Adsorption and epitaxial growth of small organic semiconductors on hexagonal boron nitride

So far, there are three main approaches to fabricate hBN substrates. First, there is mechanical exfoliation in a similar way as introduced for Gr by Novoselov and Geim [35, 36, 120]. The second is epitaxial hBN growth by a CVD method on various metal catalyst surfaces [36, 39, 121–126]. The third technique is liquid exfoliation [127] which allows larger scale fabrication of hBN nanosheets. Besides the established main techniques also novel approaches are explored like ion-intercalation assisted exfoliation [128], physical vapor deposition [129], surface segregation-based methods [130], pulsed laser deposition [131], pyrolysis [132], or unwrapping of BN nanotubes [36, 133].

Mechanic exfoliation can be done from larger hBN single crystals [41] or from crystalline powder [35, 134]. The results are usually flakes with strongly varying vertical and lateral dimensions. In vertical dimension, the flakes can be single layer, few layers, or several 100 or 1000 layers thick. The lateral dimensions usually range from a few µm to tens of µm as can be seen in figures 1(c)–(e)). As support for the flakes often SiO₂/Si wafers are used with thermal oxide thickness of either 80 nm or 285 nm. These oxide thicknesses are chosen to obtain a sufficiently high optical contrast making even single layered flakes visible by optical microscopy. Even though direct mechanic exfoliation is a comparably fast and simple way to receive thin hBN flakes it needs practice to be routinely applied. A major drawback of the method is the poor control over the flake dimensions, which complicates a targeted substrate preparation. However, the obtained flakes are usually pure and of high crystalline quality with minimal wrinkling.

A way for large-scale fabrication of 2D hBN nanosheets is liquid exfoliation [135, 136]. Within this approach boron nitride powder is sonicated in an appropriate solvent that minimizes the energy of exfoliation. For hBN, a variety of possible solvents facilitating mechano-chemical exfoliation but also flake's functionalization has been reported [135–143].

As classical surface science technique, epitaxial growth by CVD is a highly controllable and scalable process, to obtain large-area, single or multilayer hBN. In the process, a gaseous precursor (e.g. borazin (B₃H₆N₃), and ammonia borane (NH₃-BH₃)) is brought into contact with a catalyst surface which triggers the chemical reaction and acts as epitaxial template. The result is usually a large-area polycrystalline hBN sheet which is in intimate contact with the catalyst and needs to be subsequently separated and transferred to the desired substrate. As catalysts, polycrystalline materials like Cu foils [39], Pt foils [59], or Ni films [121] have been used in flow reactors or furnaces. Well-ordered monolayer hBN films can be generated by CVD on single crystalline substrates like for example Ni(1 1 1) [144], Rh(1 1 1) [54], Ir(1 1 1) [125], Cu(1 1 1) [145], or others [36] under UHV conditions. In case a transfer is desired, the CVD grown hBN film is covered by a polymer layer, usually polymethylmethacrylate (PMMA), for protection and handling. The hBN/PMMA assembly is then removed from the catalyst either by etching away of the catalyst or by electrochemical delamination [58, 59]. After that the hBN is brought into intimate contact with the desired substrate—most often SiO₂/Si- and the PMMA is removed by dissolving it in acetone.

Basically, all classical vapor based or solution based deposition processes can be applied to deposit organic molecules on the hBN surface. The main methods applied in the framework of this article are vapor deposition techniques like organic molecular beam epitaxy (OMBE) [146], hot wall epitaxy (HWE) [147], and vapor deposition in tube furnaces with and without carrier gases. OMBE—operated under ultra-high vacuum (UHV) conditions—is suitable to controllably deposit small amounts of molecules, which is necessary for studying very small coverages and for the development of 2D molecular networks.

HWE is a vapor deposition technique operated under high vacuum conditions (~10⁻⁶ mbar) that potentially allows to grow very close to thermodynamic equilibrium conditions and turned out to provide very good results for crystalline organic thin films [11–13, 148]. The closed bottom of a semi open, vertical quartz tube is filled with the organic material in powder form. The open side of the tube is closed by the substrate. The tube bottom, the substrate, and the tube wall can be heated independently. In this way, the evaporation rate, vapor temperature, and substrate temperature can be adjusted separately for optimal growth conditions.

Closely related to HWE is the vapor deposition technique in furnaces or tube reactors when operated without carrier gas [14, 104]. There, the source material and the substrate are arranged adjacent but at some distance to each other in a horizontal tube. A heating system generates a temperature gradient, where the temperature is highest at the position of the source material and lowest at the tube endings. The organic vapor deposits on the substrate which is at lower temperature than the source. The growth parameters are optimized by controlling the temperature and adjusting the distance between source material and substrate. Often, carrier gases like Ar, N₂, H₂, or mixtures of these are used that flow from the source region towards the substrate [149–153].

Solution based deposition can happen via immersion, drop casting, spin coating, etc, of solutions which contain organic molecules [154]. The film properties can be adjusted by the constituent concentration, choice of solvent, immersion time, and temperature. Solution based techniques are insofar interesting since they do not need elaborate vacuum equipment, which saves time and costs.

For insulating thicker hBN layers on metals supports or dielectric substrates, atomic force microscopy (AFM) [155, 156] methods are often applied. Unlike STM, which does not work on insulators, the AFM relies on a force feedback which also works on insulating surfaces. Most AFM systems operate under ambient conditions and are used to resolve structure sizes from several 10 nm to several 10 µm. Since AFM systems providing molecular resolution are meanwhile commercially available they are frequently used to determine molecular arrangements and surface unit cells. Thus, they play a similar role on insulating surfaces as STM does on conductive systems.

For epitaxial, single-layer hBN still attached to the metal substrate, scanning tunneling microscopy (STM) [157] is a frequently used standard method to characterize molecular assembling on hBN at molecular and atomic resolution [24, 50]. It also allows to perform local manipulations by application of voltage pulses, which can be used to trigger and to investigate local chemical reactions between the molecules on the surface [75, 76, 158].

