J. Phys. Chem. C 2009, 113, 20927–20933
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Kinetic Phase Selection of Rubrene Heteroepitaxial Domains Marcello Campione,* Massimo Moret, Luisa Raimondo, and Adele Sassella Department of Materials Science & CNISM, UniVersity of Milano Bicocca, Via R. Cozzi 53, I-20125 Milan, Italy ReceiVed: June 18, 2009; ReVised Manuscript ReceiVed: August 4, 2009
The high carrier mobility measured in rubrene single crystals makes its growth in the form of crystalline thin films a challenge of great interest in view of their use as semiconductor layers in integrated organic devices. Here, heteroepitaxial thin films of rubrene are grown by molecular beam epitaxy under different conditions, showing that the film structure can be tuned from disordered to crystalline by the sole control of the deposition rate. Highly oriented crystalline films are demonstrated to grow undergoing a transition characterized by a strong mass transport among different domains, persisting even during the postgrowth stages. This mechanism is rationalized within a thermodynamic model which gives quantitative predictions starting from empirical force field calculations and atomic force microscopy topographical data. By an appropriate setting of the growth parameters, an epitaxial rubrene monolayer is obtained at room temperature. These results, rationalizing the concrete possibility of growing crystalline and oriented rubrene thin films, open new perspectives to organic-based thin film technology and devices. I. Introduction Rubrene (RUB, C42H28, 5,6,11,12-tetraphenyltetracene) is rapidly becoming one of the most attractive organic semiconductors thanks to the extremely promising charge carrier mobility of about 20 cm2/(V s) measured in single crystals,1 the highest ever measured in organic materials. Due to the consequent great interest in developing a thin film technology based on RUB as a semiconductor, researchers started facing the challenge of growing crystalline thin films with highly controlled properties. Despite the propensity of this material to give well-formed single crystals when grown by vapor transport techniques, the majority of the attempts to grow crystalline films of RUB have been unsuccessful: the material is observed to solidify in an amorphous form, giving rise to ill-connected domains with poorly controlled morphologies. Polycrystalline films were obtained only on properly deposited buffer layers2 or by performing postgrowth annealing treatments3 aimed at inducing the amorphous-to-crystalline transition; in a recent paper, crystalline homoepitaxial RUB films were successfully grown from the vapor phase by vacuum sublimation, establishing the adequacy of molecular beam methods for obtaining crystalline thin films of RUB.4 Following the organic epitaxy approach,5 we demonstrated the successful growth of an epitaxial nanostructure of RUB on the (001) surface of tetracene (TEN) single crystals, kept at 100 °C, achieving a single crystalline order of the overlayer.6 The experimental epitaxial relation corresponds to an heterojunction interface where the direction and spacing of the main corrugations of substrate and overlayer surfaces coincide. Furthermore, thanks to the p1 plane symmetry of the substrate surface, a single azimuthal orientation of the RUB domains is nucleated, giving rise to a film phase with single crystalline order. In this paper, we analyze in detail the growth dynamics of this system by using a model developed on the basis of empirical * Corresponding author. Address: Department of Geological Sciences and Geotechnologies, University of Milano Bicocca, Piazza della Scienza 4, I-20126 Milan, Italy. E-mail:
[email protected]. Phone: +39 02 64485012. Fax: +39 02 64485400.
force field calculations and experimental morphological parameters of the grown films. By carrying out growth experiments of the RUB/TEN heterojunction by organic molecular beam epitaxy (OMBE) under different deposition rates at room temperature, we realized that the structure of the RUB overlayer can be tuned from disordered, at low rates, to single crystalline, at high rates. In particular, the possibility to select the crystalline phase is ensured by a substantial mass transport occurring from disordered to crystalline domains, both during growth and during the postgrowth stages. The proper tuning of the deposition parameters allowed us to obtain a uniform crystalline monolayer of RUB at room temperature and without any postgrowth treatment. II. Experimental and Computational Methods TEN single crystals were grown with the physical vapor transport method7 by using nitrogen as a gas carrier and heating the source at 170 °C; they are thin flakes exposing a wide, clean, molecularly flat {001} surface, a few mm2 in size, which does not require cleavage before use. The TEN crystal structure belongs to the triclinic system, with two molecules in the unit cell and parameters a ) 6.06 Å, b ) 7.84 Å, c ) 13.01 Å, R ) 77.13°, β ) 72.12°, and γ ) 85.79°.8 Commercial RUB (Acros Organics, 98%) was purified through several steps of vacuum crystallization and then used as the source in a Knudsentype effusion cell. All of the thin films were grown at room temperature by OMBE under high vacuum using a quartz microbalance to monitor the deposition rate,9 which was changed for the different samples by properly controlling the source temperature. RUB thin films crystallize on TEN(001) as (100)oriented domains in the orthorhombic form,6 having four molecules in the unit cell with parameters a ) 26.86 Å, b ) 7.19 Å, and c ) 14.43 Å.10 The morphology of the grown films was analyzed by atomic force microscopy (AFM), collecting images in tapping mode with silicon cantilevers (force constant, 40 N/m; resonance frequency, 250 kHz; tip apex radius of curvature, 10 nm) under a dry nitrogen atmosphere (evaporated from the liquid) with a MultiMode Nanoscope IIIa (Veeco).
