, 3, 3116.
210. Y. Zhao, S. M. Gali, C. Wang, A. Pershin, A. Slassi, D. Beljonne and P. Samorì, Adv. Funct. Mater., 2020, 30, 2005045.
211. L. Yang, K. Majumdar, H. Liu, Y. Du, H. Wu, M. Hatzistergos, P. Y. Hung, R. Tieckelmann, W. Tsai, C. Hobbs and P. D. Ye, Nano Lett., 2014, 14, 6275.
212. J. Lu, A. Carvalho, X. K. Chan, H. Liu, B. Liu, E. S. Tok, K. P. Loh, A. H. Castro Neto and C. H. Sow, Nano Lett., 2015, 15, 3524.
213. S. Cho, S. Kim, J. H. Kim, J. Zhao, J. Seok, D. H. Keum, J. Baik, D.-H. Choe, K. J. Chang, K. Suenaga, S. W. Kim, Y. H. Lee and H. Yang, Science, 2015, 349, 625.
214. H. Yang, Y. Zhao, Q. Wen, Y. Mi, Y. Liu, H. Li and T. Zhai, Nano Res., 2021, 14, 4814.
215. P. K. Gogoi, Z. Hu, Q. Wang, A. Carvalho, D. Schmidt, X. Yin, Y.-H. Chang, L.-J. Li, C. H. Sow, A. H. C. Neto, M. B. H. Breese, A. Rusydi and A. T. S. Wee, Phys. Rev. Lett., 2017, 119, 077402.
216. H. Qiu, Z. Liu, Y. Yao, M. Herder, S. Hecht and P. Samorì, Adv. Mater., 2020, 32, 1907903.
217. X. Liu, D. Qu, L. Wang, M. Huang, Y. Yuan, P. Chen, Y. Qu, J. Sun and W. J. Yoo, Adv. Funct. Mater., 2020, 30, 2004880.
218. C. Wang, X. Wu, X. Zhang, G. Mu, P. Li, C. Luo, H. Xu and Z. Di, Appl. Phys. Lett., 2020, 116.
219. A. V. Kuklin, L. V. Begunovich, L. Gao, H. Zhang and H. Ågren, Phys. Rev. B, 2021, 104, 134109.
220. L. Daukiya, J. Teyssandier, S. Eyley, S. El Kazzi, M. C. Rodríguez González, B. Pradhan, W. Thielemans, J. Hofkens and S. De Feyter, Nanoscale, 2021, 13, 2972.
221. S. Presolski and M. Pumera, Mater. Today, 2016, 19, 140.
222. X. Chen, C. Bartlam, V. Lloret, N. Moses Badlyan, S. Wolff, R. Gillen, T. Stimpel-Lindner, J. Maultzsch, G. S. Duesberg, K. C. Knirsch and A. Hirsch, Angew. Chem., Int. Ed., 2021, 60, 13484.
223. R. Canton-Vitoria, Y. Sayed-Ahmad-Baraza, M. Pelaez-Fernandez, R. Arenal, C. Bittencourt, C. P. Ewels and N. Tagmatarchis, npj 2D Mater. Appl., 2017, 1, 13.
224. S. Chakraborty, J. Chattopadhyay, W. Guo and W. E. Billups, Angew. Chem., Int. Ed., 2007, 46, 4486.
225. J. M. Englert, C. Dotzer, G. Yang, M. Schmid, C. Papp, J. M. Gottfried, H.-P. Steinrück, E. Spiecker, F. Hauke and A. Hirsch, Nat. Chem., 2011, 3, 279.
226. Z. Syrgiannis, B. Gebhardt, C. Dotzer, F. Hauke, R. Graupner and A. Hirsch, Angew. Chem., Int. Ed., 2010, 49, 3322.
227. I. Gómez-Muñoz, S. Laghouati, R. Torres-Cavanillas, M. Morant-Giner, N. V. Vassilyeva, A. Forment-Aliaga and M. Giménez-Marqués, ACS Appl. Mater. Interfaces, 2021, 13, 36475.
228. M. Morant-Giner, R. Sanchis-Gual, J. Romero, A. Alberola, L. García-Cruz, S. Agouram, M. Galbiati, N. M. Padial, J. C. Waerenborgh, C. Martí-Gastaldo, S. Tatay, A. Forment-Aliaga and E. Coronado, Adv. Funct. Mater., 2018, 28, 1706125.