Due to the limited thickness of the organic films, structural investigations using x-rays are often performed using gracing incidence x-ray diffraction (GIXD) [159, 160]. In GIXD, a collimated x-ray beam is directed at an incidence angle  <1° to the sample surface. In order to determine the in-plane structure, the sample is rotated in-plane and the diffracted intensity is recorded as a function of rotation angle. For organic mono or few-layer samples, GIXD is often performed at synchrotron radiation sources to realize sufficient signal intensity [159, 161] but can also be performed with standard laboratory equipment [162]. However, care has to be taken when using highly intense x-ray beams at synchrotrons, since they might cause damage of the relatively sensitive organic films [163].

Angle-dependent near edge x-ray adsorption fine structure spectroscopy (NEXAFS) is a method to determine the local geometry and chemical state of atoms or molecules adsorbed on surfaces [164]. Since for NEXAFS the energy of the incident x-ray has to be tuned to track an energy window around a specific absorption edge, a synchrotron light source is necessary. For organic materials, often the carbon K-edge is investigated (280–300 eV). Intensity changes as a response to a variation of the beam incidence angle can be used to determine the molecular tilt angle with respect to the substrate [165].

Transmission electron microscopy (TEM) is a powerful technique for structure characterization with potential atomic-resolution imaging in real space [166]. In order to enable electron transmission for TEM measurements, specimens have to be thinner than ~100 nm (material dependent). Thin organic layers on hBN flakes are naturally thin samples which only can be transferred to, or prepared directly on the TEM grid without further treatment. In this way, selected area electron diffraction (SEAD) patterns can be recorded, which is useful for the determination of local unit cell parameters and orientations [49]. However, organic films are comparably sensitive and electron beam damage effects can limit the reliable use of TEM for organic materials [167].

The majority of the techniques mentioned so far mainly aim on structural/geometrical properties. They can be complemented by techniques which are sensitive to chemical changes and (opto) electronic properties. A frequently used standard tool for qualitative and quantitative chemical analysis of surfaces is x-ray photoelectron spectroscopy (XPS) [168] which can be performed with standard lab apparatus or at a synchrotron if very high energy resolution or high sensitivity are necessary. For the measurements UHV is required. X-rays with energies around 1.5 keV for lab sources (for synchrotrons up to 20 keV) generate photoelectrons from core levels which are analyzed by their kinetic energy which is specific for their binding energy. Chemical interaction can cause core level shifts, which are then detected. Sensitive to binding configurations and chemical surrounding is also Raman spectroscopy [169]. Raman spectroscopy relies on inelastic scattering of monochromatic light at molecules or solids. Laser light is focused to the sample and interacts with molecular vibrations. Laser photons can change their energy by either exciting or absorbing phonons which is detected as frequency shift in the scattered light. From the spectrum potentially parameters like crystallinity, crystal orientation, chemical composition, strain, temperature, doping, etc, can be deduced.

Also light absorption spectra recorded in the ultra-violet, visible and near-infrared region (UV–vis-NIR) can be used to gather information about the electronic properties of the organic compounds. The wavelength dependent absorption provides information about important parameters like the energy gap between the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO) [107]. For example, from the comparison of the HOMO–LUMO gap measured in solution or gas phase with the gap of the adsorbed material conclusions about the binding strength can be drawn [170, 171]. It is also a powerful tool when polarized light is used. Many organic molecules are anisotropic in shape and light absorption. Polarization direction dependent absorption provides information about preferential molecular orientation and can reveal domain structures [172].

In order to properly interpret experimental data, density functional theory (DFT) [173] is widely applied. DFT has been very successfully applied to many solid state problems. DFT is especially powerful for the calculation of equilibrium structures and, for example, has been successfully applied to inorganic thin films on metal surfaces [174, 175]. However, when dealing with organic molecules on surfaces like hBN a special approach is needed to account for vdW interactions, e.g. by employing the Tkatchenko–Scheffler [176, 177] or the Grimme schemes [178]. Structural models, adsorption geometries, or spatial arrangement of molecular orbitals can be simulated. Different scenarios/structural models can be tested, and the results are then compared with the experimental evidence. DFT is especially useful to identify preferential adsorption geometries of individual molecules by comparing the corresponding binding energies of possible adsorption configurations. This can help to understand the alignment of the evolving molecular structures [12, 14, 104]. Further, analysis of the orbital overlap helps to understand the structural-electrical conductivity relation [14].

An exceptionally high field effect mobility of µ ~ 11.5 cm² V⁻¹ s⁻¹ has been achieved for high-quality rubrene films grown on exfoliated hBN [161]. This value lies already within the range of free standing rubrene single crystal FETs (µ ~ 15 cm² V⁻¹ s⁻¹) [179] and clearly above those of rubrene thin film FETs with conventional SiO₂ gate dielectric (µ ~ 8  ×  10⁻⁶ cm² V⁻¹ s⁻¹ [180]). In this case, vdW-Gr bottom contacts were employed in order to ensure minimum distortion of the film and low contact resistance. The values almost reached those of devices with high quality single crystal rubrene suspended over an air-gap [179], which was attributed to the atomically sharp, low trap density hBN/rubrene interface. As substrates exfoliated hBN flakes transferred on SiO₂/Si supports were used. For device fabrication, smooth flakes of uniform height of 10–60 nm and lateral sizes larger than 10 µm were selected. 500–1000 nm thick rubrene films were grown by a vapor transport method at a substrate temperature of ~450 K and a growth rate between 16 and 33 nm s⁻¹ [161]. The resulting rubrene films formed large single crystalline domains essentially limited only by the flake size. A structural model of the rubrene molecular arrangement on hBN is provided in figures 3(a) and (b). AFM revealed the existence of atomically smooth, uniform terraces with a typical step height of (1.4  ±  0.4) nm (figure 3(c)), which well matches the molecular length of the rubrene c-axis as sketched in figure 3(a). TEM and SAED confirmed the crystallinity of hBN and rubrene layers as shown in figure 3(d), and indicated that no amorphous interface layer between hBN and rubrene is formed. GIXD revealed the growth of an orthorhombic rubrene crystal with the c-axis perpendicular to the hBN plane. The deduced lattice parameters were found to be very close to those of a free standing rubrene crystal indicating that the intermolecular interactions dominate over the molecule/hBN interactions. The 6-fold symmetry of the GIXD patterns seems to emerge from 2-fold rubrene domains epitaxially locked to the directions of the 3-fold symmetric hBN crystal. Apparently the weak rubrene/hBN interaction is still sufficiently strong to promote epitaxial ordering of rubrene on hBN, but not strong enough to severely distort the rubrene lattice. GIXD and SAED revealed that the a-axes of hBN and the rubrene lattices are rotated by ~3° (GIXD) to ~4° (SAED) with respect to each other. First-principle binding-energy calculations for a single rubrene layer on hBN as a function of the relative rotation angle between their a-axes revealed—consistently with the experimental data—energy minima at rotation angles smaller than 10°. Calculations including vdW interactions yielded optimum misorientation angles of ~6° which corresponds to reduction in energy of 0.2 eV per molecule in comparison to the computations excluding vdW contributions [161]. This further indicates that the weak vdW interaction can be important for epitaxial growth.

