Crystal Shape Engineering of Silicon Nanoparticles in a Thermal

Article pubs.acs.org/crystal

Crystal Shape Engineering of Silicon Nanoparticles in a Thermal Aerosol Reactor Richard Körmer,† Benjamin Butz,‡ Erdmann Spiecker,‡ and Wolfgang Peukert*,† †

Institute of Particle Technology, Friedrich-Alexander University Erlangen-Nuremberg, Cauerstrasse 4, 91058 Erlangen, Germany Center for Nanoanalysis and Electron Microscopy (CENEM), Friedrich-Alexander University Erlangen-Nuremberg, Cauerstrasse 6, 91058 Erlangen, Germany

‡

ABSTRACT: In this work, the capability of gas phase synthesis for crystal shape engineering of silicon nanoparticles (SiNPs) in a hot wall reactor is demonstrated. Therefore, the necessary boundary conditions for the formation of monodisperse spherical SiNPs as well as octahedral-shaped particles from silane pyrolysis are systematically deduced. The different shapes of the SiNPs are ascribed to different growth regimes (reaction limitation, diffusion limitation) depending on the global process parameters. Single crystalline, defect-free, spherical particles with a mean diameter of 30 nm (geometric standard deviation below 1.1) and octahedra with a mean edge length of about 100 nm could be obtained solely by process parameter adjustment. Particle size and shape as well as crystallinity were characterized by scanning electron microscopy and X-ray diffraction. The inner structure and faceting of the particles were analyzed in detail by high resolution transmission electron microscopy. A model that elucidates the orientation relation between the inner silicon structure and the particle shape is derived for the octahedra.

1. INTRODUCTION Challenges in current nanoparticle synthesis aim far beyond the control of particle size. Especially for functional particle systems, for example, thin films for electronic applications, where the dense packing of primary building blocks to hierarchical superstructures is preferred, the particle shape becomes of primary importance. Since most functional particle systems are crystalline materials, the term crystal shape engineering comes up. It means a process design by which the crystallization environment is adjusted to obtain the desired crystal shapes.1 Huge progress has been made in liquid phase synthesis of nanoparticulate building blocks: A large variety of shapes, for example, rods, cubes, and tetrapods, have been synthesized for many materials (e.g., ZnO, PbS or CdSe and others). A broad overview is given in the review by Talapin et al.2 The successful oriented deposition of these building blocks to densely packed films has been demonstrated for ZnO nanorods,3 which resulted in distinct improved performance of thin film transistors. Nanomaterials from group-IV semiconductors (especially silicon or germanium) lag behind this rapid evolution due to the unsolved problem of surface degradation. Nevertheless silicon is the dominating material in the conventional semiconductor industry because of its superior © 2012 American Chemical Society

properties such as doping capability, nontoxicity, and abundance. Some progress has been made in the field of electrical application by the usage of single silicon nanowires as field effect transistors4,5 or solar cells6 because of the absence of interparticulate boundaries. However, nanowires are not compatible with printing techniques and require extensive templateinduced growth processes. On the other hand, dispersions made from freestanding silicon nanoparticles (SiNPs) turned out to be suitable candidates for device fabrication.7−9 Freestanding SiNPs are typically synthesized via a classical gas phase route due to purity reasons, whereupon the most promising results have been obtained by thermal,10−12 laser-heated,13−15 and microwave16,17 or plasma18 assisted pyrolysis of silane. In those studies, SiNPs of undefined shape or spherical particles  separated as well as aggregated  are observed. Spherical SiNPs with narrow particle size distribution (PSD), which is not self-evident for aerosol processing, were reported by Shen et al.19 and Körmer et al.20 The mechanism that leads to such particles is explained by the dominance of the surface growth Received: October 21, 2011 Revised: December 16, 2011 Published: January 4, 2012 1330