229. K. C. Knirsch, N. C. Berner, H. C. Nerl, C. S. Cucinotta, Z. Gholamvand, N. McEvoy, Z. Wang, I. Abramovic, P. Vecera, M. Halik, S. Sanvito, G. S. Duesberg, V. Nicolosi, F. Hauke, A. Hirsch, J. N. Coleman and C. Backes, ACS Nano, 2015, 9, 6018.
230. M. Lihter, M. Graf, D. Iveković, M. Zhang, T.-H. Shen, Y. Zhao, M. Macha, V. Tileli and A. Radenovic, ACS Appl. Nano Mater., 2021, 4, 1076.
231. K. Greulich, A. Belser, S. Bölke, P. Grüninger, R. Karstens, M. S. Sättele, R. Ovsyannikov, E. Giangrisostomi, T. V. Basova, D. Klyamer, T. Chassé and H. Peisert, J. Phys. Chem. C, 2020, 124, 16990.
232. S. Padgaonkar, S. H. Amsterdam, H. Bergeron, K. Su, T. J. Marks, M. C. Hersam and E. A. Weiss, J. Phys. Chem. C, 2019, 123, 13337.
233. J. Henzie, J. E. Barton, C. L. Stender and T. W. Odom, Acc. Chem. Res., 2006, 39, 249.
234. J. Wang, H. Yu, X. Zhou, X. Liu, R. Zhang, Z. Lu, J. Zheng, L. Gu, K. Liu, D. Wang and L. Jiao, Nat. Commun., 2017, 8, 377.
235. T. Yoshida, K. Watanabe, M. Petrović and M. Kralj, J. Phys. Chem. Lett., 2020, 11, 5248.
236. S. Fu, R. Wang, D. Tang, X. Zhang and D. He, ACS Appl. Mater. Interfaces, 2021, 13, 18372.
237. S. H. Amsterdam, T. LaMountain, T. K. Stanev, V. K. Sangwan, R. López-Arteaga, S. Padgaonkar, K. Watanabe, T. Taniguchi, E. A. Weiss, T. J. Marks, M. C. Hersam and N. P. Stern, J. Phys. Chem. Lett., 2021, 12, 26.
238. F. Ceballos, M. Z. Bellus, H.-Y. Chiu and H. Zhao, Nanoscale, 2015, 7, 17523.
239. B. Li, Z. Wan, C. Wang, P. Chen, B. Huang, X. Cheng, Q. Qian, J. Li, Z. Zhang, G. Sun, B. Zhao, H. Ma, R. Wu, Z. Wei, Y. Liu, L. Liao, Y. Ye, Y. Huang, X. Xu, X. Duan, W. Ji and X. Duan, Nat. Mater., 2021, 20, 818.
240. L. Meng, Z. Zhou, M. Xu, S. Yang, K. Si, L. Liu, X. Wang, H. Jiang, B. Li, P. Qin, P. Zhang, J. Wang, Z. Liu, P. Tang, Y. Ye, W. Zhou, L. Bao, H.-J. Gao and Y. Gong, Nat. Commun., 2021, 12, 809.
241. H. Liu and Y. Xue, Adv. Mater., 2021, 33, 2008456.
242. D. Capitano, Z. Hu, Y. Liu, X. Tong, D. Nykypanchuk, D. DiMarzio and C. Petrovic, ACS Omega, 2021, 6, 10537.
243. A. Kimel, A. Zvezdin, S. Sharma, S. Shallcross, N. de Sousa, A. García-Martín, G. Salvan, J. Hamrle, O. Stejskal, J. McCord, S. Tacchi, G. Carlotti, P. Gambardella, G. Salis, M. Münzenberg, M. Schultze, V. Temnov, I. V. Bychkov, L. N. Kotov, N. Maccaferri, D. Ignatyeva, V. Belotelov, C. Donnelly, A. H. Rodriguez, I. Matsuda, T. Ruchon, M. Fanciulli, M. Sacchi, C. R. Du, H. Wang, N. P. Armitage, M. Schubert, V. Darakchieva, B. Liu, Z. Huang, B. Ding, A. Berger and P. Vavassori, J. Phys. D: Appl. Phys., 2022, 55, 463003.
244. K. S. Novoselov, A. K. Geim, S. V. Morozov, D. Jiang, Y. Zhang, S. V. Dubonos, I. V. Grigorieva and A. A. Firsov, Science, 2004, 306, 666.
245. K. I. Bolotin, K. J. Sikes, J. Hone, H. L. Stormer and P. Kim, Phys. Rev. Lett., 2008, 101, 096802.
246. S. V. Morozov, K. S. Novoselov, M. I. Katsnelson, F. Schedin, D. C. Elias, J. A. Jaszczak and A. K. Geim, Phys. Rev. Lett., 2008, 100, 016602.
