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J. Phys. Chem. 1996, 100, 3144-3149
Spatially Confined Chemistry: Fabrication of Ge Quantum Dot Arrays J. R. Heath,* R. S. Williams, and J. J. Shiang UCLA, Department of Chemistry and Biochemistry, 405 Hilgard AVenue, Los Angeles, California 90095-1569
S. J. Wind,* J. Chu, C. D’Emic, W. Chen, C. L. Stanis, and J. J. Bucchignano IBM T.J. Watson Research Center, P.O. Box 218, Yorktown Heights, New York 10598 ReceiVed: July 6, 1995; In Final Form: NoVember 6, 1994X
We report a technique for investigating nucleation and growth confined to nanometer scale surfaces. Lithographic and etching processes were used to create arrays of 100 and 150 nm holes through a thin SiO2 layer onto Si(100). Ge dots were nucleated and grown to a few nanometers in diameter within the patterned wells. Transmission electron and atomic force microscopic analyses revealed the presence of 0-1 Ge quantum dots in each of the 100 nm wells and 2-4 dots in the 150 nm wells. For the latter case, size-distance correlations indicated the effective radius of the diffusion field around a growing Ge particle was much larger than for growth on an infinite surface.
Introduction Electron beam fabrication and processing technology in the last decade has led to the creation of structures with precise dimensional control to less than 50 nm. This scale is not sufficiently small either for the fabrication of true size-confined quantum structures or for the fabrication of quantum effect devices for room temperature operation. In both cases, the required structure is often well below 20 nm in size. Semiconductor quantum structures of size 2-20 nm have been realized by a variety of wet1 and gas-phase2 chemical techniques. Some of these methods are capable of spectacular size control (+/-0.2 nm),3 and Langmuir-Blodgett techniques have even been utilized to produce quantum dot arrays.3,4 However, these quantum dots are only in proximal contact with the underlying substrate, and precise positioning of the dots or dot arrays is difficult. Leonard et al. have demonstrated uniform particle size control for 30 nm InAs islands directly grown on unpatterned GaAs substrates.5 These islands are epitaxially oriented with the underlying substrate, although their location on the substrate is, again, not well controlled. In general, randomly dispersed particles may find applications as optical filters or detectors. However, single-particle device-related tasks have wiring, geometry, and insulation requirements which mandate intimate electrical contacts and precise positioning of the dots. Lithographic resolution limits, while larger than the length scales for quantum confinement, are smaller than the diffusion length scales which govern a surface chemical reaction. Thus, by utilizing lithography to define templates for confining a heteroepitaxial chemical vapor deposition (CVD) process, it should be possible to investigate the effects of spatial confinement on surface chemistry. Such chemistry may provide a route toward the fabrication of positionally ordered and electrically contacted quantum dot structures of a size substantially smaller than lithographic resolution. In this paper, we report a process for the fabrication of ∼6-30 nm single-crystal Ge quantum dot arrays in which the size, crystal orientation, and spatial position of the dots are all controlled. Lithographic and etching processes were utilized to expose arrays of 100 and 150 nm diameter circular areas of Si through a SiO2 thin film. These * Authors to whom correspondence should be addressed. X Abstract published in AdVance ACS Abstracts, January 15, 1996.