dx.doi.org/10.1021/cg201394y | Cryst. Growth Des. 2012, 12, 1330−1336

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combination with an overfocus of a few nanometers results, providing a small sample thickness, that is, at the tip of the octahedral SiNP, in high intensity at positions where the atomic columns are situated (reduced intensity in between) and a negligible delocalization of image contrast at interfaces and surfaces.30 For TEM investigation, the particles were dispersed onto Lacey C-films by dropping an ultrasonically treated solution (SiNP in ethanol) directly onto the supporting film. This locally led to particles or small agglomerates that protruded from the supporting film. 2.4. X-ray Diffraction (XRD). XRD scans were carried out in Bragg−Brentano geometry with a Bruker AXS Advance D8 X-ray diffractometer equipped with a VANTEC-1 detector using Cu Kα radiation. For analysis, the powder was placed in low background sample cups with a vicinal (911) Si crystal of 25 mm diameter. The XRD patterns of the samples were recorded using a step size of 0.02° and a counting time of 3 s per step. The average crystallite size was determined by Rietveld refinement using Bruker AXS TOPAS 3 software. The XRD patterns were numerically corrected for Kα1/Kα2 splitting of the reflections.

reaction at negligible collision rates for the particles among themselves.21 This enables the production of SiNPs with a welldefined size and shape. Nevertheless, a broad variety of SiNPs building blocks with different shapes, similar to prospects accessible by liquid phase routes,22,23 is yet not accessible via an aerosol route. The concepts of liquid phase synthesis, where suitable capping agents or oriented attachment enable the growth of distinct particle shapes, cannot be transferred to gas phase processes. An indication for octahedral SiNPs has only been mentioned by Murthy et al.24 who examined undesired particle formation in epitaxial growth processes based on silanes and chlorosilanes. Kortshagen's group observed faceted SiNPs with a cubic shape from a gas phase process utilizing a nonthermal plasma generated by radiofrequency (RF).25,26 They explain the origin of the cubic-shaped particles by a restructuring of amorphous, cauliflower-shaped particles in the plasma zone due to surface diffusion.26 A key role is ascribed to the hydrogen terminated surfaces, as proposed by Barnard & Zapol:27 The cubic shape was predicted as equilibrium shape for fully hydrogenated SiNPs by taking the Gibbs free energy for bulk, surface, edges, and corners into account. Another interesting approach is the work from Hawa & Zachariah, who performed molecular dynamics simulations on SiNPs.28 Their results confirm the cubical shape as the most stable one for hydrogenated SiNPs. In this study, we demonstrate the synthesis of highly symmetrical octahedral SiNPs. The inner structure and shape are characterized by X-ray diffraction, scanning electron microscopy, and high-resolution transmission electron microscopy to derive a structural model (crystal structure, shape, faceting). Since spherical SiNPs with excellent quality (narrow PSD, monocrystalline) have been obtained using the same reactor system at different experimental parameters,20 the formation of both, spherical as well as octahedral SiNPs, is compared to gain information about the differences in crystal growth.

3. RESULTS We reported the production of spherical, crystalline SiNPs with narrow PSD in the same experimental setup before.20 This has been achieved for the following set of process conditions: Reduced total pressure (ptotal) of 25 mbar, furnace temperature (Tfurnace) in the range of 900−1100 °C, mean residence times in the furnace (τR) between 80 and 420 ms, and silane partial pressures (psilane) equal or less than 1 mbar. Slightly aggregated particles with primary particle sizes within the size range of 20− 40 nm at a geometric standard deviation (GSD) below 1.1 have been produced in this process parameter regime. Mostly separated spherical particles with a mean diameter of approximately 30 nm and a GSD of 1.05 can be produced by reducing the total pressure to 15 mbar as representatively shown in Figure 1a. The larger agglomerates visible in Figure 1a

2. EXPERIMENTAL SECTION 2.1. Reactor Setup. SiNPs are produced in a resistively heated, vertical hot wall reactor by thermal pyrolysis of monosilane (SiH4). Monosilane (Linde, grade 5.0) is diluted in argon (Linde, grade 6.0 additionally cleaned via gettering) and injected at the bottom of the reactor. The reactor consists of a quartz tube with an inner diameter of 36 mm and the heated length is 500 mm. Behind the reactor the aerosol is passed through a thermophoretic sampler that consists of a water-cooled quartz tube (length: 1000 mm). This device allows the sampling of a particle fraction close to the wall due to the thermal gradient. The remaining particles are separated by a standard membrane filter. Additionally, aerosolized particles at the reactor exit can be sampled locally at arbitrary positions in the exit cross section by extraction with a hooked probe (4 mm inner diameter). These samples are deposited on silicon wafer substrates by the use of a low pressure impactor (LPI). A schematic drawing of the reactor and more details can be found elsewhere.20 2.2. Scanning Electron Microscopy (SEM). Topographic SEM images were obtained using a Zeiss GEMINI ULTRA 55. The measurements were carried out at an acceleration voltage of 8 kV using the in-lens detector. The particles were either deposited by the use of the LPI or via dropping a solution of dispersed particles in ethanol on a silicon wafer substrate. The samples were examined by SEM without further treatment. 2.3. Transmission Electron Microscopy (TEM). Transmission electron microscopy was carried out using an image-side aberrationcorrected FEI Titan3 80−300 microscope. To minimize electron-beam induced knock-on damage, the accelerating voltage was set to 200 kV. For high-resolution TEM (HRTEM) imaging the NCSI (negative Cs imaging) technique was applied.29 A negative Cs value of −5 μm in