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The fact that the vdW forces between molecule and substrate are often significantly weaker than the mutual molecular interaction can be exploited to obtain quasi free standing molecular crystals with minimum interference from the substrate. This way, high-quality few layer thick films of dioctylbenzothienobenzothiophene C₈-BTBT exhibiting record charge carrier mobilities have been fabricated on Gr and hBN [104, 106]. In this case, exfoliated hBN substrates transferred onto SiO₂/Si supports (285 nm oxide thickness) were used. The C₈-BTBT crystals were grown in a tube furnace at ~5.3  ×  10⁻⁶ mbar without carrier gas by heating C₈-BTBT powder to 100 °C–120 °C. The substrates were just positioned a few inches away from the source. Unfortunately, the sample temperature is not specified, but should lie below the source temperature and above room temperature. To study the evolution of the film, the growth process was interrupted in order to perform ex-situ AFM measurements at different growth stages. The thickness of the initial three layers were determined to be ~0.53 nm (interface layer), ~1.68 nm (first layer) and ~2.85 nm (second layer) as shown in figures 4(b)–(d). The subsequent higher layers exhibited slightly higher heights around ~3 nm which corresponds well with the layer distance specified for C₈-BTBT bulk (~2.9 nm) [101, 105, 181]. The interface layer with a thickness of only ~0.53 nm indicates a molecular packing different from the bulk. STM studies on the C₈-BTBT interface layer on Gr, which behaves very similar as the C₈-BTBT/hBN system, suggest molecular packing in a rectangular lattice with d₁  =  2.52 nm and d₂  =  0.66 nm [104]. Corresponding DFT calculations resulted in an optimized geometry for a single C₈-BTBT on Gr where the alkyl chains are almost parallel to the Gr and the molecule exhibits a leaning angle of ~35° with respect to the 'stand-up' configuration. (In the 'stand-up' configuration, the benzothiophene is parallel to the surface whereas the alkyl chains stick out of the Gr plane with a height of ~1 nm.) In this configuration, the CH-π [182] and π-π [183] interactions are maximized. A fully relaxed structure including molecule–molecule interaction resulted in a tilt angle of the benzothiophene plane with respect to Gr of only 10° and a lattice with d₁  =  2.47 nm and d₂  =  0.64 nm and a layer thickness of ~0.52 nm which coincided perfectly with the STM data. A schematic representation of the approximate position of the C₈-BTBT molecules of the IL with respect to the hBN is shown in figure 4(a)). Interestingly, the alkyl chains tend to align roughly with the hBN's zigzag direction. A perfect alignment of the alkyl chain with the zigzag direction was found for the adsorption of hexacontane on hBN [184] (see later). High-resolution AFM images (figures 4(e) and (f)) of the first and second layer growing on top of the interface layer revealed a monoclinic crystal structure with herring-bone stacking similar to the bulk configuration. The lattice parameters were found to be a ~ 0.67 nm and b ~ 0.79 nm which is slightly different from the bulk values. Further, the molecules in the first layer appeared to be more inclined as would be expected for the bulk [102]. Apparently, the substrate's influence is already reduced in the first layer and already negligible for the second layer. That is also reflected in the calculated binding energy of an individual molecule in the first layer on top of the interface layer which is 0.048 eV for the hBN substrate. The corresponding intermolecular binding energy in the first layer is 0.704 eV. That clearly demonstrates again that the influence of the substrate (interface layer) is very small compared to the intermolecular interaction. This fact is also reflected in the morphology where a small nucleation density and large islands with smooth rims were observed indicating high diffusion rates on the substrate and along the island rims. The high crystalline quality obtained on the smooth and only weakly interacting hBN substrate is also evident in the charge carrier mobility, which is very sensitive to structural defects. Mobility measurements were performed on C₈-BTBT OFETs with hBN gate dielectric and transferred Au top contacts as a function of film thickness and temperature [104, 106]. Even though the C₈-BTBT channel only comprised from the interface layer and the first layer, field effect mobility values of up to µ ~ 10 cm² V⁻¹ s⁻¹ were reached [104]. These high values basically correspond to monolayer devices since it can be assumed that the interface layer is non-contributing. The observed field effect mobility outperforms typical values found for monolayer devices by orders of magnitude (µ ~ 10⁻⁶–10⁻¹ cm² V⁻¹ s⁻¹). Measurements of the intrinsic hole mobility utilizing a gated four-point probe technique revealed values up to µ ~ 30 cm² V⁻¹ s⁻¹ for first layer devices [106]. After optimization of the contacts by insertion of a doped Gr layer, C₈-BTBT OFET devices showed clear band like transport and superior characteristics. Interestingly, devices with a higher number of C₈-BTBT layers performed worse, which was attributed to the low charge transport between the layers [106]. However, the high mobility values obtained are a consequence of the high crystal quality and large domain sizes. It could be further shown that C₈-BTBT preferentially grows in a layer-by-layer fashion on hBN and that an out-of-plane roughness just larger than 0.6 nm is sufficient to trigger additional nucleation [185]. At coverages beyond two monolayers, the roughness is already significantly larger than 0.6 nm resulting in increased nucleation and smaller domain sizes which could reduce mobility.