0022-3654/96/20100-3144$12.00/0
exposed regions, bounded on the circumference by walls of SiO2, served as ‘nanosurfaces’ for investingating spatially confined nucleation and growth of Ge islands. Ultrahigh vacuum (UHV) CVD of GeH4 was used to selectiVely grow Ge quantum dots within these templates to a size of 6-25 nm. Between 2 and 4 quantum dots grew in each of the 150 nm wells, largely at defects (caused by an over-etch) around the internal perimeter of the wells. Transmission electron microscopy (TEM) measurements of nearest neighbor distances, coupled with island sizes, indicated that these islands act as sinks, creating diffusion fields that extend to the boundaries of the well. The 100 nm wells, which were substantially smaller than this diffusion field, were not over-etched. A high-resolution ‘supertip’ atomic force microscopy (AFM) analysis of these wells revealed the presence of a single Ge quantum dot in three out of four wells imaged. Process Methodology This fabrication technique used in this work to grow Ge nanocrystal arrays is described in Figure 1. The substrates were 〈100〉 silicon wafers which were cleaned and thermally oxidized to a thickness of ∼20 nm. The wafers were then coated with 70 nm of PMMA (poly(methylmethacrylate) and baked at 80 °C for 1 h. Arrays of holes were patterned in the PMMA using an IBM Vector Scan electron beam lithography system operating at 25 KeV. The arrays consisted of point exposures with pitches of 200 or 400 nm, covering areas from 1 mm2 to 1 cm2. The size of the holes was determined by the exposure dose, which was controlled by the dwell time of the beam on each hole site. Doses ranged between 50 and 75 µC/cm2 and, following development in a mixture of isopropyl alcohol and methyl isobutyl ketone (3:1), resulted in holes in the PMMA with approximate diameters of 60 and 100 nm. The thin oxide layer exposed by the dot pattern was then reactive ion etched (RIE) in CF4/CHF3 to expose the underlying Si substrate. There was a partial loss of anisotropy in the oxide etch, causing an enlarging of the hole diameters to 100 and 150 nm. An effect known as ‘RIE lag’ leads to differential etch rates, with smaller structures etching more slowly. No end-point detection was used in this etch, and the oxide is over-etched for the 150 nm wells by about 10 nm so that these wells were ∼30 nm deep, as measured by AFM. However, the 100 nm holes were etched © 1996 American Chemical Society
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J. Phys. Chem., Vol. 100, No. 8, 1996 3145 temperature of 600 °C for 5.0 min. Following the CVD step, the wafer was diced into several pieces for subsequent analysis. Analysis of the Ge Quantum Dots
Figure 1. Cartoon description of the process steps utilized to fabricate the nanosurfaces for the investigation of nucleation and growth in confined dimensions. In step 1, e-beam lithography defines arrays of ∼60-150 nm wells in photoresist. This is followed by an RIE step (2) to transfer the lithographic image through the SiO2 layer, exposing an array of nucleation sites on the Si substrate. The resist is removed, and Ge quantum dots are nucleated and grown on the exposed Si surface sites using UHV/CVD techniques.
The particle arrays were investigated using spectroscopic and microscopic techniques. The crystallinity of the Ge islands was confirmed in three ways. The 301 cm-1 Ge LO phonon was observed by micro-Raman spectroscopy. This mode was observed only on the patterned portion of the substrate. In addition, Moire´ fringes were imaged using plan view TEM (JEOL 4000 operating at 400 kV), indicating that the particles not only were crystalline but also were aligned with the substrate. For the larger particles, discontinuities were visible in the Moire´ fringes, indicating the presence of defects. As a final check, the characteristic Ge surface state absorption peak at 5890 cm-1 was identified via near-IR absorption spectroscopy. The individual wells were imaged using AFM and TEM. The AFM data were collected with a Topometrix TMX 2010 MultiView atomic force microscope in contact mode with a high aspect ratio ‘supertip.’ For ‘supertip’ AFM, a carbon whisker is grown onto the end of a standard silicon nitride AFM tip to produce a narrow probe capable of mapping crevices and valleys on a substrate with minimal loss in resolution. Particle size measurements and nearest neighbor distances (for the 150 nm diameter wells) were taken from relatively low resolution TEM data. Well-to-well centers were taken as the length standard for these measurements. Results and Discussion
Figure 2. ‘Supertip’ AFM image of an array of 150 nm wells produced according to the scheme outlined in Figure 1. The massively parallel nature of this experiment is critical for obtaining reliable statistics on the kinetics of nucleation and growth in confined areas.