Figure 1. SEM images of (a) spherical SiNPs at low reactor total pressure (ptotal = 15 mbar, Tfurnace = 1100 °C, psilane = 0.5 mbar, τR = 80 ms), (b) mostly octahedral SiNPs from the near wall region (ptotal = 100 mbar, Tfurnace = 1100 °C, psilane = 1 mbar, τR = 80 ms). The insets (edge length 150 nm) show single NPs at higher magnification.

are mainly due to the deposition technique via low pressure impaction. The situation changes significantly if the process conditions are modified. For increased total pressure of 100 mbar (Tfurnace = 1100 °C, psilane = 0.4−2.5 mbar, τR = 80 ms) faceted, mostly octahedral-shaped particles are observed close to the reactor wall as shown in the SEM image in Figure 1b. Freestanding octahedra as well as slightly aggregated ones have been found. Apparently a high fraction of octahedral SiNPs of approximately 50% is visible. The octahedra exhibit eight equilateral facets with an average edge length of around 100 nm as it is 1331



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shown in the upper inset in Figure 1b. These particles are efficiently separated from the core flow at the beginning of the thermophoretic sampler or by positioning the hooked probe close to the wall. Nevertheless, a fraction of differently shaped particles remains in those samples. These faceted SiNPs are reproducibly observed only in the above-mentioned, narrow regime of operation parameters of the reactor. To exclude any influence of the specific reactor system on the occurrence of the octahedral SiNPs, similar particles were synthesized by operating a comparable, second reactor system at the same experimental conditions. Therefore we attribute the appearance of octahedral SiNPs to the applied crystallization conditions rather than to any contamination effects. Table 1. Typical Experimental Parameters for the Synthesis of Spherical and Octahedral SiNPs Tfurnace/°C ptotal/mbar psilane/mbar τR/ms

spherical SiNPs

octahedral SiNPs

1100 15 0.5 80

1100 100 1 80

Figure 3. Bright-field TEM image of mostly octahedral SiNPs from the near wall region as shown in Figure 1b. Several nonoctahedral SiNPs are indicated: undefined SiNPs with inner defects are indicated by arrows. A polycrystalline aggregate is marked by the dotted boundary.

octahedra appear in various shapes (square, rectangle, hexagon, or rhomb) depending on their orientation on the Lacey C-film. The alternating and highly symmetrical contrast (thickness fringes) is due to the symmetry of each octahedron and strongly depends on its specific orientation. A few undefined SiNPs with inner defects are exemplarily marked by arrows  the dashed border indicates a larger agglomerated and sintered particle. The homogeneous contrast of the octahedral SiNPs, except the alternating thickness-dependent fringes, clearly confirms that these particles are  as it was recently found for the spherical SiNPs20  single crystalline and free of extended defects such as stacking faults and dislocations. Deviating from the ideal octahedral shape, the synthesized SiNPs exhibit slightly radiused edges and flattened tips. This can be seen in detail in the following HRTEM images. To correlate the inner crystal structure and shape, Figure 4 representatively shows the bright-field image of a single octahedral SiNP aligned along the [110] zone axis. The periodic thickness fringes are due to the linear increase in thickness from both tips (top and bottom) to the inner part of the particle. Two characteristic features are marked by white squares (b, c), from which HRTEM images were recorded: (b) a tip of the octahedron, (c) the projection of the SiNP along an [110] edge. Both HRTEM images (Figure 4b,c) indicate that the large facets of the SiNP are determined by {111} planes. This can also be seen from the calculated FFT of Figure 4c shown as an inset in Figure 4a. Furthermore, the rounded tips, observed in SEM and bright-field TEM (Figure 1b, Figure 3), are additionally faceted on the nanoscale. Despite the native oxide around the SiNP (amorphous layer in Figure 4b,c), the local tip shape is revealed in Figure 4b. Apparently, well-marked {113} and {001} facets terminate the tip of the SiNP. The thickness distribution at the tip (see Figure 4b, thickness increases from top to bottom) corresponding to the (001), (113), and (1̅1̅3) facets (along beam direction) was verified by simulating the high-resolution pattern in this region. Therefore, a super cell (4 unit cells wide, surfaces according to faceting) of the central slab of the tip was modeled. The simulated contrast variations due to (i) the change in local thickness and (ii) the change of the local defocus value (inclined surfaces of the NP) are in good agreement with those, observed in the high-resolution image. For comparison the simulation (Cs = −5 μm, defocus

