J. Phys. Chem. C 2007, 111, 779-786
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Probing Interactions of Ge with Chemical and Thermal SiO2 to Understand Selective Growth of Ge on Si during Molecular Beam Epitaxy Qiming Li,† Joshua L. Krauss,‡ Stephen Hersee,§ and Sang M. Han*,† Chemical and Nuclear Engineering Department, UniVersity of New Mexico, Albuquerque, New Mexico 87131, Materials Science and Engineering Department, UniVersity of Wisconsin, Madison, Wisconsin 53706, and Center for High Technology Materials, Electrical Engineering and Computer Engineering Department, UniVersity of New Mexico, Albuquerque, New Mexico 87131 ReceiVed: May 15, 2006; In Final Form: September 20, 2006
We have previously demonstrated that Ge selectively grows on Si over a SiO2 mask during molecular beam epitaxy. To determine the surface phenomena responsible for the selectivity, we probed the stability of SiO2 upon Ge exposure and Ge diffusion through thin SiO2, using X-ray photoelectron spectroscopy, ellipsometry, atomic force microscopy, transmission electron microscopy, and multiple internal reflection Fourier transform infrared spectroscopy. We observe that the consumption of SiO2 occurs only with chemical oxide upon Ge exposure via Ge + Si + 2SiO2 f GeO(g) + 3SiO(g) at the SiO2/Si interface where all three reactants are present. This erosion is initiated by the Ge diffusion only through thin chemical SiO2, and the diffusion is attributed to a larger concentration of SiOH groups (∼1 × 1022 cm-3) and a greater porosity in chemical oxide than that in thermal oxide. For thermal SiO2, where Ge diffusion and subsequent oxide degradation are not observed, we have determined that the selectivity stems from the low desorption activation energy (Edes) of Ge adspecies from the thermal SiO2 surface. The experimentally measured Edes is 42 ( 3 kJ/mol on the order of Van der Waals force. The low Edes entails a low activation barrier (∼13 kJ/mol) for the surface diffusion of Ge adspecies on thermal SiO2, leading to a characteristic diffusion length greater than 1 µm. Under typical Ge growth conditions where the interdistance between exposed Si areas is much less than 1 µm, the large diffusion length would cause Ge adspecies to migrate over SiO2 and preferentially aggregate on the exposed Si surface.
1. Introduction Selective epitaxial growth (SEG) occurs when deposited material readily nucleates on the desired substrate surface, while nucleation is completely suppressed on the mask. The SEG technique provides precise dimensional and positional control over epitaxial structures by utilizing a masking template created by lithography.1-7 Such control is well suited for improving the epilayer quality, when the artificially engineered template simultaneously prevents dislocations from propagating and decreases the interfacial strain.8,9 For instance, the epitaxial lateral overgrowth (ELO) technique9-14 relies on high-aspectratio windows in a patterned mask. The epitaxial islands grow in these windows, eventually surpassing the window openings, and coalesce laterally over the masking material to form an epilayer. The ELO technique, when applied to the Ge/Si bimaterial system, improves the optical and electrical properties of the epilayer by lowering the threading dislocation density.5,8,9,15 The necking of threading dislocations that propagate along directions is largely responsible for the improvement. Another viable SEG technique for the Ge/Si system is the “touchdown” process,5,8,15,16 where the Ge beam opens nanoscale windows in a 1.2-nm-thick chemical SiO2 layer and forms Ge seeds that coalesce over the remaining SiO2. The * Corresponding author. E-mail:
[email protected]. † Chemical and Nuclear Engineering Department, University of New Mexico. ‡ University of Wisconsin. § Center for High Technology Materials, Electrical Engineering and Computer Engineering Department, University of New Mexico.