to very nearly the exact thickness of the oxide and were measured by AFM to be ∼17 nm deep. The PMMA was then removed via an acetone rinse. A low-resolution AFM image of a 25 µm2 section of the wafer patterned with with 100 nm diameter wells is presented in Figure 2. Both the 100 and 150 nm diameter arrays were patterned onto the same substrate. The wafer was briefly dipped in a 10% HF solution to remove any native oxide from the exposed silicon sites and then introduced into the load lock of a UHV/CVD growth chamber with a base pressure below 5 × 10-6 Torr. The wafer was then exposed to 10% GeH4 in He at a flow rate of 90 sccm, a deposition pressure of 1-2 mTorr, and a growth
In order to understand and control Ge quantum dot growth in confined dimensions, it is important to appreciate the physical chemistry of the relevant CVD process. Heteroepitaxy, according to Bauer,6 is characterized by three growth modes: Frank-van der Merwe (F-vdM), or 2D layer by layer growth; Volmer-Weber (V-W), or 3D (island) growth; and, StranskiKrastanov (S-K), which is initial 2D growth followed by 3D growth. These modes are rationalized from thermodynamic considerations involving the surface energies of the substrate and epilayer, the interfacial energy, and the lattice-mismatch related strain energy. The surface free energy of Ge(100) is calculated to be about 30% less than that of Si(100),7 a condition which favors F-vdM growth. However, the 4% lattice mismatch between Ge and Si leads to an increasing strain energy with film thickness, favoring S-K growth. For UHV/CVD, Ge/Si heteroepitaxy proceeds smoothly from a F-vdM growth mode below T0 ) 350 °C to a S-K and/or V-W mode at higher growth temperatures.8 The growth conditions for these arrays were determined by growing a series of Ge films on planar unpatterned Si(100) for various times and temperatures, followed by TEM analysis of cross-sectioned wafers. It was found that growth in the 560600 °C range led to the most uniform distribution of Ge islands on the Si surface. In addition, an incubation period of just under 5 min of wafer exposure to GeH4 near 600 °C was necessary to initiate any island growth. After 8 or 9 min of exposure, the Ge particles were already quite large (∼0.1 µm) and the substrate coverage was nearly complete (70-90%). Thus, the 5.0 min of exposure used here represents only the earliest stages of real growth. For an infinite surface, individual islands are all coupled to each other through their diffusion fields. The present experiment may be considered as a large number of isolated nucleation and growth events proceeding in parallel, with the SiO2 barriers between the Si nanostructures effectively decoupling the nucleation and growth processes among the various wells.
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General experimental conditions, such as the growth temperature, growth time, and adatom flux, are invariant from well to well. However, the number and nature of the sites available for heterogeneous nucleation may change from well to well. For this reason, measurements such as overall particle size distributions will be inhomogeneously broadened and not particularly informative. However, the nature of the experiment does yield rich information relevant both to the earliest stages of growth within a well and to the diffusion field around a growing island. Within this diffusion field, the concentration of adatoms is strongly affected by the presence of a growing island, which acts as a ‘sink’ for adatoms. The end result is a decreased rate for subsequent nucleation events within the field. The kinetics of 3-D island growth on a 2-D surface are well described by an activated nucleation step followed by diffusioncontrolled growth. In the present case, nucleation is likely to occur at defects and is thus a heterogeneous process. The size distribution at the earliest stages of growth (the period most relevant to this work) is broadened by the nucleation of new islands concurrent with the continued growth of older islands. For such kinetics, the rate of particle growth (after nucleation) may be expressed as
Rn - R0n ) k(t - t0)
(1)
Here, R is the size of the growing island, R0 is the critical island size for growth, and t - t0 is the time elapsed after the critical island size has been reached. The dependence of the exponent (n) in the growth law has been discussed by Ngo and Williams for the cases of n ) 2, 3, and 4.9 The n ) 3 case, which has been observed experimentally,10 is the most likely description of growth for the present experiment, as it implies a diffusion field around a growing 3-D particle. The diffusion field length increases with increasing island size, such that the (diffusion field size)/(island size) relationship maintains a constant aspect ratio and growth is dominated by diffusion kinetics.11 For all cases, certain observations are expected. Once the nucleation sites are saturated, smaller particles grow more quickly than large ones, so that the size distribution narrows with increasing growth time.12 Indeed, for islands grown on unpatterned Si substrates, we observe island size distributions to narrow substantially with increasing island size. An implication from this argument (independent of the value of the exponent in eq 1) is that the size of an island within a well, normalized to other islands within that same well, is a relative measurement of when that island nucleated and began to grow. Thus, it is possible to identify islands which represent the first, second, third, etc. nucleation events within a well by their relative size. A measurement of the distance between the first and second nucleation events (dn1fn2) within a well should give an estimate for the spatial influence a growing island exerts on a subsequent nucleation event. For the rest of this paper, two implicit assumptions are made in modeling the data. The first assumption is that the kinetics of island growth do not change as the islands grow. One might expect that an activation barrier toward forming a defect could prevent highly strained islands from growing.13 However, the temperature (600 °C) chosen for these experiments is relatively high for Ge CVD and should be enough to overcome any such activation barriers so that the growth is kinetically, rather than thermodynamically, limited.14 The second assumption is that GeH4 falling on the oxide does not contribute to the growth of islands within the wells. 600 °C is more than sufficient to grow Ge directly on SiO2. However, the weak interaction between SiO2 and GeH4 requires relatively high pressures of GeH4 to build up a sufficient
Figure 3. TEM dark-field image of an array of 150 nm wells. In dark-field, the contrast between the oxide overlayer and the Si wells is not very high, although the contrast between the Ge islands and the Si substrate is amplified. Moire´ fringes are observed, indicating epitaxial registry between the Ge islands and the Si. Note that the vast majority of the Ge islands have nucleated around the perimeter of the wells due to the ring of defects produced by the RIE over-etch.