The structural homogeneity of the batch of octahedral SiNPs is demonstrated by the XRD pattern (Figure 2) that exclusively

Figure 2. XRD patterns of spherical SiNPs in comparison to octahedral-shaped SiNPs.

shows the characteristic reflections due to the cubic silicon lattice. The peak widths for spherical SiNPs (Figure 1a) are broader in comparison to those for the batch of mainly octahedral SiNPs (Figure 1b). This can be explained by a smaller average crystallite size in contrast to the octahedra. For the spherical SiNPs, the crystallite size from XRD analysis is in good agreement with the SEM images due to the narrow PSD (dcrystallite,spheres= 29 nm). This leads to the conclusion that the spheres are monocrystalline and largely free of extended defects. Since the octahedra show more broadly distributed sizes (and in addition partly aggregated, undefined SiNPs), it is difficult to compare the SEM images with the value for the mean crystallite size (dcrystallite,octahedrons = 49 nm). Nevertheless the SEM images already indicate much larger structures for the octahedra compared to the spheres. The bright-field TEM image in Figure 3 shows an ensemble of randomly oriented, octahedral SiNPs. In the projection, the 1332



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showed the inner 4-fold symmetry axis || ⟨001⟩ from tip to tip (Figure 4). The {111} type of the major facets of the SiNP was clarified (Figure 4b,c). Hence, the shape is mainly characterized by the 8 {111} facets that are defined by the {111} planes of the cubic silicon structure. Furthermore, the presence of small {113} and {001} facets at the flattened tips could be revealed. These smaller facets are representatively indicated for one enlarged tip in Figure 5. The formation of pronounced surface facets points to a growth mechanism that allows the SiNPs to establish a lowenergy shape (see below). The equilibrium (or Wulff) shape of silicon has been extensively studied, both experimentally and by computer simulation, in the literature. Eaglesham et al.31 observed similar truncated octahedra for faceted voids in silicon (several 10 nm in diameter) which formed after high-energy helium implantation upon annealing at 700 °C. Bermond et al.32 in situ annealed as-prepared silicon bulbs (∼800 nm; unfaceted, round shape) at temperatures up to 1050 °C and observed well-defined {111} and {113} facets. These findings are in nice agreement with the present observations. However, the fractional area of the individual facets, which is determined by the relative surface energies, strongly depends on the experimental conditions (temperature, gas environment, etc.) indicating that details of the surface, such as reconstruction and adsorption, play an important role in the competing development of the facets. Unfortunately, in the present work oxidization of the SiNPs after extraction impeded a more detailed study on surface reconstruction and atomic faceting at the tips and edges. Hence, any instability of the {113} facets and their faceting into low-energy {111} and {100} facets on the atomic scale, as controversially discussed in the literature,33−35 could not be investigated. 4.2. Dependence of the Crystal Shape on the Experimental Growth Conditions. A large diversity of differently shaped SiNPs at various sizes has already been reported in the literature for aerosol processing. Mostly aggregated particles with undefined shapes were observed for thermal hot wall reactors operated at atmospheric pressure or slightly reduced total pressure at relatively high silane partial pressure as summarized in Table 2. In this table, an overview of the relevant process parameters is given. From the mechanistic point of view, the formation of such particles can be attributed to classical concepts of particle formation in gases.36,37 The chemical reaction of the precursor leads to the formation of a supersaturation that subsequently drives nucleation. Since the nucleation rates are typically very high in aerosol synthesis, agglomeration and sintering sets in. All these  typically simultaneously ongoing  processes lead to aggregated particles with a broad distribution in size and shape. For the synthesis of spherical particles two concepts can be utilized: The first concept relies on the sintering of aggregated particles at high temperatures. This has been demonstrated by the group of Flagan12,41 (hot wall reactor, Tfurnace = 1250 °C) and by Flint & Haggerty42 (laser reactor) in a temperature range of 1280−1605 °C, partly above the melting point of silicon (Tmelt, Si = 1410 °C). Nevertheless, this concept does not result in particles with narrow size distributions. The second concept is based on a mechanism governed by surface growth due to chemical vapor deposition (CVD) and leads to spherical nanoparticles with a narrow PSD. Holunga et al.43 described a thermal, atmospheric, turbulent mixing aerosol reactor operated at a silane concentration on the order of 100 ppm for the synthesis of monodisperse, spherical SiNPs. Similar observations