touchdown process is uniquely different from the previously established ELO technique in that the characteristic dimension (3-8 nm) of the Ge-Si junction pads ensures the growth of coherent Ge seeds;16 therefore, the high-aspect-ratio windows are not necessary for the touchdown process. The remaining amorphous masking material also serves as artificially introduced 60° dislocations that relieve the strain near the Ge-Si heterojunction.16 The resulting threading dislocation density is below 105 cm-2. Herein, we focus on the surface phenomena responsible for the random, yet self-limiting formation of nanoscale windows in thin chemical SiO2 films during touchdown. We also differentiate the consumption of chemical SiO2 upon Ge exposure from the physical mechanisms responsible for the selective growth of Ge on Si over SiO2. The success of SEG relies on a thorough understanding and manipulation of the surface phenomena occurring on the substrate versus the mask. For example, the SEG of SixGe1-x (0 e x < 1) has been demonstrated in chemical vapor deposition (CVD) systems, where a SiO2 mask is used together with selectivity-controlling agents such as H and Cl.17-22 The following surface phenomena are thought to be responsible for the selectivity in such CVD systems: (1) Cl etches Ge nuclei on SiO2.17,20,21 (2) H absorbed on SiO2 prevents Ge nucleation.18,19,22 (3) Ge reacts with SiO2 to form volatile SiO and GeO.18,23 In molecular beam epitaxy (MBE) systems, where the selectivity controlling agents are often absent, the SEG of SixGe1-x on Si over a SiO2 mask has also been demonstrated.5,24-26 On the basis of the experimental observation that Ge molecular beam exposure removes ultrathin SiO2 films,8,27-29
10.1021/jp062966o CCC: $37.00 © 2007 American Chemical Society Published on Web 12/07/2006
780 J. Phys. Chem. C, Vol. 111, No. 2, 2007 the selectivity has been attributed to the reaction Ge + SiO2 f GeO(g) + SiO(g), which does not require elemental Si.8,25,27-29 However, other findings contradict this mechanism at substrate temperatures less than 700 °C. First, oxygen transfer occurs from Ge to Si at elevated temperatures because the reaction enthalpy of oxygen with Si is larger than that of oxygen with Ge.27,30,31 Second, if the starting materials are GeO2 and Si, then the oxygen transfer from Ge to Si accelerates with increasing temperature, predominantly producing SiO2 rather than volatile GeO and SiO.31 Third, Ge promotes SiO2 degradation only in the presence of Si, and Ge alone does not degrade SiO2.26 Thus, it is unlikely that GeO and SiO are formed via reduction of SiO2 by Ge at a substrate temperature lower than 700 °C. We therefore focus on the interaction of Ge with both chemically and thermally grown SiO2 films to understand the surface phenomena responsible for the SEG of Ge on Si over SiO2 during MBE. The Ge interaction with SiO2 is studied using a variety of spectroscopic and materials characterization techniques, such as X-ray photoelectron spectroscopy (XPS), ellipsometry, atomic force microscopy (AFM), transmission electron microscopy (TEM), and multiple internal reflection Fourier transform infrared spectroscopy (MIR-FTIRS). These techniques helped determine the instability of chemical SiO2 upon Ge exposure, Ge diffusion through thin chemical SiO2, the Ge desorption rate from thermal SiO2, and stable Ge adspecies on thermal SiO2. 2. Experimental Section 2.1. SiO2 Growth and Erosion Under Ge Exposure. Ultrathin chemical SiO2 films are grown on Si by immersing Si wafers in a H2SO4/H2O2 solution. The H2SO4/H2O2 solution is prepared by mixing H2SO4 (2 M) with H2O2 (30 wt %) in a 4:1 volumetric ratio. Depending on the crystallographic orientation of Si and the composition of H2SO4/H2O2 solution, the thickness of the chemical oxide increases up to 1.0-1.4 nm with increasing immersion time.32,33 In our experiments, the Si(100) substrates are first immersed in the H2SO4/H2O2 solution for 5 min. The chemical oxide is subsequently removed in a HF solution (11 wt %). The H2SO4/H2O2 and HF treatments are repeated three times to remove organic contaminants from the Si surface. The Si substrate is finally treated in a fresh batch of H2SO4/H2O2 solution at 80 °C for 10 min to grow the chemical oxide. The thickness of the resulting chemical oxide is approximately 1.2 nm, measured by cross-sectional transmission electron microscopy (XTEM).8 After deionized H2O rinsing and N2 blow drying, the substrates are immediately loaded into an ultrahigh vacuum molecular beam epitaxy (UHV-MBE) chamber through a high-vacuum transfer load lock. The base pressure of the UHV-MBE chamber is maintained at 3 × 10-10 Torr. During the deposition, the chamber pressure increases to 1 × 109 Torr. A Ge effusion cell, operated at 1050 °C, provides a flux of 1.1 × 10-2 equivalent monolayers (1 equiv ML ) 6.3 × 1014 atoms/cm2) per second. The loaded substrate is heated to 620 °C for Ge exposure. We interrupt the Ge exposure intermittently to collect Si2p and Ge3d X-ray photoelectron spectra to quantify the erosion rate of chemical SiO2 upon Ge exposure. The takeoff angle is set to 20° to obtain a relatively strong photoelectron signal from the film surface, rather than from the underlying substrate. The takeoff angle is defined here as the angle between the sample surface and the orientation of the electron energy analyzer. For XPS, we use an Al KR excitation source and scan the Si2p and Ge3d regions with a 0.2 eV step size and 10 s dwell time. The thermal SiO2 films are grown on Si substrates to thicknesses that range from 6 to 200 nm. The dry oxidation is