10.1021/jp905752r CCC: $40.75 2009 American Chemical Society Published on Web 08/27/2009
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Calculations based on atom-atom empirical potentials were performed to study the dependence of the interaction potential energy on the in-plane azimuthal orientation of a RUB(100) slab on top of a bulk terminated TEN(001) surface. A modified version of the AutoDock3 molecular docking package11 was used in combination with the UNI empirical force field12 for organic systems. The modifications are aimed at accepting very large systems. The maximum number of grid points in the interaction potential maps calculated with Autogrid3 has been increased to 3513 in order to have a fine sampling grid. The maximum number of atoms comprising the substrate and overlayer has been increased to 49152 and 4608 atoms, respectively. A simulation box comprising 3513 grid points (sampling grid 0.217 Å) was used to map the interaction potential between the TEN substrate and the epitaxial RUB crystallites. The TEN(001) substrate surface was modeled with a slab of 19 × 15 × 2 unit cells along the a, b, and c axis, respectively, giving rise to a total of 1240 molecules. The RUB crystallite was modeled by one RUB(100) layer comprising 16 molecules; a total of 2261 docking runs were performed, giving satisfactory final sampling statistics of the in-plane orientations of the RUB crystallites. Approximate energy minima values found with AutoDock3 were improved by minimization with the program Orient 4.6.11 using the same interaction potentials.13 III. Results and Discussion III.a. Interface Energies of the RUB(100)/TEN(001) Heterostructure. The detailed study in ref 6 shows that AFM gives all the details on the microscopic surface and interface properties of the heterostructure RUB/TEN. In particular, (i) the RUB molecules arrange on the TEN(001) surface in the single crystal structure, with RUB(100) as the contact plane, and (ii) the epitaxial relationship, well assessed both experimentally and by crystallo-chemical arguments,6 makes the directions RUB[021] and TEN[1-10] align (Figure 1a). On the basis of the AFM data analysis, the interactions between aromatic hydrogen atoms emerging at the TEN(001) surface and RUB phenyl rings attached to the TEN core are identified as the main actors on deciding which azimuthal orientations are best suited to give a profitable interaction energy. To clarify this, adding new insights to the description of the interface properties of the heterostructure RUB/TEN, potential energy calculations were performed. A RUB(100) crystallite comprising 16 molecules, randomly placed above an unreconstructed TEN(001) surface,14 was free to move and interact with it and the potential minima searched for: the results are summarized in Figure 1b. The resulting adhesion energy strongly depends on the azimuthal orientation of the RUB crystal with respect to the TEN substrate. The plot of potential energy as a function of azimuthal rotation shows few marked peaks associated with local energy minima. In particular, the best energy minimum of -342 meV per molecule in the cluster appears at an azimuthal angle of RUB[010] vs TEN[100] of 78.4°. Other secondary energy minima are present at 170.2° (-329 meV), 3.5° (-312 meV), 146.6° (-306 meV), and 31.7° (-300 meV). Experimental observation by means of high resolution AFM revealed only the azimuthal orientation at ∼78°, while all other orientations were not observed at all.6 As to the peak at ∼170°, it is reasonable to expect it from modeling, once a peak at ∼78° is found, due to the almost perfect 2:1 ratio between the c and b lattice parameters of RUB,10 which makes the two orientations similar in terms of close contacts and hence potential energy. However, intermolecular interactions at work during the nucle-
Figure 1. (a) Molecular structures of RUB and TEN and model for the molecular arrangement at the interface of the RUB (red layer)/ TEN (yellow layer) heterostructure. (b) Empirical force field calculations showing the potential energy for many different in-plane alignments of (100)-oriented RUB epitaxial islands (positive rotations are counterclockwise). The asterisks below the minima indicate the azimuth and energy obtained after minimization. Energy has been normalized to a single RUB molecule contacting the substrate.