247. C. Qin, Y. Gao, Z. Qiao, L. Xiao and S. Jia, Adv. Opt. Mater., 2016, 4, 1429.
248. K. F. Mak, C. Lee, J. Hone, J. Shan and T. F. Heinz, Phys. Rev. Lett., 2010, 105, 136805.
249. A. K. Geim and I. V. Grigorieva, Nature, 2013, 499, 419.
250. C.-J. Park, H. J. Park, J. Y. Lee, J. Kim, C.-H. Lee and J. Joo, ACS Appl. Mater. Interfaces, 2018, 10, 29848.
251. H. J. Park, C.-J. Park, J. Y. Kim, M. S. Kim, J. Kim and J. Joo, ACS Appl. Mater. Interfaces, 2018, 10, 32556.
252. S. P. Tiwari, R. Verma, M. B. Alam, R. Kumari, O. P. Sinha and R. Srivastava, Vacuum, 2017, 146, 474.
253. S. Andleeb, X. Wang, H. Dong, S. Valligatla, C. N. Saggau, L. Ma, O. G. Schmidt and F. Zhu, Nanomaterials, 2023, 13, 1491.
254. J. Pak, J. Jang, K. Cho, T.-Y. Kim, J.-K. Kim, Y. Song, W.-K. Hong, M. Min, H. Lee and T. Lee, Nanoscale, 2015, 7, 18780.
255. G. Ghimire, K. P. Dhakal, G. P. Neupane, S. Gi Jo, H. Kim, C. Seo, Y. Hee Lee, J. Joo and J. Kim, Nanotechnology, 2017, 28, 185702.
256. W. L. Chow, P. Yu, F. Liu, J. Hong, X. Wang, Q. Zeng, C.-H. Hsu, C. Zhu, J. Zhou, X. Wang, J. Xia, J. Yan, Y. Chen, D. Wu, T. Yu, Z. Shen, H. Lin, C. Jin, B. K. Tay and Z. Liu, Adv. Mater., 2017, 29, 1602969.
257. S. Dey, H. S. S. R. Matte, S. N. Shirodkar, U. V. Waghmare and C. N. R. Rao, Chem. – Asian J., 2013, 8, 1780.
258. S. Andleeb, A. Kumar Singh and J. Eom, Sci. Technol. Adv. Mater., 2015, 16, 035009.
259. Y. Du, H. Liu, A. T. Neal, M. Si and P. D. Ye, IEEE Electron. Dev. Lett., 2013, 34, 1328.
260. D. M. Sim, M. Kim, S. Yim, M.-J. Choi, J. Choi, S. Yoo and Y. S. Jung, ACS Nano, 2015, 9, 12115.
261. K. Heo, S.-H. Jo, J. Shim, D.-H. Kang, J.-H. Kim and J.-H. Park, ACS Appl. Mater. Interfaces, 2018, 10, 32765.
262. D.-H. Kang, M.-S. Kim, J. Shim, J. Jeon, H.-Y. Park, W.-S. Jung, H.-Y. Yu, C.-H. Pang, S. Lee and J.-H. Park, Adv. Funct. Mater., 2015, 25, 4219.
263. A. Tarasov, S. Zhang, M.-Y. Tsai, P. M. Campbell, S. Graham, S. Barlow, S. R. Marder and E. M. Vogel, Adv. Mater., 2015, 27, 1175.
264. C. J. L. D. L. Rosa, R. Phillipson, J. Teyssandier, J. Adisoejoso, Y. Balaji, C. Huyghebaert, I. Radu, M. Heyns, S. D. Feyter and S. D. Gendt, Appl. Phys. Lett., 2016, 109, 253112.
265. B. Tang, Z. G. Yu, L. Huang, J. Chai, S. L. Wong, J. Deng, W. Yang, H. Gong, S. Wang, K.-W. Ang, Y.-W. Zhang and D. Chi, ACS Nano, 2018, 12, 2506.