Li et al. performed in an N2/O2 ambient at 1 atm and 1000 °C for a few seconds to 25 min. The mass flow rate for both O2 and N2 is 60 sccm per square-centimeter cross section of the quartz-tube furnace. The oxidized samples are exposed to Ge at a variety of fluxes (1.1 × 10-2 to 2.4 × 10-1 equiv ML/s) and substrate temperatures (300-800 °C). The thickness and surface roughness are compared before and after the Ge exposure in order to determine the existence and extent of SiO2 degradation. The SiO2 thickness is measured by single-wavelength ellipsometry, and the surface roughness is measured by atomic force microscopy (AFM). 2.2. Characterization of Ge Diffusion Through SiO2. Scanning transmission electron microscopy (STEM) is used to characterize the Ge diffusion through ultrathin chemical SiO2 and dry thermal SiO2 films. A 100-nm-thick thermal SiO2 layer is first grown on Si(100) at 1000 °C. An array of holes is created in the SiO2 layer by interferometric lithography34-36 and reactive ion etching to expose the underlying Si. The sample is then treated in a H2SO4/H2O2 solution to form an ultrathin layer of chemical oxide on the exposed Si surface. The sample is finally exposed to Ge at 1.0 × 10-1 equiv ML/s for 12 h. The substrate temperature is held at 700 °C during Ge exposure. The sample is then studied by STEM to map the Ge distribution at the interfaces between chemical oxide and Si as well as between thermal oxide and Si. 2.3. Quantitative Estimation of SiOH Concentration in SiO2. Multiple internal reflection Fourier transform infrared spectroscopy (MIR-FTIRS) is used to compare the quality of chemical oxide compared to that of thermal oxide. For the MIRFTIRS analysis, a chemical oxide film is grown on a Si(100) MIR crystal. The dimensions of the Si crystal are 5 cm long, 1 cm wide, and 2 mm thick with 45° bevels on the opposing edges of the longest dimension. The IR beam, upon entering the Si crystal from the first beveled edge, makes 25 internal reflections at the top and bottom surfaces before exiting the second beveled edge. The IR signal is then collected by a liquid-N2-cooled HgCdTe detector. The internal reflections create an evanescent electromagnetic field penetrating through the chemical oxide. This evanescent field interacts with the dipole moments of IRactive species (e.g., SiOH) and induces their vibrational excitation, which subsequently appears as absorbance in the IR spectra. A similar experimental setup is described in detail in our previous publication.37 Both background and sample spectra are taken with 2 cm-1 resolution. Each spectrum is averaged over 300 scans. 2.4. Measurement of Desorption Activation Energy by X-ray Photoelectron Spectroscopy. To estimate the desorption activation energy (Edes) of Ge adspecies from the SiO2 surface, we measure the Ge desorption rate from the thermal SiO2 surface as a function of substrate temperature (Ts). At a fixed Ge flux, we incrementally decrease Ts and intermittently analyze the growth surface until the Ge3d X-ray photoelectron signal is detected. For XPS analysis, the Ge impingement is first turned off by closing the effusion cell shutter. The substrate heater is then turned off, and the sample is transferred in vacuum to the XPS analysis chamber. The sample transfer typically takes less than 5 min, and the substrate is quenched to room temperature during this period. We assume that the Ge desorption rate (Rd) perfectly counterbalances the absolute Ge flux (FGe) at the transition point where the Ge3d signal appears. For instance, when the effusion cell is maintained at 1130 °C, and the corresponding absolute Ge flux is 6.0 × 10-2 equiv ML/s, this transition point occurs near 630-650 K. We trace similar temperature excursions for other effusion cell temperatures, plot
Probing Interactions of Ge
J. Phys. Chem. C, Vol. 111, No. 2, 2007 781 the attenuation is not considered in calculating the film thickness. The effective oxide thickness (δ), plotted as a function of Ge exposure in Figure 1, is approximated by simply comparing the integrated areas of XPS peaks according to
δ ) 1.2 nm ×
Figure 1. Si2p photoelectron spectra collected from the chemicaloxide-covered Si(100) surface with increasing Ge exposure. From spectrum A to spectrum H, the Ge exposure leads to the disappearance of the Si4+ peak, corresponding to the erosion of chemical SiO2. The dotted line (---) represents our model fit.