ation of RUB islands atop the TEN(001) surface are able to suppress the appearance of this orientation due to the difference in adhesion energy, and the relative statistical weight, for a critical sized cluster composed of a few molecules. The energy minimum at 78.4° provides the specific energy of adhesion of (100)-oriented RUB domains on TEN(001) which results 6.6 meV/Å2. It is worth comparing this energy with the surface energy of the crystal faces involved with the heterostructure: as indicated by a preliminary periodic bond chain analysis15 of the crystal structure of RUB carried out with the same UNI potential, {200} faces are the most stable ones with 4.61 meV/ Å2, followed by {020}, {002}, {202}, and {111} with values of 5.57, 6.14, 6.46, and 7.14 meV/Å2, respectively. With the same procedure, the energy of the TEN(001) deposition surface has been estimated to be 5.0 meV/Å2. All of these energies fall well within the range 3-10 meV/Å2, enclosing the available theoretical estimates for the surface energies of crystals of other organic semiconductors having a molecular packing motif similar to those of RUB and TEN.16 The present modeling is based on both rigid and static molecules and aggregates/crystals; therefore, it does not take into account likely variations of small cluster structures compared to the final film phases. The same kind of simulation was in full agreement with experimental observations in the case of epitaxial growth of para-sexi-phenyl on (010)KAP.17 Then, compared to less computing intensive lattice geometry calculations18 (almost always telling us that organic surface lattices are incommensurate), these results are more informative and reliable in that they properly take into account the most
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Figure 2. 15 × 15 µm2 AFM images showing the as-grown surface morphology evolution with deposition rate of 1 nm thick RUB thin films deposited on the TEN(001) surface at room temperature. The deposition rates are (a) 0.025, (b) 0.10, (c) 0.20, (d) 0.26, and (e) 0.38 nm/min. Below each image, a cross-sectional profile taken along a horizontal image line is reported.
relevant driving force, i.e., local intermolecular interactions determining the outcome of organic-organic heteroepitaxy. Indeed, it would be difficult to explain on the basis of lattice geometry calculations a rotation of 80° of the RUB lattice with respect to TEN. On the contrary, best matching and hence best binding energy is found with our modeling scheme. III.b. Growth Dynamics. Figure 2 shows five images collected by AFM over a 15 × 15 µm2 region of three submonolayer thick RUB films (1 nm nominal thickness) grown on TEN crystals at different deposition rates, from a few hundredths of nm/min up to about 0.4 nm/min. The figure immediately suggests a strong dependence of film morphology on the deposition rate I. Indeed, for I ) 0.025 nm/min, RUB thin films are composed of cap-shaped nanodots having a contact-base diameter of 400 ( 100 nm, a height of 30 ( 8 nm, and a number density of 0.48 ( 0.04 µm-2 (Figure 2a). One can also observe the cross section of the nanodots and the typical roughness of the TEN(001) surface, characterized by the presence of monomolecular holes (depth: 1.21 nm).19 The spherical shape of the nanodots is indicative of an isotropic, disordered phase. Two isotropic phases of RUB are known, namely, amorphous and spherulitic.2a,20 However, due to the metastability of the dots (see below), their in-depth structural characterization is severely hindered. However, the dots formed on TEN(001) show close similarities with the amorphous dots grown and characterized by X-ray diffraction on sapphire substrates;2a on the contrary, the spherulitic phase has been identified as a late-time evolution of the amorphous phase, displaying domains up to some tens of micrometers wide with
fiber morphology.20 From these observations, we assume the dots formed on TEN to have an amorphous structure. At I ) 0.10 nm/min (Figure 2b), a fascinating phenomenon is observed: branched flat islands extending over areas as large as 50 µm2 start nucleating and coexist with the nanodots, which display here a smaller diameter of 320 ( 40 nm, and a reduced height of 24 ( 5 nm, while their number density increases to 3.0 ( 0.2 µm-2; the crystalline island height is 1.3 ( 0.2 nm, corresponding to one monomolecular layer of the orthorhombic phase of RUB (d200 ) 1.34 nm).10 The number density of crystalline islands is of the order of 10-3 µm-2 (see Figure S1 in the Supporting Information). For I ) 0.20 nm/min, the crystalline phase becomes the dominant one, as shown in Figure 2c; here, the film surface is constituted by impinged islands, and second layer nucleation appears. Further on, for I > 0.2 nm/min (Figure 2d and e), the second layer coverage decreases down to zero when I ) 0.38 nm/min. This last sample is indeed formed only by a one-monolayer-high crystalline layer of RUB, covering almost all of the substrate surface. As a final but not less important remark on the film morphology, during image collection, we observed some postgrowth processes, particularly for films grown at I ) 0.10 nm/min, a condition that permits visualization of the regions where disordered dots and crystalline islands coexist over scales accessible to AFM (Figure 2b). Figure 3 shows the comparison of images collected after 12 h from extraction from the growth chamber over the same areas of the samples reported in Figure 2b and c. Some characteristics of these images deserve attention: (i) the crystalline RUB island observable on the top left of Figure
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Figure 3. Postgrowth evolution of the surface morphology of RUB thin films deposited on TEN at I ) 0.10 nm/min (a and b) and 0.20 nm/min (c and d). The as-grown films are reported in panels a and c, whereas the images in panels b and d are taken after 12 h. Cross-sectional profiles taken along the numbered horizontal white lines are also reported.
Figure 4. Surface morphology of a RUB thin film deposited on TEN at I ) 0.20 nm/min after permanence in an ultrahigh vacuum for 48 h.