Figure 4. (a) Bright-field TEM image of a single silicon octahedron oriented along the [110] zone axis. (b, c) HRTEM images of the tip as well as along an edge of the octahedron (as indicated in a) revealing the inner crystal structure and faceting. The amorphous outer layer with a thickness of less than 1 nm is attributed to the native silicon oxide. The FFT of such HRTEM image (inset in a) shows the orientation of the lattice with respect to the facets of the SiNP.

ΔZ = +8 nm, high tension 200 kV) is superimposed to the experimental data in the inset in Figure 4b.

4. DISCUSSION 4.1. Crystal Structure and Shape. In conclusion to the experimental findings, a model (Figure 5) is deduced that

Figure 5. Model of the octahedral SiNP with slightly truncated tips: the directions and facets are indexed in the cubic indexing scheme of the silicon structure. For better visibility one tip is shown enlarged.

describes the orientation relation between the inner silicon structure and the shape of the octahedral SiNPs. The cubic silicon structure (Fd3m̅ , a = 5.43 Å) of the mainly octahedral SiNPs was verified by XRD (Figure 2) and HRTEM (Figure 4). Bright-field TEM in combination with HRTEM in detail 1333



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Table 2. Synthesis Parameters of SiNPs without Defined Shape (Literature Overview) reference Tfurnace/°C ptotal/mbar psilane/mbar reactor type

Wiggers10

Nguyen12

Onischuk38

Onischuk39

Cannon13

van Erven40

1000 ∼1000 100−400 hot wall

800−1000 ∼1000 1.7 hot wall

517−727 390 0.06−19.5 hot wall

567−667 66 3.3 hot wall

955−1060 200 200 laser

800 40−80 laser

have been made by Shen et al.19 for a plasma reactor at very low total pressures of pure silane (0.01−0.1 mbar). They reported that argon dilution significantly increases the particle production rate. This finding agrees well with the results of Körmer et al.20 for a hot wall reactor, where spherical SiNPs with narrow PSDs could exclusively be obtained if the reactor total pressure (ptotal = 25 mbar) as well as the silane partial pressure (psilane ≤ 1 mbar) were low. The growth mechanism is described by a short nucleation burst and subsequent growth mainly via surface reaction and condensation.21 A crucial role was attributed to the total pressure since argon atoms, which are used for dilution, act as collision bodies for the pyrolysis of the precursor molecules.44 The low total pressure leads to high diffusion rates of reactive species present in the gas phase. The spherical shape of the particles is therefore explained by the presence of a reaction limited growth regime according to classical crystallization theory.45 The deposition rate of the crystallization units onto the surface of the growing SiNPs exceeds the rate of surface diffusion by far. Hence, the incorporation in the crystal lattice cannot occur at its energetically favored sites. These statements are valid in the same manner for the spherical particles discussed in this work. The further reduced total pressure in this study (ptotal = 15 mbar) is responsible for a pronounced equalization of the reactive species leading to almost no dependence of the particle properties (size, shape) on the radial sampling position. Consequently, faceted particles can be obtained if the incorporation of crystallization units at energetically favored sites is facilitated. This has rarely been observed for gas phase synthesis. Kortshagen's group was able to synthesize exclusively cubic SiNPs in a nonthermal plasma generated by radiofrequency excitation.25,26 The silicon cubes originated from the restructuring of amorphous, caulif lower-shaped particles in the plasma zone due to surface diffusion.26 RF plasma reactors are able to produce a significant amount of atomic hydrogen that leads to an effective hydrogen passivation of the SiNP surface.28,46 Barnard & Zapol27 theoretically investigated the equilibrium shape for hydrogenated as well as for bare SiNPs in a size range of 1.6−7.3 nm by taking the Gibbs free energies for bulk, surface, edges, and corners into account. On the basis of the modeling, the cube appeared to be the energetically favored particle shape for a hydrogenated silicon surface which is consistent with the experimental findings of Kortshagen et al.26 It is pointed out here that Barnard’s model predicts equilibrium shapes for purely thermodynamic stability arguments. The thermodynamically most stable shape is not necessarily obtained in kinetically dominated processes (such as gas phase synthesis), which often result in nonequilibrium structures. Since we use a very different reactor system for particle synthesis and different mechanisms are present, the above-mentioned explanation for the formation of faceted particles may not be applicable for our system. In contrast to the RF plasma reactor, a thermal reactor does not produce fully hydrogen covered SiNPs.47 It is wellknown that elevated temperatures lead to hydrogen desorption