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However, based on the C₈-BTBT/hBN OFET system [104, 106] it could convincingly be demonstrated that these kinds of single layer devices are advantageous because of the direct contact with the charge transport layer, the conserved intrinsic mobility and the minimized parasitic contact resistance.

Exfoliated hBN has also been used as vdW substrate for the growth of high-quality, few layer pentacene (5A) in order to study 5A's intrinsic electrical properties in the 2D limit [14, 65]. For this study, hBN was especially suitable because, first, it is an atomically smooth vdW material with low defect density and, second, it has a comparably low dielectric constant which yields only weak polaronic coupling. The smooth vdW surface is crucial for the growth of high-quality crystals with little disturbance from the interface, thus minimizing disorder, charge trapping, etc. The low polaronic coupling at the organic/hBN interface reduces the possible Fröhlich polaron [186] related reduction of the mobility in the film [187, 188], which might cause deviations from the intrinsic values.

Sample preparation was performed the same way as explained for C₈-BTBT. Just the furnace temperature was 130 °C–160 °C in this case. A typical growth time for a 5–10 layer thick 5A crystal was 3 h. Analogous to C₈-BTBT, the procedure resulted in layered crystals as confirmed by AFM and TEM investigations. Here, again an interface layer of only 0.5 nm height followed by two layers with heights of 1.14 nm (first layer) and 1.58 nm (second layer) occurs (see figures 5(a)–(d)). All subsequent layers had the same height as the second layer which corresponded well to the terrace height in the 5A thin film phase [189]. The small height of the interface layer implies a flat-lying configuration of 5A with its long molecular axis oriented parallel to the hBN's zigzag direction (figure 6). A flat-lying configuration of molecules is frequently observed for 5A on a variety of substrates [190–192]. High-resolution AFM measurements revealed a unit cell of the first layer which is ~3.3% larger along the a-axis than that of the second layer. TEM and SAED of 5A/hBN films revealed lattice constants for the second layer (a  =  5.98  ±  0.09 Å, b  =  7.61  ±  0.13 Å) (see figures 5(e) and (f)) and angle between the unit cell vectors (α  =  88.3  ±  1.2°) that correspond well to the 5A thin film phase. Analysis of the SAED patterns from several different few-layer 5A crystals showed mostly an angle of 16° between the 5A(0 1 0) and the hBN(1 0 0) crystal orientation indicating a quasi-epitaxial relationship between 5A and hBN. The quasi-epitaxy was explained to arise from the weak vdW interaction with the substrate and the incommensurability between 5A and hBN. However, the presence of other angles indicated that the epitaxy is not perfect. DFT calculations of the relaxed first and second layer structures agreed within 1.5% with the experimentally found structures. The 5A packing in each layer is determined by the competition between interlayer and intralayer interactions. The flat lying configuration in the interface layer can be explained by the high binding energy of 2.35 eV for a flat lying molecule with its long molecular axis oriented in the hBN [1 1  −  2] direction. In contrast, the interlayer interaction energy in the second layer is only 0.3 eV which leads to the upright thin film configuration. The molecules in the first layer tilt their long axes towards the interface layer in order to increase the π-π interaction. At the same time, the short molecular axes align more parallel with the a-axis of the unit cell resulting in repulsion of the molecules which in turn leads to an expansion along the a-axis as experimentally observed.

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Combined polarization dependent adsorption microscopy and high-resolution AFM images revealed the existence of first layer single domains as large as ~60 µm only limited by the hBN flake size. The good crystalline quality was further confirmed by photoluminescence measurements which resulted in free exciton binding energies 0.3 eV smaller than those found in 5A thin films [193] and an estimated exciton radius of ~5.7 nm.

5A crystals of different thickness were electrically tested using a back gated OFET geometry. It turned out that the interface layer was insulating consistently with the lack of π-π intralayer stacking in this layer. In contrast, devices containing the interface layer and the first layer had excellent OFET characteristics with field effect mobility of µ ~ 1.6 cm² V⁻¹ s⁻¹ at room temperature. In this case, 2D hopping transport was found as charge transport mechanism [14, 194, 195]. Interestingly, second layer devices showed bandlike charge transport with room temperature mobility of µ ~ 3 cm² V⁻¹ s⁻¹ and µ ~ 5.2 cm² V⁻¹ s⁻¹ at around 100 K, which by far outperformed typical polycrystalline 5A OFETs at similar temperatures [196]. The difference in transport mechanisms of the first and second layer was rationalized by calculating the molecular orbitals of the intermolecular bonding states. Indeed, continuous overlap was found for the second layer whereas the overlapping orbitals only span five molecules in the first layer reducing the travel length of a charge carrier to only ~1 nm. 5A OFETs with three or more layers had similar maximum room temperature mobility as two-layer devices indicating that the interface induced modulation of the molecular packing and transport are insignificant beyond the second layer. The significance of hBN as gate dielectric becomes apparent when comparing the device performance of single layer 5A FETs with classical SiO₂ dielectric. First single and double layer 5A FETs with SiO₂ gate dielectrics were realized by physical vapor deposition with nitrogen as carrier gas [197]. Using classical top source-drain electrode contacts fabricated by vapor deposition of 30 nm of gold, mobility values of only µ ~ 0.01 cm² V⁻¹ s⁻¹ were achieved for single layer 5A FETs and µ ~ 0.15 cm² V⁻¹ s⁻¹ for double layer devices. A mobility improvement for a single layer device to µ ~ 0.015 cm² V⁻¹ s⁻¹ was realized by employing bottom gold contacts modified with self-assembled monolayers of perfluoro octanethiol to optimize contact resistance [198]. A tremendous mobility improvement for single layer 5A to µ ~ 0.31 cm² V⁻¹ s⁻¹ was achieved when a polystyrene layer was spin coated on the SiO₂/Si substrate prior to 5A deposition and when a special soft imprinting method for placing the gold source and drain contacts was applied. Even though significant improvements have been demonstrated, the mobility values are low compared to the 5A FET with hBN dielectric. The reason is that the single layer 5A films on conventional and modified SiO₂ are still polycrystalline with a significant density of grain boundaries that dominate the charge transport and permit hopping transport only.