3a is constituted by a single molecular layer (d200 ) 1.343 nm, see cross-sectional profile 1); (ii) the thickness of the same island is doubled in Figure 3b, whereas its contact area is fairly reduced (see profile 2); this effect is also evidenced in the sample deposited at higher I in profile 3 (left arrow); (iii) the dots surrounded by the crystalline island in Figure 3a disappear in Figure 3b, leaving an uncovered region; this is also shown in profiles 3 and 4 (right arrow) for the other sample; furthermore, a depletion region between the island and the dots is formed. We analyzed also the morphology of films kept in ultrahigh vacuum after growth for more than 48 h. As can be observed in Figure 4, dots are not present at all in these films whose surface is characterized by the presence of a high density of holes. The results reported in Figures 3 and 4 bring us to the following conclusions, which will be fundamental for validating the discussion below: (i) the nucleation of crystalline RUB on TEN occurs through a two-dimensional (2D) mechanism; (ii) nanodots supply spontaneously (both in air and in vacuum) molecules to crystalline RUB islands though a diffusion-limited process which increases the size and thickness of islands, while leaving voids at their original positions and a depletion region between them and the remaining dots (if any); (iii) monomo-
lecular RUB islands are morphologically unstable: spontaneous mass transport from the first to the second layer reduces their contact area with the substrate while forming a threedimensional (3D) crystalline domain. This same mass redistribution is observed in films grown at higher rates, as those corresponding to Figure 2d and e (see Figure S2 in the Supporting Information). As far as the disordered phase is concerned, it must be noted that the contact angle θ of the nanodots, as calculated from their height h and diameter D values with the formula θ ) 2 tan-1(2h/ D), is 17 ( 3°, independent of I; for such a low angle, the tipinduced diameter broadening is of the order of a few nanometers, then negligible with respect to the measured dot diameters. In the framework of classical nucleation thermodynamics, the free energy variation ∆Ga associated with the nucleation of an isotropic (amorphous) cap-shaped nucleus of N molecules from the vapor follows the expression:21
∆Ga ) -N∆µa + γa(36πf(θ))1/3(ΩsN)2/3
(1)
where γa is the energy of the amorphous solid-vapor interface, Ωs is the volume occupied by a molecule in the solid, and ∆µa is the energy associated with the transfer of a molecule from the vapor to the infinite amorphous solid (thermodynamic supersaturation); f(θ) ) [2 - 3 cos(θ) + cos3(θ)]/4, with θ being the contact angle of the nucleus with the substrate, is a geometric factor accounting for the specific adhesion energy βa of the amorphous nucleus with the substrate; indeed, in accordance with Young’s law:
cos θ ) (γsub - γint)/γa ) (βa /γa) - 1
(2)
with γsub being the energy of the substrate-vapor interface and γ int being the energy of the substrate-amorphous solid interface. The free energy barrier ∆G*a associated with the formation of amorphous cap-shaped nuclei on a foreign surface is obtained calculating the stationary point of the ∆Ga(N) function:
Kinetic Phase Selection of RUB Heteroepitaxial Domains
∆G*a )
16πγa3Ωs2 3∆µa2
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(3)
f(θ)
For films deposited under the same conditions, a variation of the number density of nuclei is ascribed to a variation of the energy barrier ∆G*a for nucleation. The kinetic expression relating the nucleation rate J with the energy barrier height is J ) A exp(- ∆G*/kBTs), where A is a kinetic constant, kB the Boltzmann constant, and Ts the substrate temperature (here 295 K); assuming a linear kinetics of the nucleation of dots, the ratio of the number density of dots before (n1) and after (n2) the increment of I is related to the variation of the energy barrier for nucleation ∆(∆G*) a through the equation
()
∆(∆G*) a ) kBTs ln
n1 n2
(4)
Considering the experimental n1 and n2 values for I ) 0.025 nm/min (Figure 2a) and I ) 0.10 nm/min (Figure 2b), eq 4 gives ∆(∆G*) a ) -46 ( 3 meV; such a decrement of the energy barrier for nucleation of amorphous dots is ascribed to the increase ∆(∆µa) of the thermodynamic supersaturation which accompanies the increment of I. In relation to solid-vapor equilibrium, the expression for the thermodynamic supersaturation is ∆µa ) kBTs ln(p/p∞-a), with p being the actual pressure of the vapor of RUB molecules in the growth chamber and p∞-a the vapor pressure of the infinite amorphous solid at Ts. Following kinetic theories, the vapor pressure p is related to the number of deposited molecules per unit surface per unit time (flux, F) through the equation p ) (2πmkBTs)1/2F, where m is the mass of one RUB molecule (8.83 × 10-25 kg). I (nm/ min) is related to F (m-2 s-1) through Ωs (for crystalline and amorphous RUB, Ωs can be assumed to be 700 Å3).10 Table 1 reports p values estimated from the measured I and the increments of supersaturation ∆(∆µ) ) kBTs ln(p2/p1) with respect to that corresponding to the lowest rate. From these values, the calculated decrease of the energy barrier ∆(∆G*) a ) -46 ( 3 meV can be related to an increment of the supersaturation ∆(∆µ) ) +35 meV. What we can try now is to calculate the actual supersaturation under the deposition conditions corresponding to Figure 2a and b. Equation 3 allows us to obtain an expression for ∆(∆G*) a as a function of ∆µ1 and ∆µ2, these being the supersaturation values corresponding to I ) 0.025 nm/min and I ) 0.10 nm/min, respectively:
∆(∆G*) a ) -
[
16πγa3Ωs2f(θ) ∆µ1 + ∆µ2 ∆(∆µ) 3 ∆µ12 + ∆µ22
]
(5)
Under the constraint ∆(∆µ) ) ∆µ2 - ∆µ1 ) +35 meV, eq 5 gives only one real positive solution for ∆µ1 (and, consequently, for ∆µ2). The major uncertainty in the quantities appearing in eq 5 rests on γa, whereas the values of all the other parameters are deduced either from the present experiment f(θ), and ∆(∆µ)) or from the literature (Ωs). Nonethe(∆(∆G*), a less, we can define a reasonable range for γa considering that, being θ ≈ 17°, eq 2 gives γa ) (γsub - γint)/0.96. Since the energy of the TEN(001) surface (γsub) is 5.0 meV/Å2, the previous relation gives an upper limit for γa, being γsub/0.96 ) 5.2 meV/Å2. Furthermore, we can consider the surface energy
TABLE 1: Influence of the Deposition Rate I on the Increment ∆(∆µ) of the Supersaturation as Well as the Estimated Supersaturation ∆µa and Energy Barrier ∆G*a (Assuming γa ) 5 meV/Å2) for the Growth of RUB Amorphous Thin Films; Flux F and Vapor Pressure p of the RUB Beam Are Also Reportedb p ∆(∆µ) ∆µa ∆G*a I F na (nm/min) (1014 m-2 s-1) (10-7 Pa) (meV) (meV) (meV) (µm-2) 0.025 0.10 0.20 0.26 0.38
6.0 24 48 62 91
0.90 3.6 7.2 9.3 14
35 53 59 70
110 140 160 170 180
120 75 55 45 35
0.43 3.0 6.2a
a Extrapolated with eq 4. b The measured and extrapolated dot density numbers na are reported in the last column.
of the most stable face of the RUB crystal (see section III.a) as the lowest limit of γa (4.61 meV/Å2). With these constraints, it follows that ∆µ1 ) 100-120 meV. Table 1 reports the estimates of the supersaturation for the nucleation of the amorphous phase for the explored interval of deposition rates. The explanation of the propensity of the TEN(001) surface to induce crystallization of RUB must be searched through a combined study of the thermodynamics and kinetics of nucleation. Being established that the RUB crystalline film phase nucleates through a 2D mechanism (Figures 2 and 3), we can express the free energy variation ∆Gc associated with the nucleation of a 2D, (100)-oriented, disk-shaped nucleus of N RUB molecules from the vapor with the following equation:21
∆Gc ) -N∆µeff + 2γl(πhΩsN)1/2
(6)
where h is the height of the cluster (13.43 Å)10 and γl is the surface energy of the edge of the cluster. The term ∆µeff ) ∆µc - Ωs(2γ200 - βc)/h ) ∆µc - ∆µ0, with ∆µc being the energy associated with the transfer of a molecule from the vapor to the infinite crystalline solid, γ200 the energy of the (200) surface, and βc the specific adhesion energy of the RUB(200) surface with TEN(001) (being 4.61 and 6.6 meV/Å2, respectively; see section III.a), represents the effective supersaturation during 2D nucleation; the quantity ∆µ0 ) 140 meV is a threshold supersaturation required for 2D nucleation. When 2γ200 - βc > 0 and ∆µc < 2∆µ0, 3D nucleation remains favored vs 2D nucleation; on the contrary, when the transition supersaturation 2∆µ0 is overcome, 2D nucleation is the favored mechanism.16a,b,22 Since all of the crystalline nuclei observed in our films are 2D, we can assume that for I ) 0.10 nm/min ∆µc g 2∆µ0 ) 280 meV. This value of supersaturation corresponds to an equilibrium vapor pressure at room temperature of bulk crystalline RUB 2 orders of magnitude lower than that of bulk amorphous RUB (p∞-a). The energy barrier ∆G*c associated with the nucleation of the 2D crystalline phase of RUB is calculated differentiating eq 6:
∆G*c )
πhγl2Ωs ∆µeff
(7)
Figure 5 reports the curves representing ∆Ga(N) and ∆Gc(N) calculated with eqs 1 and 6 for the deposition rate interval I ) 0.025-0.38 nm/min, setting γa ) 5 meV/Å2 and γl ) 6 meV/ Å2, the lowest surface energy calculated for the edge of a (100)oriented cluster (see section III.a).
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Figure 5. Free energy variation ∆G associated with the formation of an amorphous (left curves) and crystalline (right curves) cluster of N RUB molecules from a vapor on the TEN(001) surface. The thermodynamic parameters are γa ) 5 meV/Å2, γl ) 6 meV/Å2, ∆µa ) 110, 140, 160, 170, and 180 meV, and ∆µeff ) 140, 160, 170, and 180 meV.