ln(Rd) versus 1/Ts, and calculate Edes from its slope. Although the steady-state surface coverage of Ge at the transition point might be finite, and the subsequent surface coverage may decrease to zero during sample transfer, we expect the steadystate coverage to be near the zero-coverage limit. As our later analysis will suggest, the characteristic diffusion length of the Ge adspecies is much greater than 1 µm within the explored temperature range, and any coalescence into an island with GeGe bond formation would be irreversible, leading to a detectable Ge3d signal. As a conservative estimate where the characteristic diffusion length is 1 µm, the steady-state Ge surface coverage would be on the order of 10-7 ML. Therefore, the steady-state Ge surface coverage is likely to be much lower than 10-7 ML, and we assume that the adsorption rate is equal to the impingement rate in the zero-coverage limit with a sticking probability of unity. 2.5. Critical Island Size and Stable Ge Adspecies on SiO2. High-resolution scanning electron microscopy (HR-SEM) (Hitachi S5200) is used to measure the saturation island density of Ge on thermal SiO2 as a function of substrate temperature and net Ge flux. The electron beam energy is adjusted to 1 keV in order to prevent charging on nonconductive materials such as SiO2, while eliminating the need to deposit a conductive material over the sample for imaging. The given electron beam energy provides 1.7 nm resolution. The substrate temperature ranges from 300 to 500 °C. The net flux is obtained by subtracting the desorption flux from the impingement flux. On the basis of the relationship between saturation island density and Ge absolute flux, we deduce the critical island size (i) using the mean field theory.38,39 3. Results and Discussion The XPS technique provides a means to quantitatively assess the stability of chemical and thermal SiO2 films upon Ge exposure based on the intensity of Si2p photoelectron emission from SiO2 versus Si and the presence of Ge3d photoelectron emission. The experimentally measured erosion rate of chemical oxide upon Ge exposure, in turn, provides a possible reaction mechanism for the oxide loss and its kinetics. The inset of Figure 1 shows a series of Si2p photoelectron spectra taken from the chemical-oxide-covered Si(100) surface with increasing Ge exposure. The peaks at 103 and 99 eV correspond to Si2p photoelectrons emitted from SiO2 and underlying Si, respectively. The Si2p photoelectron attenuation length in SiO2 is 3.448 nm from Al KR irradiation.40 This value is much larger than the thickness of the chemical oxide layer (1.2 nm); therefore,
I103eV I1.2nm 103eV
(1)
where I103eV and I1.2nm 103eV represent the integrated intensity of the Si2p peak at 103 eV after the Ge exposure and that for the 1.2nm-thick chemical SiO2 before the Ge exposure, respectively. This simple approximation provides an effective thickness as a measure of the remaining amount of chemical SiO2 during Ge exposure. According to the above approximation (Figure 1), the oxide thickness initially increases from 1.2 to 1.6 nm before the erosion takes over. Although abundantly available SiOH groups (see our later discussion) in the chemical oxide may release H2O and form Si-O-Si bridges, this reaction will not lead to a net increase in the Si4+ signal. We speculate that the porous chemical SiO2 contains Si dangling bonds41-43 and that these Si dangling bonds bind with O on Si-O and form Si-O-Si bridges upon Ge exposure. The excess Si-O formed would lead to an increase in the Si4+ signal. This speculation also entails that the Si2p signal from suboxides (i.e., SiOx where x < 2), whose binding energy peaks appear in the 100.51-102.14 eV range,44 must decrease during SiO2 formation. However, the peaks corresponding to Si3+ (102.14 eV) and Si2+ (101.34 eV), which are likely surface species, overlap with the shoulder of the Si4+ (103.34 eV) peak. Because of the limited resolution of our XPS measurements with a non-monochromated source, we were not able to reliably distinguish