266. L. Qi, Y. Wang, L. Shen and Y. Wu, Appl. Phys. Lett., 2016, 108, 063103.
267. B. Lei, Z. Hu, D. Xiang, J. Wang, G. Eda, C. Han and W. Chen, Nano Res., 2017, 10, 1282.
268. H. Im, A. Bala, B. So, Y. J. Kim and S. Kim, Adv. Electron. Mater., 2021, 7, 2100644.
269. B. Radisavljevic, A. Radenovic, J. Brivio, V. Giacometti and A. Kis, Nat. Nanotechnol., 2011, 6, 147.
270. W. S. Hwang, M. Remskar, R. Yan, V. Protasenko, K. Tahy, S. D. Chae, P. Zhao, A. Konar, H. Xing, A. Seabaugh and D. Jena, Appl. Phys. Lett., 2012, 101, 013107.
271. R. Kroon, D. Kiefer, D. Stegerer, L. Yu, M. Sommer and C. Müller, Adv. Mater., 2017, 29, 1700930.
272. G. M. Rangger, O. T. Hofmann, L. Romaner, G. Heimel, B. Bröker, R.-P. Blum, R. L. Johnson, N. Koch and E. Zojer, Phys. Rev. B: Condens. Matter Mater. Phys., 2009, 79, 165306.
273. Y. Krupskaya, M. Gibertini, N. Marzari and A. F. Morpurgo, Adv. Mater., 2015, 27, 2453.
274. J. Lin, J. Zhong, S. Zhong, H. Li, H. Zhang and W. Chen, Appl. Phys. Lett., 2013, 103, 063109.
275. Z. Yu, Y. Pan, Y. Shen, Z. Wang, Z.-Y. Ong, T. Xu, R. Xin, L. Pan, B. Wang, L. Sun, J. Wang, G. Zhang, Y. W. Zhang, Y. Shi and X. Wang, Nat. Commun., 2014, 5, 5290.
276. R. Kappera, D. Voiry, S. E. Yalcin, B. Branch, G. Gupta, A. D. Mohite and M. Chhowalla, Nat. Mater., 2014, 13, 1128.
277. V. K. Sangwan, D. Jariwala, I. S. Kim, K.-S. Chen, T. J. Marks, L. J. Lauhon and M. C. Hersam, Nat. Nanotechnol., 2015, 10, 403.
278. D. Jariwala, V. K. Sangwan, C.-C. Wu, P. L. Prabhumirashi, M. L. Geier, T. J. Marks, L. J. Lauhon and M. C. Hersam, Proc. Natl. Acad. Sci. U. S. A., 2013, 110, 18076.
279. M. M. Furchi, A. Pospischil, F. Libisch, J. Burgdörfer and T. Mueller, Nano Lett., 2014, 14, 4785.
280. C.-H. Lee, G.-H. Lee, A. M. van der Zande, W. Chen, Y. Li, M. Han, X. Cui, G. Arefe, C. Nuckolls, T. F. Heinz, J. Guo, J. Hone and P. Kim, Nat. Nanotechnol., 2014, 9, 676.
281. F. Xiong, In Encyclopedia of RF and Microwave Engineering, John Wiley & Sons, Inc., 2005.
282. J. Sun, Y. Choi, Y. J. Choi, S. Kim, J.-H. Park, S. Lee and J. H. Cho, Adv. Mater., 2019, 31, 1803831.
283. P. C. Y. Chow and T. Someya, Adv. Mater., 2020, 32, 1902045.
284. S. Mukherjee, D. Bhattacharya, S. Patra, S. Paul, R. K. Mitra, P. Mahadevan, A. N. Pal and S. K. Ray, ACS Appl. Mater. Interfaces, 2022, 14, 5775.
285. J. Ji and J. H. Choi, Nanoscale, 2022, 14, 10648.
286. K. Cho, J. Pak, S. Chung and T. Lee, ACS Nano, 2019, 13, 9713.
287. B. Zhao, Z. Gan, M. Johnson, E. Najafidehaghani, T. Rejek, A. George, R. H. Fink, A. Turchanin and M. Halik, Adv. Funct. Mater., 2021, 31, 2105444.
288. T. Schmaltz, A. Y. Amin, A. Khassanov, T. Meyer-Friedrichsen, H.-G. Steinrück, A. Magerl, J. J. Segura, K. Voitchovsky, F. Stellacci and M. Halik, Adv. Mater., 2013, 25, 4511.