The comparison of the curves in Figure 5 provides the explanation of the growth dynamics of RUB thin films. The nucleation of crystalline RUB films is a possible event for any I value explored in this experiment; nonetheless, the free energy barrier associated with this nucleation is considerably higher than that associated with the nucleation of amorphous nanodots (note the different ordinate scales in Figure 5). This implies the nucleation rate of crystalline RUB to be extremely low compared to that of amorphous dots. Consistently with this description, Figure 2a does not show the presence of any crystalline island, whereas Figure 2b, corresponding to films grown with a 35 meV increment of ∆µ, does show crystalline islands, but with a number density 3 orders of magnitude lower than that of nanodots. The scenario undergoes an abrupt change by a further 18 meV increment of ∆µ, giving rise to a film phase almost completely composed by crystalline RUB (Figure 2c). This supersaturation-dependent selection of amorphous and crystalline domains can find a quantitative explanation from Figure 5. It is observed that the increment of ∆µ gives rise to a decrement of ∆G*c and ∆G*a to fairly different extents (eqs 3 and 7): in the crucial step between Figure 2b and c, ∆G*a is decreased by 20 meV, whereas ∆G*c is decreased by a quantity about 1 order of magnitude larger. Following eq 4, the number density of dots is expected to double (see last column of Table 1), whereas that of islands is expected to augment to more than 3 orders of magnitude. In accordance with this prediction, the film in Figure 2c presents a crystalline layer of well connected domains, whereas in Figures 2d and e the density of nuclei is so high that different domains cannot be distinguished any more; on the contrary, the number density of dots in Figure 2c is found to be of the order of 10-1 µm-2, whereas dots are absent in Figure 2d and e. What is the reason for this reduction of the number of dots? The explanation comes from the analysis of the postgrowth behavior of the heterostructure, presented in Figures 3 and 4. As previously described, the postgrowth changes of the film morphology are the result of the onset of interdomain and interlayer mass transport; the former phenomenon consists of the migration and/or incorporation of molecules from dots to islands, the driving force for this effect being the difference of the chemical potential of a molecule in a dot and that of a molecule in a 2D island. Starting from a dot with a certain size, the transfer of a molecule from it to a nearby island gives rise to an increase of the chemical potential of the molecules in the dot, while that of molecules in the island decreases, with an overall loss of free energy. This implies that the process is self-sustained, bringing to an efficient and rapid
Campione et al. transition toward the crystalline phase. The transformation of clusters with a size larger than a critical value from amorphous to crystalline is expected to be a spontaneous event.23 This event is not demonstrated to occur in our films; rather, molecules are observed to transfer from amorphous to crystalline domains. This process must be all the more so active during the deposition experiment, being enhanced by the thermalization energy of incoming molecules; this demonstrates that, besides being a spontaneous event, the transfer of molecules from amorphous to crystalline domains occurs through a rapid kinetics, contrary to the direct transformation of dots from amorphous to crystalline, which is probably kinetically hindered. The literature provides several examples of thickness-dependent amorphous-to-crystalline transformations occurring in thin film deposition experiments. An extensively studied system is represented by Mo thin films which above a thickness as low as 2 nm undergo a amorphous-to-crystalline transition on a variety of substrates and under different growth conditions.24 Similarly, metal thin films on silica substrates were observed to display such a size-dependent transformation,23 even under athermal conditions.25 The trivial case of thermally induced crystallization by agglomeration of an amorphous layer was reported for example for Si and Ge thin films,26 and it was documented also in the case of RUB thin films.27 The interesting phenomenon of spherulitic growth from an amorphous organic layer during deposition by thermal evaporation in a vacuum was recently reported for p-nitrophenyl nitronyl nitroxide.28 However, as far as we are aware, the peculiar growth dynamics described and rationalized here from the thermodynamic and kinetic standpoint for RUB thin films on TEN, involving competing nucleation events and