the increase in Si4+ from the decrease in suboxide peaks by deconvoluting the peaks near 103.34 eV. During the erosion of SiO2, the time-averaged oxide removal rate is approximately 0.5 Å/s. Note that the diminishing intensity of Si4+ photoelectron peak is not due to Ge nucleation on the SiO2 surface, which may potentially cause the attenuation of the Si4+ signal. The Ge nucleation on the thermal SiO2 surface is not detected by XPS under the given exposure condition.8 We will discuss the conditions that lead to Ge nucleation within the context of measuring the desorption activation energy of Ge adspecies from SiO2. Under the low Ge flux (1.1 × 10-2 equiv ML or 6.9 × 1012 atoms/cm2‚s) and at the substrate temperature of 620 °C, the erosion of chemical SiO2 completes without Ge nucleation on SiO2. On the basis of cross-sectional TEM images, the Ge arriving at the exposed Si surface diffuses into Si. The Ge nucleation, however, may occur on the exposed Si after the chemical SiO2 is consumed completely. We have previously shown that the erosion of chemical SiO2 proceeds in a random, non-homogeneous fashion.8 The Ge exposure first creates nanoscale (3- to 7-nm-wide) windows, whose areal density is on the order of 1011 cm-2, in the chemical oxide layer, while the thickness of the remaining oxide remains constant.8,16 We expect that the continued, low-flux Ge exposure enlarges these windows laterally. Figure 2 conceptually captures this mechanism where a circular window in an ultrathin SiO2 layer expands, as Ge adspecies on the exposed Si react with Si substrate and chemical SiO2 at an available reaction site (*) along the perimeter of the window to form volatile SiO and GeO:
Ge + Si + 2SiO2 + * ) GeO v + 3SiO v
(2)
782 J. Phys. Chem. C, Vol. 111, No. 2, 2007
Li et al.
Figure 2. Conceptual diagram showing an expanding window in an ultrathin SiO2 film to expose the underlying Si during Ge exposure.
The merging of windows and subsequent decrease in SiO2 loss rate during the final stage of window expansion will be neglected here to focus on the initial stage of window formation and expansion. On the basis of the proposed mechanism and the approximation, the SiO2 removal rate in mol/s can be defined as
R≡
FSiO2δ dASiO2 MWSiO2
dt
)-
FSiO2δ MWSiO2
2πr
dr dt
(3)
where ASiO2 represents the area covered by SiO2. FSiO2, δ, and MWSiO2 stand for density, thickness, and molecular weight of SiO2, respectively. r represents the window radius. We assume that the rate of Ge adspecies arriving at the window perimeter is much faster than the consumption rate of Ge via the reaction described in eq 2. The surface concentration of Ge, Si, and SiO2 along the perimeter is therefore assumed to remain constant near their saturation value, and the available reaction site (*) becomes the rate-limiting factor during the initial stage of window formation and expansion. With the above assumptions, the oxide removal rate is linearly proportional to the perimeter length by
R ) kCGeCSiCSiO2C*2πr
(4)
where CGe, CSi, and CSiO2 represent the surface concentration of Ge, Si, and SiO2, respectively. C* denotes the density of active sites per unit length of the perimeter. k is the reaction constant. The perimeter length is given by 2πr. By combining eqs 3 and 4, we obtain
MWSiO2 dr )k′ dt FSiO2δ
(5)
where k′ ) kCGeCSiCSiO2C*. Because r ) 0 when t ) 0, eq 5 can be expressed as
r)-
MWSiO2 FSiO2δ
k′t
(6)
After substituting eqs 5 and 6 into eq 3, we obtain
R)-
2πMWSiO2 k′2t FSiO2δ
(7)
Therefore, the remaining amount of SiO2, defined by NSiO2 ≡ 0 loss - NSiO , is given by NSiO 2 2 0 + NSiO2 ) NSiO 2
0 ∫0t R dt ) NSiO
2
πMWSiO2 k′2t2 FSiO2δ
(8)
Figure 3. Cross-sectional z-contrast STEM image of a Si substrate, portions of which are covered with 1.2-nm-thick chemical oxide (A) and 100-nm-thick dry oxide (B). The sample is exposed to a Ge flux at 1.0 × 10-1 equiv ML/s for 12 h. Note that the Ge diffusion through the chemical oxide layer forms an intermixed region (C) in Si, whereas no Ge intermixing is observed in Si below the dry thermal oxide.