289. D. Rhodes, S. H. Chae, R. Ribeiro-Palau and J. Hone, Nat. Mater., 2019, 18, 541.
290. O. Sato, Nat. Chem., 2016, 8, 644.
291. H. Koshima, N. Ojima and H. Uchimoto, J. Am. Chem. Soc., 2009, 131, 6890.
292. S. Kobatake, S. Takami, H. Muto, T. Ishikawa and M. Irie, Nature, 2007, 446, 778.
293. K. Kang, K.-H. Lee, Y. Han, H. Gao, S. Xie, D. A. Muller and J. Park, Nature, 2017, 550, 229.
294. L. J. McGilly, A. Kerelsky, N. R. Finney, K. Shapovalov, E.-M. Shih, A. Ghiotto, Y. Zeng, S. L. Moore, W. Wu, Y. Bai, K. Watanabe, T. Taniguchi, M. Stengel, L. Zhou, J. Hone, X. Zhu, D. N. Basov, C. Dean, C. E. Dreyer and A. N. Pasupathy, Nat. Nanotechnol., 2020, 15, 580.
295. K. S. Mali, J. Greenwood, J. Adisoejoso, R. Phillipson and S. De Feyter, Nanoscale, 2015, 7, 1566.
296. S. Carr, S. Fang and E. Kaxiras, Nat. Rev. Mater., 2020, 5, 748.
297. R. Bistritzer and A. H. MacDonald, Proc. Natl. Acad. Sci. U. S. A., 2011, 108, 12233.
298. L. Zhang, R. Ni and Y. Zhou, J. Appl. Phys., 2023, 133.
299. Y. Zhong, B. Cheng, C. Park, A. Ray, S. Brown, F. Mujid, J.-U. Lee, H. Zhou, J. Suh, K.-H. Lee, A. J. Mannix, K. Kang, S. J. Sibener, D. A. Muller and J. Park, Science, 2019, 366, 1379.
300. L. Tong, Z. Peng, R. Lin, Z. Li, Y. Wang, X. Huang, K.-H. Xue, H. Xu, F. Liu, H. Xia, P. Wang, M. Xu, W. Xiong, W. Hu, J. Xu, X. Zhang, L. Ye and X. Miao, Science, 2021, 373, 1353.
301. Z. Hao, H. Wang, S. Jiang, J. Qian, X. Xu, Y. Li, M. Pei, B. Zhang, J. Guo, H. Zhao, J. Chen, Y. Tong, J. Wang, X. Wang, Y. Shi and Y. Li, Adv. Sci., 2022, 9, 2103494.
302. L. Sun, W. Wang and H. Yang, Adv. Intell. Syst., 2020, 2, 1900167.
303. J. Zhang, S. Dai, Y. Zhao, J. Zhang and J. Huang, Adv. Intell. Syst., 2020, 2, 1900136.
304. J. Ghosh, S. Parveen, P. J. Sellin and P. K. Giri, Adv. Mater. Technol., 2023, 8, 2300400.
305. M. Lee, W. Lee, S. Choi, J.-W. Jo, J. Kim, S. K. Park and Y.-H. Kim, Adv. Mater., 2017, 29, 1700951.
306. A. A. Bessonov, M. N. Kirikova, D. I. Petukhov, M. Allen, T. Ryhänen and M. J. A. Bailey, Nat. Mater., 2015, 14, 199.
307. T. Ohno, T. Hasegawa, T. Tsuruoka, K. Terabe, J. K. Gimzewski and M. Aono, Nat. Mater., 2011, 10, 591.
308. X. Hou, H. Zhang, C. Liu, S. Ding, W. Bao, D. W. Zhang and P. Zhou, Small, 2018, 14, 1800319.
309. C. Liu, X. Yan, X. Song, S. Ding, D. W. Zhang and P. Zhou, Nat. Nanotechnol., 2018, 13, 404.
310. D. Sarkar, J. Tao, W. Wang, Q. Lin, M. Yeung, C. Ren and R. Kapadia, ACS Nano, 2018, 12, 1656.
311. Y. Yang, Y. He, S. Nie, Y. Shi and Q. Wan, IEEE Electron Dev. Lett., 2018, 39, 897.
312. Y. Yang, J. Wen, L. Guo, X. Wan, P. Du, P. Feng, Y. Shi and Q. Wan, ACS Appl. Mater. Interfaces, 2016, 8, 30281.
313. S. Qin, F. Wang, Y. Liu, Q. Wan, X. Wang, Y. Xu, Y. Shi, X. Wang and R. Zhang, 2D Mater., 2017, 4, 035022.
314. J. Zhu, Y. Yang, R. Jia, Z. Liang, W. Zhu, Z. U. Rehman, L. Bao, X. Zhang, Y. Cai, L. Song and R. Huang, Adv. Mater., 2018, 30, 1800195.