a strong mass transport from amorphous to crystalline domains continuing even during the postgrowth stage, was not observed in other systems. III.c. Concluding Remarks. The interlayer mass transport within a single crystalline domain observed in Figure 3 is a phenomenon recently observed also in TEN thin films deposited on silicon oxide.29 Monolayer islands of TEN were observed to increase their thickness with time while changing their morphology: molecules adsorbed at the edge of the island migrate on top of the island and nucleate the second layer. This migration was believed to be activated by the presence of small second layer nuclei and continues until completion of three monolayers. These phenomena of molecular migration, both from a lower to an upper crystalline layer and from an amorphous dot to a crystalline domain, are observed ex situ under ambient conditions; hence, water, oxygen, and other atmospheric agents adsorbed on the film surface can be proposed to play a mediation role; moreover, it can not be excluded that a significant fraction of the migrating molecules undergo oxidation; further investigations are needed for individuating the actual dynamics and driving forces. However, Figure 4, showing a hole-rich morphology of RUB thin films kept in a vacuum after growth, demonstrates that molecular migration from dots to crystalline domains can occur also in the absence of atmospheric agents. In conclusion, the favorable adhesion of the RUB(100) face on TEN(001), achieved through the matching of the direction and spacing of their main surface corrugations, is the only requirement making the nucleation of crystalline films of RUB possible. During deposition, amorphous domains coexist with crystalline ones; however, the mass transfer from the former to the latter ones is a spontaneous and rapid process and, by a proper tuning of the growth conditions, it can be exploited for forming a highly crystalline and uniform thin film. On the other
Kinetic Phase Selection of RUB Heteroepitaxial Domains hand, the unsuccessful growth of crystalline RUB thin films reported in several experiments can only be ascribed to an insufficient adhesion of the crystalline overlayer on the substrate. This appears in contrast to the recent literature, which designates the conformational change RUB molecules must undergo during the transition from vapor to crystal, requiring to overcome an additional kinetic barrier for nucleation of 0.21 eV, as the main mechanism hindering the growth of crystalline thin films.30 However, this kinetic barrier does not have any effective role in preventing the nucleation of crystalline RUB films as homoepitaxial RUB thin films grown under very low rates, at which the material forms only amorphous dots on TEN (I ) 0.025 nm/min, and even at I ) 0.010 nm/min), show the presence of exclusively crystalline 2D islands nucleated at defect-free positions (see Figure S3 in the Supporting Information). On the contrary, if βc is too low for any orientation of the overlayer, the threshold supersaturation making crystalline nucleation competitive with amorphous nucleation may not be achievable under the OMBE conditions. IV. Conclusions By a combined approach based on AFM characterization, force field calculations, and the application of classical nucleation thermodynamics, we elucidated a new growth mechanism for thin films of molecular organic materials possessing a metastable amorphous solid phase. This involves an amorphousto-crystalline transition occurring through the diffusion of molecules from amorphous to crystalline domains. This growth dynamics has been demonstrated for the OMBE growth of heteroepitaxial films of RUB at room temperature: a morphological transition from amorphous nanodots to a uniform crystalline monomolecular layer can be gradually induced by tuning the molecular beam flux. These findings can serve as valuable guidelines in the controlled nanostructuring of organic interfaces and in the development of an integrated organic electronics based on one of the most prominent organic semiconductors, i.e., RUB. Furthermore, the described growth mode may in principle occur for all molecular organic materials displaying significant differences between the equilibrium molecular conformation in the vapor phase and that in the crystalline solid. Acknowledgment. This work was supported by Fondazione Cariplo (Grant No. 2007/5205). Enrico Fumagalli is kindly acknowledged for his support in the AFM measurements. Supporting Information Available: Figures showing the AFM morphology of heteroepitaxial RUB films grown with a low deposition rate, postgrowth morphology of heteroepitaxial thin films grown at a high rate, and AFM morphology of homoepitaxial RUB films grown with a low deposition rate. This material is available free of charge via the Internet at http:// pubs.acs.org. References and Notes (1) (a) Podzorov, V.; Menard, E.; Borissov, A.; Kiryukhin, V.; Rogers, J. A.; Gershenson, M. E. Phys. ReV. Lett. 2004, 93, 086602. (b) Zeis, R.; Besnard, C.; Siegrist, T.; Schlockermann, C.; Chi, X. L.; Kloc, C. Chem. Mater. 2006, 18, 244.