Fitting a parabolic curve to the experimental data (--- in Figure 1) yields a correlation coefficient (R) of 0.998, suggesting that our model captures the SiO2 removal process with reasonable accuracy until the complete SiO2 removal. In comparison, when the Ge flux is increased by an order of magnitude (e.g., 5.0 × 10-1 equiv ML/s or 3.1 × 1014 atoms/ cm2‚s), Ge nucleation occurs on the exposed Si through the windows created in SiO2. Ge grows laterally over the remaining chemical SiO2 and coalesces into a Ge film before the erosion of chemical oxide completes. The above results suggest that the nucleation of Ge on Si and the reaction that laterally consumes chemical SiO2 are two competing surface phenomena. When the Ge flux exceeds the rate at which it can be consumed by the loss of chemical SiO2, the selective growth of Ge on Si dominates, thereby forming mushroom-shaped epitaxial Ge islands on Si.8 Several researchers have shown that Ge exposure degrades ultrathin (0.3-1.2 nm) thermal SiO2 films28,45,46 grown on Si at relatively low substrate temperatures (620-800 °C). The degradation was attributed to the sublimation of volatile GeO and SiO produced from a reaction between SiO2 and Ge. Contrary to this observation, we experimentally observe that Ge exposure does not degrade thermal SiO2 grown at a temperature over 1000 °C. For instance, when a 6-nm-thick thermal SiO2 is exposed to a Ge flux of 1.1 × 10-2 equiv ML/s at 700 °C for 48 h, the oxide thickness change is below the ellipsometer measurement error ((1 nm), and the surface roughness change is below the AFM measurement error ((0.3 nm). These results are repeatably obtained for different SiO2 thickness (6-200 nm), substrate temperatures (560-800 °C), and Ge fluxes (1.1 × 10-2 to 2.38 × 10-1 equiv ML/s). We therefore deduce that Ge does not react with SiO2 alone at the surface. Only in the case of ultrathin chemical SiO2 and presumably ultrathin thermal SiO2 of lesser quality grown at low substrate temperatures (620-800 °C), the following experimental observation suggests that Ge diffuses through the SiO2 layer to cause the SiO2 decomposition at the SiO2/Si interface. Figure 3 is a typical cross-sectional z-contrast STEM image of a Si substrate after Ge exposure. Before the Ge exposure, the center region of the substrate surface is covered by 100nm-thick dry thermal oxide, while the rest of the surface is covered by 1.2-nm-thick chemical oxide. In z-contrast STEM imaging, the characteristic scattering angle of electrons depends on the atomic number of chemical species present in the scanned domain, such that heavy atoms appear bright, and light atoms appear dark. Because Si and O have smaller atomic numbers than Ge, the chemical and thermal oxides appear as a dark line
Probing Interactions of Ge
J. Phys. Chem. C, Vol. 111, No. 2, 2007 783
Figure 4. Typical infrared absorbance spectrum of chemical SiO2 on Si MIR crystal shows characteristic peaks of isolated and bulk SiOH groups, while the negative absorbance of SiHx groups indicates that hydrides present after the HF treatment of the Si surface are removed after growing the chemical SiO2. The spectrum is taken after a 1.2nm-thick chemical oxide layer is grown on a hydrogenated Si surface.