Even though the vdW interaction between hBN and organic compounds is weak, it is evidently still sufficient to promote a quasi-epitaxial alignment of the organic crystal with the hBN substrate. This has been also exploited to grow self-assembled para-hexaphenyl (6P) needle networks with discrete needle orientations with respect to the substrate [12]. As substrates, hBN exfoliated from powder and subsequently transferred on SiO₂/Si (80 nm oxide thickness) were employed. Typical lateral hBN flake sizes in the range of 5–30 µm were obtained with varying heights between 0.6 nm to several tens of nm. 6P was deposited via hot-wall epitaxy [147] under high vacuum conditions (p  ~ 2  ×  10⁻⁶ mbar) at substrate temperatures (T_D) between 300 K and 400 K at a growth rate of ~9  ×  10³ molecules µm⁻² s⁻¹. After a typical growth times of ~5 min, the 6P coverage was ~0.6 monolayers (one monolayer (ML) refers to the packing density in the 6P (0 0 1) plane: ~4.4  ×  10¹⁴ molecules cm⁻² [199, 200]). The resulting film morphologies were characterized using intermittent contact mode AFM under ambient conditions. The morphology was dominated by needle-like structures as can be seen in the AFM image in figure 7(a). Their average lengths ranged—depending on the growth temperature—from several 100 nm (T_D  =  298 K) to several µm (T_D  =  393 K). Occasionally, at T_D  >  380 K individual needles of up to 30 µm length were found, only limited by the hBN flake's size. 6P needles have been frequently observed and were identified to be crystallites composed of molecules with their long axes oriented parallel to the substrate surface [64, 201]. Similar to the behavior of 6P on Gr [11, 32, 97, 172], the needles' long axes on hBN followed six discrete growth directions which grouped in three pairs of two with around ~10° in between, and ~60° in between the three pairs (see fast Fourier transformation image in figure 7(b)). DFT calculations revealed that individual 6P molecules adsorb with their long axes parallel to hBN's armchair direction with the phenyl rings centered on the nitrogen atoms (see figure 6). With increasing coverage, the individual molecules—aligned with the armchair direction—attach along their long molecular axes and form stable nuclei. From there, growth proceeds preferentially in the direction perpendicular to the long molecular axes (in this case, the hBN zigzag direction). However, the evolving structure is 3D, and the vdW molecule-substrate interaction is weak. Thus, intermolecular interactions dominate and drive the molecular system towards its bulk equilibrium structure. For 6P under the applied temperature regime this is the 6P β-phase herring-bone structure (Baker structure) [90, 96]. In order to minimize the system energy, a contact plane of the 6P bulk is chosen that causes the smallest possible strain between the interfacing molecules in their preferred adsorption sites and the 6P bulk. In this case, the 6P contact plane fulfilling the requirements best was the 6P(−6 2 9) plane [96] which was, for example, also found for 6P on Cu(1 1 0) [202]. As shown in figure 7(c), starting from the (−6 2 9) plane with a single 6P in its preferred adsorption geometry on hBN we obtain a needle growth direction which is ~6° tilted with respect to the hBN zigzag direction. The mirror symmetry of the system implies a second equivalent growth direction which is rotated by 2  ×  6°  =  12°. Thus, a  ±6° splitting from the zigzag direction is expected in good agreement with the observations.

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Further, with increasing substrate temperature more and more needles followed the discrete growth directions indicating that such an alignment is energetically preferred. A comparison of the 6P growth morphologies between different hBN thicknesses revealed that hBN flakes thinner than three layers did not support growth of long needles [12]. This was attributed on the one hand to surface roughness which was translated from the interface with the SiO₂/Si support through the very thin layers and on the other hand to the low dielectric constant of hBN that caused insufficient screening of dipole fields arising from trapped water at the interface. Thus, a flake thickness of at least ~1.5 nm was needed to cancel out interference from the support.

For such ordered 6P needles on hBN, the influence of the vdW interface on the nanoscale friction between organic crystallites and Gr and ultra-thin hBN has been studied by means of AFM manipulation and was compared to 6P needles on Gr [203]. It was shown that crystalline 6P nano-needle fragments preferentially moved along their preferred growth directions rather than directly following the primary pushing direction given by the AFM tip movement. Further, application of a torque by the AFM resulted in rotation of the crystallites. When crossing the preferential sliding/growth directions of the needles a significant increase of the torsional signal of the AFM cantilever was obtained. Supported by molecular dynamics simulations, it turned out that the clear frictional anisotropy and preferred sliding directions are determined by the complex epitaxial relation between organic crystallite and 2D substrate. It takes higher forces to reorient the crystal from its original directions because it has to leave its commensurate state. Once the crystal is out of the commensurate state, it can be easily moved until another commensurate state is hit and the crystallite locks again. Even though the interaction between the vdW substrate and the organic molecules is comparably weak, it is still sufficient to have significant influence on the frictional properties of the system [203].

Networks of needle like crystallites with discrete growth directions were also found for the heptacene derivate dihydrotetraazaheptacene (DHTA7) HWE grown on exfoliated hBN/SiO₂/Si [204]. The substrate temperature was kept in the range of 355 K—395 K at growth rates of ~0.35 ML h⁻¹. (1 ML refers to a compact single layer of upright standing molecules.) At a final coverage of ~0.6 ML, crystallites up to 300 µm long with uniform heights between 5–10 nm and widths between 50 nm and 100 nm were observed. As substrates 80–160 nm thick flakes were chosen in order to enhance the optical contrast between DHTA7 crystallites and the substrate facilitating optical microscopy. Average lateral flake sizes were around ~200 µm.