J. Phys. Chem. C, Vol. 113, No. 49, 2009 20933 (2) (a) Haemori, M.; Yamaguchi, J.; Yaginuma, S.; Itaka, K.; Koinuma, H. Jpn. J. Appl. Phys. 2005, 44, 3740. (b) Hu, W.-S.; Weng, S.-Z.; Tao, Y.-T.; Liu, H.-J.; Lee, H.-Y. Org. Electron. 2008, 9, 385. (3) Park, S. E.; Jeong, S. H.; Choi, J.-M.; Hwang, J. M. Appl. Phys. Lett. 2007, 91, 033506. (4) Zeng, X.; Wang, L.; Duan, L.; Qiu, Y. Cryst. Growth Des. 2008, 8, 1617. (5) (a) Le Moigne, J.; Kajzar, F.; Thierry, A. Macromolecules 1991, 24, 2622. (b) Koma, A.; Yoshimura, K. Surf. Sci. 1986, 174, 556. (c) Wang, H.; Zhu, F.; Yang, J.; Geng, Y.; Yan, D. AdV. Mater. 2007, 19, 2168. (d) Mannsfeld, S. C. B.; Leo, K.; Fritz, T. Phys. ReV. Lett. 2005, 94, 056104. (e) Campione, M.; Sassella, A.; Moret, M.; Papagni, A.; Trabattoni, S.; Resel, R.; Lengyel, O.; Marcon, V.; Raos, G. J. Am. Chem. Soc. 2006, 128, 13378. (6) Campione, M. J. Phys. Chem. C 2008, 112, 16178. (7) Laudise, R. A.; Kloc, Ch.; Simpkins, P. G.; Siegrist, T. J. Cryst. Growth 1998, 187, 449. (8) Holmes, D.; Kumaraswamy, S.; Matzger, A. J.; Vollhardt, K. P. C. Chem.sEur. J. 1999, 5, 3399. (9) Campione, M.; Cartotti, M.; Pinotti, E.; Sassella, A.; Borghesi, A. J. Vac. Sci. Technol., A 2004, 22, 482. (10) Jurchescu, O. D.; Meetsma, A.; Palstra, T. T. M. Acta Crystallogr. 2006, B62, 330. (11) Morris, G. M.; Goodsell, D. S.; Halliday, R. S.; Huey, R.; Hart, W. E.; Belew, R. K.; Olson, A. J. J. Comput. Chem. 1998, 19, 1639. (12) Gavezzotti, A. Molecular Aggregation: Structure Analysis and Molecular Simulation of Crystals and Liquids; Oxford University Press: Oxford, NY, 2007; p 215. (13) Stone, A. J.; Dullweber, A.; Engkvist, O.; Fraschini, E.; Hodges, M. P.; Meredith, A. W.; Nutt, D. R.; Popelier, P. L. A.; Wales, D. J. Orient: a Program for Studying Interactions between Molecules, version 4.6; University of Cambridge: Cambridge, U.K., 2002. Enquiries to A. J. Stone at
[email protected]. (14) Overney, R. M.; Howald, L.; Frommer, J.; Meyer, E.; Gu¨ntherodt, H.-J. J. Chem. Phys. 1991, 94, 8441. (15) Hartman, P.; Perdock, W. G. Acta Crystallogr. 1955, 8, 49. (16) (a) Campione, M.; Sassella, A.; Moret, M.; Marcon, V.; Raos, G. J. Phys. Chem. B 2005, 109, 7859. (b) Verlaak, S.; Steudel, S.; Heremans; Janssen, D.; Deleuze, M. S. Phys. ReV. B 2003, 68, 195409. (c) Nabok, D.; Puschnig, P.; Ambrosch-Draxl, C. Phys. ReV. B 2008, 77, 245316. (17) Haber, T.; Resel, R.; Thierry, A.; Campione, M.; Sassella, A.; Moret, M. Physica E 2008, 41, 133. (18) (a) Hillier, A. C.; Ward, M. D. Phys. ReV. B 1996, 54, 14037. (b) Haber, T.; Muelleger, S.; Winkler, A.; Resel, R. Phys. ReV. B 2006, 74, 045419. (19) Campione, M.; Raimondo, L.; Sassella, A. J. Phys. Chem. C 2007, 111, 19009. (20) Luo, Y.; Brun, M.; Rannou, P.; Grevin, B. Phys. Status Solidi A 2007, 204, 1851. (21) Mutaftschiev, B. The Atomistic Nature of Crystal Growth; Springer Series in Materials Science; Springer: Berlin, 2001. (22) Sassella, A.; Campione, M.; Papagni, A.; Goletti, C.; Bussetti, G.; Chiaradia, P.; Marcon, V.; Raos, G. Chem. Phys. 2006, 325, 193. (23) Hu, M.; Noda, S.; Komiyama, H. J. Appl. Phys. 2003, 93, 9336. (24) (a) Bajt, S.; Stearns, D. G.; Kearney, P. A. J. Appl. Phys. 2001, 90, 1017. (b) Schubert, E.; Ma¨ndl, S.; Neumann, H.; Rauschenbach, B. Appl. Phys. A 2005, 80, 47. (c) Nayak, M.; Lodha, G. S.; Nandedkar, R. V. J. Appl. Phys. 2006, 100, 113709. (d) Lo¨hmann, M.; Klabunde, F.; Bla¨sing, J.; Veit, P.; Dru¨sedau, T. Thin Solid Films 1999, 342, 127. (25) Ekinci, K. L.; Valles Jr, J. M. Acta Mater. 1998, 46, 4549. (26) (a) Grom, G. F.; Lockwood, D. J.; McCaffrey, J. P.; Labbe´, H. J.; Fauchet, P. M.; White Jr, B.; Dlener, J.; Kovalev, D.; Koch, F.; Tsibeskov, L. Nature 2000, 407, 358. (b) Wakayama, Y.; Tagami, T.; Tanaka, S. Thin Solid Films 1999, 350, 300. (27) (a) Hsu, C. H.; Deng, J.; Staddon, C. R.; Beton, P. H. Appl. Phys. Lett. 2007, 91, 193505. (b) Park, S.-W.; Jeong, S. H.; Choi, J.-M.; Hwang, J. M.; Kim, J. H.; Im, S. Appl. Phys. Lett. 2007, 91, 033506. (c) Nothaft, M.; Pflaum, J. Phys. Status Solidi B 2008, 245, 788. (28) Caro, J.; Fraxedas, J.; Figueras, A. J. Cryst. Growth 2000, 209, 146. (29) Shi, J.; Qin, X. R. Phys. ReV. B 2008, 78, 115412. (30) Ka¨fer, D.; Ruppel, L.; Witte, G.; Wo¨ll, Ch. Phys. ReV. Lett. 2005, 95, 166602.
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