(arrow A) and a dark region (arrow B), respectively. In comparison, Ge exhibits bright contrast against Si and O. This contrast shows up below the thin chemical oxide pointed by arrow C, indicating that Ge diffuses through the chemical oxide and into the Si substrate. Because of the low solid solubility (200 nm). The calculated SiOH concentration is 1.2 × 1022 cm-3, which is 4 orders of magnitude greater than the SiOH concentration (5 × 1017 cm-3)57 in dry thermal oxides grown at a temperature over 1000 °C. The high SiOH concentration strongly indicates that the chemical oxide is comparatively more porous than the thermal oxide. We speculate that these pores serve as diffusion paths through which Ge arrives at the SiO2/ Si interface. On the basis of the above experimental observations, the selective growth of Ge on Si over SiO2 cannot be explained by the reaction between Ge and SiO2 forming volatile GeO and SiO byproducts. The selectivity must therefore rely on a fundamentally different mechanism. To determine the surface phenomena responsible for the selectivity, we experimentally measure the desorption activation energy (Edes) of Ge adspecies from SiO2. The absolute impingement flux of Ge (Fi) is first calibrated against the Ge effusion cell temperature. At a fixed Ge flux, we incrementally decrease the SiO2/Si substrate temperature (Ts) and intermittently analyze the growth surface until the Ge3d photoelectron signal is detected. In this case, the XPS technique is used only to detect the Ge3d signal and the beginning of Ge island formation on the SiO2 surface, thus demarking the transition temperature at which the Ge adsorption rate equals the Ge desorption rate. For instance, when the effusion cell is maintained at 1130 °C, and the corresponding absolute Ge flux is 0.06 equiv ML/s, the Ge3d signal (Figure 5) appears after the SiO2 substrate is cooled down to 633 K and exposed to Ge for 30 min. Another 30-min Ge exposure at 613 K gives rise to a significant increase in the Ge3d signal. If we plot the integrated area of the Ge3d peak as a function substrate temperature, then a transition point is obtained at 648 K (see the inset of Figure 5), where the Ge3d intensity sharply rises above zero. We assume that the Ge desorption flux (Fd) perfectly counterbalances the Ge impingement flux (Fi) at this transition point. We also assume that the Ge sticking probability on the SiO2 surface is unity and that the Ge impingement flux is equal to the adsorption flux (Fa). That is, Fd ) Fa ) Fi ) 0.06 equiv ML/s at 648 K. Similar transition temperatures are
784 J. Phys. Chem. C, Vol. 111, No. 2, 2007
Figure 6. Plot of ln Fd vs 1/Ts. The slope yields Edes at 42 ( 3 kJ/ mol.
recorded for other Ge impingement fluxes. We finally plot ln Fd versus 1/Ts (see Figure 6) and calculate Edes from the slope. The magnitude of Edes is found to be 42 ( 3 kJ/mol, which is on the order of van der Waals forces,58-61 rather than a strong chemical bond. The pre-exponential factor calculated from Figure 6 based on the y intercept is approximately 103 equiv ML/s. Because the unit of Fd is equiv ML/s (1 equiv ML/s ) 6.3 × 1014 atoms/ cm2‚s), this pre-exponential factor can be divided by the steadystate Ge coverage (,10-7 ML) to estimate the attempt frequency of desorption (.1010 s-1). This attempt frequency is commonly reported as the pre-exponential factor for desorption rate constant. Because of the uncertainty in the absolute steadysteady surface coverage, the attempt frequency is left only as an estimate nearing the typical vibrational frequency of the surface adspecies (∼1013 s-1). The selectivity of Ge on Si over SiO2, when carefully delineated from Ge diffusion through lesser-quality, porous SiO2 and subsequent erosion of SiO2, is a direct result of the low desorption activation energy of Ge adspecies from the SiO2 surface. At typical growth temperatures of Ge on Si, the low desorption activation energy gives rise to a high desorption flux that exceeds the Ge impingement flux. The Ge adspecies thus desorb before forming stable nuclei on SiO2. If the surface temperature is decreased, however, then the desorption flux of Ge adspecies decreases exponentially. Conversely, when the impingement flux exceeds the desorption flux, the net positive Ge adspecies concentration on the surface leads to the formation of stable islands even on SiO2. The saturation Ge island density (Nmax) on SiO2 is measured at substrate temperatures of 300 and 400 °C to extract the critical island size (i) based on the mean field theory.38,39 Plan-view SEM images of Ge islands on SiO2 are analyzed to obtain the island density evolution at different stages. At a fixed substrate temperature and Ge flux, the surface undergoes morphological transitions, starting from island formation, to saturation of island density, to island coalescence as a function of exposure time. Figure 7 shows the time-dependence of the density of Ge islands grown at 0.5 equiv ML/min Ge flux and 300 °C substrate temperature. Typical plan-view SEM images for the different stages of island growth at 300 °C are also shown in Figure 8a-c. During the island formation stage, the Ge islands are sparse and far apart (Figure 8a), and the average island-island spacing is presumably longer than the characteristic diffusion length of Ge adspecies on SiO2. As a result, the Ge adspecies cannot be readily incorporated into the existing islands by surface diffusion. The Ge adspecies rather form stable nuclei between the existing islands. This new island generation process continues
Li et al.