monomer surface coverage to effect nucleation. Below 0.1 Torr, Ge CVD on Si/SiO2 is selective for Si.15 The GeH4 partial pressures employed here were a factor of 1000 below this limit, in an effort to intentionally keep the monomer coverage on SiO2 low. The end result is that the chemical events occurring within each of the wells were decoupled by the SiO2 barriers separating the wells. 150 nm Wells A TEM ‘dark-field’ plan view micrograph of a set of 150 nm wells containing Ge particles is shown in Figure 3. The ‘dark-field’ technique increases image contrast between the two crystalline materials present, the Ge islands and the underlying Si substrate. The circular Si/SiO2 well boundaries, which show up well in bright field, are difficult to discern here, since the SiO2 is amorphous. Note that the majority of the particles have nucleated and grown at the perimeter of the wells, consistent with the presence of radial defects introduced by the RIE overetch. The presence of Moire´ fringes indicates epitaxial alignment between the Ge dots and the underlying Si substrate. For some of the larger particles, discontinuities in these fringes indicate the presence of misfit dislocations. Island sizes and island-island distances were tabulated from images similar to Figure 3. The separation between each island was measured from the island centers. This, coupled with the minimum island sizes involved in the measurements, limits the possible values of dn1fn2 to between 15 and 135 nm. For most wells, the maximum possible separation was actually closer to 125 nm. As indicated above, the diffusion field around a growing island should scale with island size. An actual estimate of this size dependence involves a knowledge of the rate constant for growth. However, a plot of dn1fn2 vs r2/r1 should somewhat reflect this dependence, and such a plot is given in Figure 4. Here, r2 and r1 are the sizes of the second and first islands,
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C(r) ) Cb[1 - Ko(r/ζ)/Ko(a/ζ)]
Figure 4. Spatial influence that a growing particle’s diffusion field exerts on subsequent nucleation events. The size of the second island to nucleate within a given well is normalized against the size of the first island. This result is plotted vs the separation of the first and second islands. The distances between the two islands were measured from the centers of the islands, and so the largest separation possible is 125-135 nm, depending on island sizes. Note that the distance between the two islands is uncorrelated with r2/r1, indicating that, for even for the smallest islands, the diffusion field effectively extends to the farthest boundaries of the well.
respectively. Data from 50 separate wells were utilized in Figure 4. Only three nearest neighbor distances were measured to be closer than 100 nm. Figure 4 highlights two important characteristics of particle growth within these confined areas: the size of the diffusion field around a growing island is not related to the island size, and the upper limit of the radius of this diffusion field is set by the diameter of these wellssi.e. 150 nm. These measurements indicate that nucleation and growth kinetics within wells larger than 150 nm would reflect confinement effects similar to those observed here. It is difficult to estimate, however, the size at which such effects would cease to be important. To put this result into perspective, consider the case of a growing particle in an unconfined area. The size of a diffusion field around a growing particle is related to the screening length ζ ) (D/k)1/2, where D is the diffusion coefficient and k is the reaction rate constant for island growth. The radial concentration dependence of monomer C(r) around a growing particle may be calculated by solving a modified nonhomogeneous diffusion equation:16
∂C/∂t - D∇2C ) S - kC
(2)
Here, S represents the incident monomer flux times the sticking coefficient (in no. of monomer s-1 cm-2)) C is the (radially dependent) concentration of monomer centered on a growing island, and kC represents the averaged sink of all other islands on the surface. Assuming the steady state approximation (∂C/ ∂t ) 0) and applying the boundary values that C ) 0 at the radius of an island (r)a) and C ) Cb (bulk concentration) at large r, eq 2 may be directly solved (for the 2-D case) to give the following expression for C(r):
(3)
Ko is a zeroth order modified Bessel function of the second kind, which decays exponentially at large r/ζ. Typically, the diffusion field extends out to a length of about 2-4 island radii from the perimeter of an island. The actual diffusion field length for the islands in the wells is probably not too different from this. However, the size of the effectiVe diffusion field is apparently amplified by reflection of monomer at the boundaries of the wells. At 600 °C, Ge adatoms on a Si surface can diffuse a distance of several microns.17 Reflection of monomer at the well boundaries implies that an adatom has several chances to pass through and be captured by a single island. If the walls were lossy, rather than reflective, the diffusion field length of an island would probably be observed to be much shorter. A second piece of information that can be gained from this data relates to the nature of the nucleation kinetics. The ∼5 min incubation period at the onset of the UHV/CVD process may imply that island growth requires a critical surface coverage of monomer for the nucleation of islands. This incubation period is quite long, however, and other processes, such as a slow etch of the well-bases by GeH4 to remove any remaining SiO2, may be responsible for all or part of this incubation period. The nucleation events clearly take place at defects, although the role that these defects play in assisting nucleation is not obvious. One possibility is that the defects fit the more extreme definition of heterogeneous nucleation sites; i.e., the defects completely remove the activation barrier for nucleation. Once such sites become available, island formation at those sites is thermodynamically favorable, limited only by diffusion kinetics. A more relaxed definition of a heterogeneous nucleation site is one that simply lowers but does not completely remove the activation barrier for nucleation. The rate of island formation at such sites will initially be limited by statistical fluctuations. This is similar to a purely homogeneous nucleation process, although the size of the critical nucleus is effectively reduced at the defect site. For the purposes of this paper, we will refer to this case as ‘assisted homogeneous nucleation’. These two processes may be differentiated from one another by looking at the temporal evolution of initial nucleation events within each well. Consider the case of pure heterogeneous nucleation. This mechanism does not require an adatom concentration to build up on the surface prior to island formation, and thus the critical nucleus contains only a single atom. Once the defect sites are ready for growth (t ) 0), then there is no subsequent incubation time, and islands begin to grow immediately. As the defect sites become saturated, the island formation rate levels off and then decreases exponentially. The rate of island formation near t ) 0 is thus a convex-up function. For the case of assisted homogeneous nucleation, the rate of island formation is quite different. Near t ) 0 this rate is limited by statistical fluctuations which describe the formation of a critical nucleus at a defect site. As the adatom concentration builds up, these statistical fluctuations become more frequent. Thus, the rate of island formation is described by a concave-up function near t ) 0, with an incubation period determined by the time-dependent adatom concentration and the size of the critical nucleus. Similar to the heterogeneous case, at longer times the nucleation rate also flattens out and exponentially decays as the sites become saturated. For an infinite surface which obeys the growth kinetics of eq 1 with the exponent n ) 3, the rate of island formation can be extracted by plotting the number of islands of radius R vs R3. Such a plot is inappropriate for this experiment, since island formation in each of the wells is uncoupled from every other well. However, the isolated wells
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Figure 5. Time dependence of the first nucleation event within each well. A growth law of dr/dt ) kr-2 is assumed, so that the cube of the radius of an island is proportional to the elapsed growth time for that island. The sizes of the islands are measured with ∼5 nm resolution, and t ) 0 is assigned to an island which is just a few nanometers larger than the largest island actually measured. The initial nucleation rate is concave-up, indicative of an ‘assisted homogeneous’ nucleation process as described in the text.
each have a first nucleation event. The probability of a first nucleation event, by definition, is not related to the diffusion fields of other islands. Thus, a plot of the number of islands of radius R vs Ri3, where Ri represents the size of the largest island in the ith well, should reflect the behavior of the initial rate of island formation. Such a plot is presented in Figure 5. For this plot, island radii were measured with an approximately 5 nm resolution. RL, a hypothetical island formed at t ) 0, was arbitrarily set to an island size a few nanometers larger than the largest observed island (40 nm). This initial rate behavior is the critical signature which separates the heterogeneous and assisted homogeneous mechanisms. This plot clearly evidences concave-upward behavior at early times and shows that Ge island growth within these wells proceeds via an assisted homogeneous mechanism. To illustrate this point, a nonlinear curve has been fit to the rising edge of the data. In terms of the previously described growth modes, this data indicate that, even for defect-induced growth, a Stranski-Krastanov rather than a Volmer-Weber description still applies. 100 nm Wells As discussed above, 150 nm wells were already sufficiently small to affect any nucleation and growth processes which occurred within the wells. Such effects are likely to be amplified in smaller dimensions. Also of importance for these experiments was the nature of the 100 nm wells, which are only about 17 nm deep. Recall that the initial oxide coverage was 20 nm. The wet etch prior to the CVD step is likely to remove some of this oxide. Thus, the 17 nm depth indicates that these smaller wells were etched to just above the Si surface, and the wet etch
Heath et al.