from the SiNP surface.48 Nevertheless, a partly hydrogenated surface remains as infrared spectroscopy revealed the presence of residual Si−Hx groups.20 It is difficult to predict the influence of partially hydrogenated surfaces on the shape (results in27 only for bare and fully hydrogenated SiNP). Thus, we attribute only minor influence to the hydrogen on the surface for the evolution of the octahedral shape. Murthy et al. studied the onset of nucleation (undesired particle formation for exceedance of critical temperature or critical silane concentration) for the optimization of epitaxial growth of silicon layers in an atmospheric pressure CVD reactor.24 Small amounts of octahedral (similar to our slightly truncated octahedra), tetrahedral, and truncated triangular bipyramidal particles were observed at a temperature of 1100 °C and a silane partial pressure of 2 mbar at hydrogen dilution. Murthy et al. ascribe the faceted shapes to a diffusion limited growth regime. Our results clearly demonstrate the feasibility of engineering the crystal shape from spherical to octahedral. Following Murthy et al.,24 we suppose that the transition from spherical SiNP (Figure 1a) to faceted SiNP (Figue 1b) relies on a change in growth conditions from a reaction limited regime to a diffusion limited regime. Both shapes were obtained for the same temperature (Tfurnace = 1100 °C) which consequently results in a similar surface diffusion rate for the crystallization units deposited on the particles. This leads to the conclusion that the difference in the growth mode has to be associated with the deposition rate of the crystallization units. Therefore, the formation of octahedra close to the wall is explained as follows: The hot wall of the reactor tube acts as an additional sink for the precursor and thus leads to a depletion zone for silane and crystallization units, respectively, close to the wall as it is quantitatively discussed in section 4.3. The evolution of a distinct depletion zone is ascribed to the higher total pressure (ptotal = 100 mbar) which results in a decreased diffusivity of the gaseous species. This means the deposited mass flux on the particles in the depletion zone is strongly reduced in contrast to the case of the production of the spheres. The temperature and thus the mobility of the crystallization units on the particle surface are sufficiently high to allow incorporation into the crystal lattice at energetically favorable sites. A model depicting the situation for dominating precursor deposition (spherical particles) and dominating surface diffusion (octahedral particles) in the depletion zone is shown in Figure 6. Interestingly, a variation of the initial silane partial pressure (other process parameters conform to the case of octahedra formation) did not alter the general situation: For increased as well as for reduced silane partial pressure octahedral SiNPs were obtained close to the wall (see Figure 8). This observation confirms the particular importance of the depletion zone. Another important prerequisite for the growth of faceted particles is ascribed to a proper axial reaction zone in the tube. It was observed that the reaction zone in the tube strongly differs in length and in the deposited silicon mass. As shown in Figure 7 the quartz tubes clearly show where silane is present 1334



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Figure 6. Model for the cases of dominating precursor deposition (formation of spheres) and dominating surface diffusion (formation of octahedra).