Even though the DHTA7 bulk structure deviates from the herringbone arrangement frequently observed for rod-like conjugated molecules very, similar growth patterns to 6P have been found as can be seen in the AFM image in figure 8(b). Again needles with six preferential growth directions (three 60° rotated chiral pairs, each split by ~18°) were clearly identified (see FFT image in figure 8(c)). According to DFT calculations individual DHTA7 molecules preferentially adsorb with the molecular long axis parallel to hBN's zigzag direction (figure 6). Considering that two adjacent molecules confined in a plane interact via dipolar interactions, steric repulsion and (N-H···N) mediated hydrogen bonding the optimum adsorption geometry is a head to tail arrangement of the molecules with their long axes parallel to each other but shifted by ~1 Å. The slight relative shift of the molecules defines the angle between the long needle axis and the long molecular axes. The fastest growth of the crystallites is along the direction of the hydrogen bonds resulting in elongated crystallites, which should theoretically grow  ±9.2° away from the zigzag direction which is consistent with the observed ~18° split. Based on the latter argumentation a tentative model of the molecular arrangement at the DHTA7/hBN interface is presented in figure 8(a). The in-plane hydrogen bonds enforce a rather planar molecular arrangement with poor in-plane orbital overlap but increased π-π stacking with subsequent layers as also observed for related systems [98]. According to DFT, a unit cell contains four molecules whereas the individual molecules are tilted by ~8.5° with respect to the DHTA7 (0 0 1) plane (figure 6(b)) which is considered to be also the present contact plane with hBN. In this case, the angle between long molecular axes and the long needle axis should be ~79°.

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Optical microscopy with linearly polarized light strongly indicated that the needles were single domains. Taking into account that light absorption is strongest when the polarization direction of the incident light matches the long molecular axis, a relative orientation angle between the long needle axes and the long molecular axes of (79  ±  3)° could be deduced, in agreement with the theoretical predictions.

Since the orbital overlap in needle growth direction is very weak, also the electrical conductivity along the needle is extremely low. Therefore, charge transport in the needles was studied using electrostatic force microscopy (EFM) [205]. In the dark, charges deposited via the biased AFM tip stayed localized around the charging point whereas under illumination charges spread out through the needle network revealing photoconductivity. Since the light absorption was depending on the light polarization, directional charge movement could be stimulated. Only those needles in the network became photoconductive that contained molecules whose long molecular axes matched the polarization direction.

In order to quantify the conductivity along the needles, time dependent EFM measurements on single needle field effect devices were carried out. The field effect devices were realized by transferring graphite flakes as contact electrodes onto the hBN flakes prior to DHTA7 deposition. That way, resistivity values in the order of ~1  ×  10⁹ Ωm in the dark and ~3.5  ×  10⁷ Ωm under illumination were found. The low conductivity along needle direction was in accordance with the growth model and the poor orbital overlap between adjacent molecules.

The adsorption of the linear, long chain alkane hexacontane (C₆₀H₁₂₂) on exfoliated hBN was investigated utilizing high resolution AFM [184]. The adsorption of low cost, non-toxic alkanes is of interest in the fields of catalysis [206], surface passivation [207], and lubrication [208], but they are also interesting systems to study supramolecular assembly [209]. Hexacontane films were prepared either by vapor deposition in high vacuum or from solution. From the monolayer morphology, a flat lying molecular orientation was deduced. The monolayer morphology revealed from AFM measurement shown in figure 9 consisted of parallel rows with a period of (7.9  ±  0.2) nm (see figure 9(b)) and the measured periodic corrugation along the rows yielded (0.44  ±  0.02) nm very close to the molecular spacing found in the bulk crystal [80, 210]. Further, it turned out that the individual molecules align with their long axes parallel to hBN's zigzag direction (see figures 9(c) and (d)), which was also supported by molecular mechanics simulations and is consistent with x-ray investigations [211]. A structural model is provided in figure 10(a). The adsorption energy for zigzag alignment is favorable by 70 meV per molecule. A rough alignment of alkyl chains with the hBN's zigzag direction was also found for C₈-BTBT (see figure 4(a)) and might be a general trend for molecules with long hydrocarbon chains. In higher layers the row configuration is continued, however, with decreasing disorder indicated by meandering of the rows. Nevertheless, the growth was mainly directed by the hBN substrate.

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Adsorption and ordering on transferred CVD hBN.

Maybe, the simplest molecule to use is C₆₀, which behaves isotropic due to its spherical symmetry. In fact, C₆₀ based OFETs with mobility values up to µ ~ 2.9 cm² V⁻¹ s⁻¹ and on/off ratios of 10⁷ have been realized on CVD grown single layer hBN transferred to a SiO₂/Si substrate [49]. In this case, hBN was used to decouple the C₆₀ film from the SiO₂. The insertion of the hBN interface layer increased the mobility by a factor of ~40 compared to C₆₀ films directly grown on SiO₂. The superior charge transport properties were attributed to the epitaxial 2D growth of C₆₀ on hBN. C₆₀ films with a nominal thickness of 20 nm were thermally evaporated at a substrate temperature of 383 K and a rate of 0.02 nm s⁻¹. AFM measurements revealed 100–200 nm large terraced grains with a typical step height of ~0.8 nm, which fits well to the height of a single C₆₀ layer. Significantly smaller grains without noticeable step structure were found on a reference film grown under identical conditions on a bare SiO₂ substrate. Further, the film roughness was much higher on SiO₂ than on hBN. From TEM and SAED measurement of C₆₀ films deposited onto TEM grids with hBN as shown in figure 11, confirmed an epitaxial relationship between hBN and the C₆₀ film. Additional grazing incidence x-ray diffraction investigations clearly revealed that the (1 1 1) planes of the face centered cubic C₆₀ crystals were aligned parallel to the hBN surface and confirmed the 2D closed structure of the C₆₀ films which indicates a 2D, layer-by-layer growth facilitated by the vdW interaction between C₆₀ and CVD hBN. Further, it was found that the C₆₀ crystal edges typically align either 10° or 26° off the hBN zigzag direction. Calculations of the most stable orientation performed via DFT calculations including vdW dispersion forces yielded consistent results with the experimental observation predicting the most stable orientations with the C₆₀ crystal edges aligned 9°–12° and 21°–25° with respect to the hBN zigzag directions (see structural model in figures 11(a) and (b)).