315. C. S. Yang, D. S. Shang, N. Liu, G. Shi, X. Shen, R. C. Yu, Y. Q. Li and Y. Sun, Adv. Mater., 2017, 29, 1700906.
316. Y. van de Burgt, E. Lubberman, E. J. Fuller, S. T. Keene, G. C. Faria, S. Agarwal, M. J. Marinella, A. Alec Talin and A. Salleo, Nat. Mater., 2017, 16, 414.
317. H.-K. He, R. Yang, W. Zhou, H.-M. Huang, J. Xiong, L. Gan, T.-Y. Zhai and X. Guo, Small, 2018, 14, 1800079.
318. Y. Wang, Z. Lv, J. Chen, Z. Wang, Y. Zhou, L. Zhou, X. Chen and S.-T. Han, Adv. Mater., 2018, 30, 1802883.
319. R. A. John, F. Liu, N. A. Chien, M. R. Kulkarni, C. Zhu, Q. Fu, A. Basu, Z. Liu and N. Mathews, Adv. Mater., 2018, 30, 1800220.
320. H. Tian, Q. Guo, Y. Xie, H. Zhao, C. Li, J. J. Cha, F. Xia and H. Wang, Adv. Mater., 2016, 28, 4991.
321. X. Wang, H. Tian, S. Shen, J. Wang, Y. Li, Y. Pang, Y. Yang and T. L. Ren, IEEE Trans. Electron Dev., 2018, 65, 3543.
322. H. Tian, W. Mi, X.-F. Wang, H. Zhao, Q.-Y. Xie, C. Li, Y.-X. Li, Y. Yang and T.-L. Ren, Nano Lett., 2015, 15, 8013.
323. P. Gkoupidenis, N. Schaefer, B. Garlan and G. G. Malliaras, Adv. Mater., 2015, 27, 7176.
324. Y. Shi, X. Liang, B. Yuan, V. Chen, H. Li, F. Hui, Z. Yu, F. Yuan, E. Pop, H. S. P. Wong and M. Lanza, Nat. Electron., 2018, 1, 458.
325. X. Yan, L. Zhang, H. Chen, X. Li, J. Wang, Q. Liu, C. Lu, J. Chen, H. Wu and P. Zhou, Adv. Funct. Mater., 2018, 28, 1803728.
326. M. T. Sharbati, Y. Du, J. Torres, N. D. Ardolino, M. Yun and F. Xiong, Adv. Mater., 2018, 30, 1802353.
327. V. K. Sangwan, H.-S. Lee, H. Bergeron, I. Balla, M. E. Beck, K.-S. Chen and M. C. Hersam, Nature, 2018, 554, 500.
328. G. Indiveri, E. Chicca and R. Douglas, IEEE Trans. Neural Networks, 2006, 17, 211.
329. R. Iyer, V. Menon, M. Buice, C. Koch and S. Mihalas, PLOS Comput. Biol., 2013, 9, e1003248.
330. Y. Yang, B. Chen and W. D. Lu, Adv. Mater., 2015, 27, 7720.
331. S. Kim, C. Du, P. Sheridan, W. Ma, S. Choi and W. D. Lu, Nano Lett., 2015, 15, 2203.
332. O. Lopez-Sanchez, D. Lembke, M. Kayci, A. Radenovic and A. Kis, Nat. Nanotechnol., 2013, 8, 497.
333. W. Zhang, J.-K. Huang, C.-H. Chen, Y.-H. Chang, Y.-J. Cheng and L.-J. Li, Adv. Mater., 2013, 25, 3456.
334. D.-S. Tsai, K.-K. Liu, D.-H. Lien, M.-L. Tsai, C.-F. Kang, C.-A. Lin, L.-J. Li and J.-H. He, ACS Nano, 2013, 7, 3905.
335. J. Schornbaum, B. Winter, S. P. Schießl, F. Gannott, G. Katsukis, D. M. Guldi, E. Spiecker and J. Zaumseil, Adv. Funct. Mater., 2014, 24, 5798.
336. D. Kufer, I. Nikitskiy, T. Lasanta, G. Navickaite, F. H. L. Koppens and G. Konstantatos, Adv. Mater., 2015, 27, 176.
337. D. Kufer and G. Konstantatos, Nano Lett., 2015, 15, 7307.
338. M. M. Furchi, D. K. Polyushkin, A. Pospischil and T. Mueller, Nano Lett., 2014, 14, 6165.
339. D. Kufer and G. Konstantatos, ACS Photonics, 2016, 3, 2197.
340. T. Mueller, F. Xia and P. Avouris, Nat. Photonics, 2010, 4, 297.

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