Figure 6. ‘Supertip’ AFM image of a single ∼100 nm well revealing the presence of a single Ge quantum dot which has nucleated and grown on the Si substrate. On the bottom is a height vs distance profile across the base of the well. The measured height of the dot (1.4 nm) is a low-limit measurement of its radius, and the measured width (33 nm) is an upper limit measurement of its diameter.
(which is specific to the oxide) subsequently exposed the Si bases. These wells then do not have a ring of defects around the perimeter, and initial island nucleation is likely to occur anywhere within the well. Unfortunately, the preparation of the wafer for plan view TEM imaging either destroyed or did not sufficiently thin the regions containing the 100 nm wells. However, it was possible to image four of these wells using supertip AFM. An image of one such well is shown in Figure 6. This image clearly shows the presence of a single structure at the center of the well. In addition, the walls of the wells are sharp, especially in comparison to those of the 150 nm wells. For those larger wells, similar measurements are affected by the presence of islands around the perimeter of the wells, which tend to flatten the imaged slope of the walls considerably. No such broadening was observed for any of the 100 nm wells. Of the four wells imaged, three of the images revealed a morphology similar to that of Figure 6, with an island at the center of the well. One of the wells did not appear to contain any islands. As discussed above, supertip AFM is a technique capable of imaging narrow valleys and crevices on a substrate. However, the scanned image is nonetheless a convolution of the probe tip geometry and the sample being scanned. Thus, the image of a particle sitting in a small diameter well is somewhat larger in width and shorter in height than the particle itself. TEM side-view images of similar Ge islands on Si reveal a hemispherical (although faceted) morphology, such that the height of the island is approximately equal to the radius. This implies that twice the AFM measured height gives a lower limit on the particle diameter, while the measured width gives an upper limit on the particle diameter (measured here to be between 3.0 and 34 nm). A detailed scan of a second well was taken, and a single dot of diameter between 3.0 and 35 nm was found. The
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lack (or greatly reduced density) of defects in these smaller wells could lead to a very different class of nucleation and growth kinetics and may be responsible for the presence of just a single island in most of these wells. However, with just AFM data, it is difficult to gain much detailed insight into these issues.
Acknowledgment. We thank Dr. Gary Williams of Topometrix for collecting some of the AFM images and Jeff Kash and Cate Chess for assistance in Raman and near-IR analysis. J.R.H. acknowledges the support of IBM, the NSF NYI Program, and a Packard Fellowship.
Summary and Conclusions
References and Notes
The chemistry of island nucleation and growth kinetics within confined areas has been investigated in a massively parallel fashion. The selective growth of Ge islands on Si was the chemical system studied, and the confined areas were 150 and 100 nm diameter Si nanosurfaces bounded by SiO2 walls. These areas were defined to be smaller than the diffusion length scales which govern the relevant kinetics. For 150 nm wells, profound effects were observed. The size of the diffusion field around a growing particle is substantially increased over that expected for growth on an infinite surface, presumably due to reflection of monomer at the boundaries of the wells. The parallel nature of this experiment allowed for a measurement of the initial rate of island formation, and this rate was shown to be dominated by ‘assisted homogeneous’ nucleation, in which the nucleation sites are defects around the perimeter of the wells. The 100 nm wells, which did not have such defects, were analyzed by supertip AFM. Three out of four wells imaged were shown to have a single Ge island of size 3-30 nm located at the center of the well. This result indicates that lithographic access of chemical length scales, rather than structural length scales, can provide a route for the fabrication of precisely positioned, electrically contacted, and size-confined quantum dots. This observation should have implications for the fabrication of Coulomb blockade devices designed for ambient operation.
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