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2D molecular self-assembling on exfoliated hBN.

The formation of self-assembled monolayers of diacetylene has been demonstrated on atomically flat exfoliated hBN sheets [81, 212]. In this case, 10,12-nonacosadiynoic acid diacetylene molecules were brought into solution with xylene and spin coated on the hBN/SiO₂/Si substrates. AFM characterization of the SAMs on hBN revealed long parallel stripes corresponding to self-assembled molecular rows of flat lying molecules with a stripe period of (7.6  ±  0.2) nm. For the molecular arrangement, a model was suggested where molecules group in pairs such that the carboxy end groups of one chain face those of the neighboring chain [81]. Further, it was suspected that the alkyl chains of the molecules align with one of the main crystal axes of the hBN lattice [81, 211]. Based on the idea that the alkyl chains tend to align with the zigzag direction as confirmed for hexacontane [184], we suggest a tentative structural model which is provided in figure 10(b). Unlike substrates as HOPG or MoS₂ where domain sizes were only in the few 100 nm range, on hBN single domains—a few µm in size—were found. This was attributed to a reduced nucleation density due to the exceptional large area flatness of hBN. It was noted that for the self-assembling of the diacetylene molecules single atomic steps of the substrate might be sufficient to interrupt the process. In addition, on these substrates also on-surface photopolymerization of the SAM to polydiacetylene was successfully performed. It was also found that the polymerization rate was two orders of magnitude higher than for graphite, which was attributed to the large difference between the hBN band gap (~5.9 eV [37]) and the excitation energy of the diacetylene (~3.1 eV).

Nanoporous networks formed from perylene tetra-carboxylic di-imide (PTCDI) and 1,3,5-triazine-2,4,6-triamine (melamine) have been prepared on hBN/SiO₂/Si substrates [108]. The hBN flakes were obtained by mechanical exfoliation from mm-size boron nitride crystals. The hydrogen bond stabilized networks were obtained via immersion of the substrates in a PTCDI and melamine containing dimethylformamide solution at 373 K. High-resolution AFM measurements revealed the formation of monolayers of honeycomb-like network areas with lateral extensions of ~0.3 µm. The lattice constant was (3.5  ±  0.1) nm and very similar to what was found for PTCDI-melamine networks grown on metal substrates [213, 214]. An ordering motif of PTCDI/melamine is provided in figure 12. The same authors demonstrated that supramolecular heterostructures with different structural and optical properties can be formed on hBN by sequential stacking of 2D hydrogen bonded arrays [109]. In this case, the stacked heterostructure components were hexagonal networks formed by melamine (M) and cyanuric acid (CA), honeycomb arrays of trimesic acid (TMA) and linear arrays of terephthalic acid (TPA). As first layer at the interface to hBN the cyanuric acid/melamine (CAM) network was formed. A model of the molecular arrangement is shown in figure 12. AFM measurements revealed µm sized (depending on immersion time) single layer hexagonal networks with a lattice constant of a_CAM  =  (0.98  ±  0.02) nm which agreed within the error margin with the bulk parameters of crystalline melamine-cyanurate (a  =  0.96 nm [215]). Also a moiré pattern with a period of (6.5  ±  0.4) nm rotated by ~21° with respect to the CAM lattice vectors was found. High resolution contact mode AFM images of a CAM network and the uncovered adjacent hBN revealed a (3  ±  1)° missorientation between the symmetry axes of the CAM network and hBN, which gave rise to the moiré structure. A well-defined epitaxial relation was also found for the TMA and TPA arrays grown on the CAM network demonstrating that hydrogen bond stabilized 2D structures with parallel interfaces can be grown on hBN to form complex vdW heterostructures.

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Multilayers of covalent organic frameworks (COFs) on exfoliated hBN transferred to SiO₂/Si or HfO₂/Si were selectively grown by a solvothermal process using TAPP and TAP precursors [107]. The work was motivated by the fact that layered 2D COFs represent planar π-conjugated systems with well-ordered stacking, which can result in exceptionally high charge carrier mobilities [216, 217]. Further, theoretical work predicts a minimum perturbation of the electronic band of the active layer on the hBN support [218] making intrinsic properties accessible. The presence of 5–8 monolayers of the COF was deduced indirectly from Raman and fluorescence measurements and has been complemented by XPS, UV–VIS-NIR absorption spectrometry, and AFM. A comparison of the Raman spectra of the COF and the corresponding precursors revealed that the aldehyde stretching vibration bands related to TPA at 1691 cm⁻¹ and 1702 cm⁻¹ were missing in the film. Further, two bands at 1570 cm⁻¹ and 1625 cm⁻¹ appeared that were attributed to the C  =  N bond stretching vibration between the phenyl groups of TPA and TAPP. These trends were considered as indication for an imine-connected structure. The presence of a C  =  N bond was additionally confirmed by XPS. A model of the formed network can be found in figure 12. Statistical evaluation of AFM topography films before and after COF deposition yielded a COF film thickness in the range between ~2.3 and ~4.6 nm on the ~40 nm thick (~120 layers) hBN flakes.