Figure 8. Growth regimes for octahedral/unfaceted (spherical or undefined) SiNPs with respect to calculated wall deposition rate/silane concentration close to the tube wall. The given numbers correspond to ptotal/mbar; Tfurnace/°C; psilane/mbar; τR/ms.

from the literature,49 it is possible to estimate mean silane concentrations close to the wall for different experiments. Although the kinetic data have to be extrapolated in terms of temperature, this will not lead to an essential error because the temperature is the same for every considered case except one. The upper limit of the deposition rate for octahedron formation (0.4 μm/min) corresponds to an average silane concentration of 1.4 mmol/m3. The lower limit of the other cases (0.7 μm/min) equals 2.6 mmol/m3.

Figure 7. Furnace temperature profile for 1100 °C; (a) quartz tube after experiment with sphere formation (ptotal = 15 mbar, Tfurnace = 1100 °C, psilane = 0.5 mbar, τR = 80 ms); (b) quartz tube after experiment with octahedron formation (ptotal = 100 mbar, Tfurnace = 1100 °C, psilane = 1 mbar, τR = 80 ms).

by the deposition of a polycrystalline silicon layer. Furthermore, the thermophoretically deposited SiNPs can be seen at the right end of the tubes. For the production of spheres (ptotal = 15 mbar) the silane is completely consumed after the first half of the tube (Figure 7 a). The tube for process conditions, at which octahedra were formed (ptotal = 100 mbar), is depicted in Figure 7b and shows a much more extended reaction zone. It turned out that it is impossible to obtain octahedra at 100 mbar with increased residence time where silane is also completely consumed in the first half of the tube (not shown). For a longer residence time, the growth of the particles is completed in the temperature gradient in the beginning of the furnace where the temperature is not high enough for a sufficient mobility of the crystallization units on the particle surface. Particles without a defined shape were obtained under such process conditions (for further details see Figure 8). 4.3. Separation of Growth Regimes for Different Particle Shapes. The last section aims toward a more quantitative description of the specific growth conditions and gives a simple approach for the separation of the growth regimes. Evaluating the deposited silicon mass and reaction length of the process tubes for various experiments (ptotal, Tfurnace, psilane, τR), we calculated an averaged silicon deposition rate on the wall for each experiment. A good correlation between the average deposition rate and the finding of octahedral-shaped particles could be deduced. Figure 8 clearly shows that octahedra only formed for wall deposition rates smaller than 0.4 μm/min. Spheres (or undefined particles) were observed for deposition rates larger than 0.7 μm/min. These considerations strongly support the qualitative statements in section 4.2. Using the calculated silicon deposition rates in combination with CVD data

5. CONCLUSIONS This work demonstrates that thermal pyrolysis of silane in a hot wall reactor is capable for crystal shape engineering of spherical as well as faceted SiNPs. Faceted nanoparticles can assist the design of densely ordered superstructures,50 which are a prerequisite for improvement of functional particle layers for application as, for example, thin film transistor. Differently shaped SiNPs were synthesized by carefully adjusting the global process parameters and thus the growth conditions. Single crystalline, defect-free, spherical particles with an average diameter of 30 nm and a GSD of 1.05 can be obtained in a reaction limited growth regime exemplarily represented by the following set of parameters: ptotal = 15 mbar, Tfurnace = 1100 °C, psilane = 0.5 mbar, mean residence time = 80 ms. The occurrence of faceted SiNPs with octahedral shape and a mean edge length of about 100 nm was observed for the following set of parameters: ptotal = 100 mbar, Tfurnace = 1100 °C, psilane = 1 mbar, mean residence time = 80 ms. It was found that the octahedra are formed close to the reactor wall since the wall acts as an additional sink for the crystallization units. Therefore a diffusion limited growth regime evolves which seems to be essential for the growth of faceted SiNPs.

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AUTHOR INFORMATION

Corresponding Author

*Phone: +49-9131-8529401; fax: +49-9131-8529402; e-mail: [email protected].

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ACKNOWLEDGMENTS The authors gratefully acknowledge the funding of the Deutsche Forschungsgemeinschaft (DFG) through the Cluster 1335



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of Excellence “Engineering of Advanced Materials” at the Friedrich-Alexander University Erlangen-Nuremberg. We thank Florian Niekiel for carrying out the HRTEM simulation shown in Figure 4b. Helpful discussions with Prof. L. Frey, Prof. H. Ryssel, and M. P. M. Jank from the Fraunhofer IISB Erlangen, and Prof. H.-J. Schmid from the University of Paderborn are gratefully acknowledged.

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