The fluorescence properties of 2D supramolecular networks of 5,10,15,20-tetrakis(4-carboxylphenyl)porphyrine (TCPP) solution deposited onto exfoliated hBN/SiO₂/Si have been studied [68]. The adsorbate-substrate vdW interaction caused a clear red-shift of the TCPP fluorescence. The hBN flakes were obtained by exfoliation from hBN single crystals and had typical lateral dimensions of 20–100 µm and exhibited thicknesses between 20 and 100 nm which corresponds to 60–300 hBN layers. After immersing the hBN/SiO₂ substrate for 10 s into a 13 µM ethanol TCPP solution, AFM images revealed the existence of about ~0.26 nm high islands with typical lateral extensions of 50–100 nm. The observed island height suggested that the TCPP molecules adsorbed flat lying with their porphyrine macrocycle oriented parallel to the surface. High-resolution AFM images revealed two distinct TCPP phases with square and hexagonal symmetry, respectively. The square lattice had a lattice constant of a ~ 2.22 nm, whereas the hexagonal lattice had one of a ~ 4.37 nm [68]. However, the hexagonal structure turned out to be unstable and disappeared within ~1 d. A model of the stable square lattice arrangement can be found in figure 12. The observed structures were modeled by hydrogen bonded junctions of two or three carboxylic acids for the square and hexagonal phases, respectively. Corresponding DFT calculations strongly supported this model and estimated the binding energies per molecule to be 2.30 eV for the square lattice and 0.68 eV for the hexagonal lattice in agreement with the observed stability of the phases [68]. However, more refined calculations using DFT including dispersion corrections [178] revealed a deformation of the porphyrine macrocycle upon adsorption on hBN resulting in a bowed geometry of the molecules. This was attributed to the aryl side groups, which cannot adopt a coplanar alignment with the porphyrine macrocycle and lift up the plane of the macrocycle resulting in a small gap between molecule and substrate. Driven by vdW forces, the macrocycle can flex and bow down towards the gap adopting the bowed geometry. A direct consequence of the non-planar conformation of the porphyrine macrocycle is a red-shift in the Soret and Q-band fluorescence of the molecule [219].

Historically, the first adsorption experiments on an hBN/Rh(1 1 1) nanomesh were performed using C₆₀ fullerenes [52]. However, in this case no systematic study from low to high coverages were conducted. Therefore, no explicit prove was provided that the C₆₀ molecules are initially trapped in the pores. Nevertheless, well ordered structures—closely following the nanomesh topography- were reported. According to STM measurements, the nanomesh wires were nicely decorated by C₆₀ whereas the very pore center stayed often—but not always—unoccupied. That might be interpreted as preferred off center adsorption as well. The C₆₀ spacing along the wires was around ~1 nm. 1 nm nearest neighbor distance was also found on C₆₀ monolayers on hBN/Ni(1 1 1) [67]. They were obtained by thermal desorption of multilayers and formed a commensurate hexagonally closed packed (4  ×  4) structure. This is consistent with the molecular spacing in the fcc C₆₀ (1 1 1) plane, which was also found for C₆₀ adsorbed on CVD grown hBN on SiO₂/Si [49]. In this respect, metal supported hBN seems to be similar to transferred hBN. However, geometrical and electrical corrugation might vary quite significantly for different metal supports, which also might lead to different adsorption geometries.

C₆₀ is a molecule noticeably smaller than the nanomesh pore but isotropic due to its spherical shape. An interesting example to compare is the 5A molecule, which is also significantly smaller than the nanomesh-pores but anisotropic due to its rod-like shape. On hBN/Rh(1 1 1), 5A initially adsorbs flat lying and is preferentially located in the pores [72]. The electric field at the pore-wire transition attracts the 5A molecules causing them to arrange along the pore rim. However, the 5As adsorb preferentially with their long axes facing an intersection of the wires and never adsorb parallel to the wires. That way, six molecules can arrange along the pore perimeter forming a six-fold geometry, which is not found for 5A on exfoliated hBN [14, 65]. However, defects or adsorption of multiple 5As in the same valley can cause other adsorption geometries [72]. Increasing the 5A coverage beyond 0.2–0.3 ML yields a transition from flat lying molecules to an upright configuration, as has been indicated by NEXAFS [165]. Accompanied with increasing film thickness also changes in the C1s photoemission peak have been observed, which were attributed to originate entirely from the gradual rearrangement from flat-lying to upright 5A configuration. On exfoliated hBN flat lying molecules formed a compact interface layer on which the succeeding layer grew in an upright fashion [14]. Further investigation with micro-low energy electron diffraction and STM on a ML thick island grown at room temperature confirmed an almost upright standing arrangement in the typical herringbone motif as frequently found for rod-like conjugated organic molecules [32, 94, 96, 222]. However, the influence of the hBN nanomesh corrugation on the resulting 5A structure appeared to be strong. Even, though the typical herringbone structure was adopted, the lattice parameters and unit cell angles deviated noticeably from the 5A bulk [223] arrangement. Since organic structures—primarily bound by vdW interactions—are pliable, it appears likely that the geometric and electrostatic corrugation of the hBN nanomesh can sufficiently modify the film in order to achieve best match with the substrate. A flat-lying to upright transition is very typically for rod-like aromatic molecules in case of weakly interacting amorphous, contaminated, or strongly corrugated surfaces. Going upright, forming herring-bone packed islands in, or close to the herringbone bulk structure is a general trend and is closely related to the possibility to form an ordered, compact wetting layer [10, 32, 97, 222, 224, 225]. Apparently, the corrugation of the nanomesh is sufficient to enforce the upright molecular configuration form the first layer on without forming a flat-lying interface layer.

5A and C₆₀ are examples of molecules which are smaller than the nanomesh pores allowing for multiple adsorption within a single pore. In contrast, Nc is almost as large as a pore (~2 nm), consequently only single occupation is possible. On hBN/Rh(1 1 1), a highly selective adsorption in the pores was observed for vapor deposited Nc at room temperature [54]. The distance between the pores was large enough to suppress molecule-molecule interaction between adjacent occupied pores. Therefore, ordering was fully dominated by the molecule-substrate interaction. This way, a well ordered array of Nc molecules was realized with the periodicity of the nanomesh (~3.22 nm). In contrast, on atomically smooth hBN a significant influence of intermolecular interaction on the ordering can be expected.