Figure 7. Plot of ln N vs texp. The substrate temperature is 300 °C, and the absolute Ge flux is 0.5 equiv ML/min. The island density saturates after 40 min and decreases after 100-min Ge exposure.
Figure 8. Typical plan-view SEM images show Ge islands grown on thermal oxide at a fixed Ge flux of 0.5 equiv ML/min. Parts a, b, and c represent the islands grown at 300 °C before, during, and after the saturation of island density, respectively. Part d shows islands grown at 400 °C at the saturation of island density.
Figure 9. Plot of ln Nmax vs ln Fi. Note that the saturation island density is not a function of Ge impingement flux, although it is a strong function of surface temperature.
until the Ge island density is high enough (Figure 8b) to reach the saturation island density where the average spacing between islands is equal to the characteristic diffusion length of Ge adspecies on SiO2. The Ge adspecies are then readily incorporated into the existing islands by surface diffusion, and no additional islands are nucleated. The continued Ge exposure leads mainly to island growth, eventually causing the coalescence of islands (Figure 8c) and a subsequent decrease in the island density.
Probing Interactions of Ge
J. Phys. Chem. C, Vol. 111, No. 2, 2007 785
The saturation island density is a strong function of substrate temperature. At 400 °C, the saturation island density decreases by 2 orders of magnitude in comparison to that at 300 °C. A typical plan-view SEM image of Ge islands grown at 400 °C is shown in Figure 8d. However, the saturation island density is not a function of Ge impingement flux. Figure 9 shows the plot of saturation island density versus Ge impingement flux in logarithmic scale. The variation in the saturation island density is within an order of magnitude when the Ge flux increases from 1.6 to 36 equiv ML/s. The range of data points for 400 °C is limited compared to 300 °C because of the limited range of effusion cell temperature (1000-1200 °C) and available Ge impingement flux that can lead to island formation. On the basis of the mean field theory,38,39 the saturation island density (Nmax) is related to the impingement flux (Fi) by
ln Nmax ∝
i ln Fi, i+2
(10)
where i is the number of Ge atoms in a critical Ge island, which is in the metastable state, but the island becomes stable upon incorporating i + 1 Ge atoms. Note that the slope of ln Nmax versus ln Fi for both 300 and 400 °C in Figure 9 is close to zero, which suggests that i is equal to zero. Therefore, a single Ge atom can exist as a stable nucleus on SiO2. If the critical island size (i.e., monomer) of Ge on SiO2 is compared to the critical island size of six Ge units46 on Si, then the nucleation of Ge appears to be easier on SiO2 than on Si. Despite the stable monomers on SiO2, the ready desorption and fast surface diffusion of Ge adspecies prevent further island growth at typical substrate temperatures for film growth, becoming determining factors for selectivity. For example, at a substrate temperature of 600 °C, if the Ge impingement flux is set to 2.7 equiv ML/s, no Ge will nucleate on SiO2. This is because the desorption flux exceeds the impingement flux at that substrate temperature. However, the same impingement flux will lead to a Ge growth rate of ∼20 nm/min on Si. The selectivity is therefore achieved regardless of the interdistance between windows in the SiO2 template, provided that the desorption flux exceeds the impingement flux on SiO2. When the Ge impingement flux increases to 3.0 equiv ML/s at 600 °C, exceeding the desorption flux, the selectivity is still obtainable as long as the interdistance between windows in the SiO2 template is smaller than the characteristic diffusion length of Ge on SiO2. We approximate the characteristic diffusion length by measuring an average Ge island spacing from planview SEM images. For instance, the characteristic diffusion length is ∼2.5 µm at a substrate temperature of 600 °C. Because the interdistance between windows is on the order of hundreds of nanometers in our SiO2 templates, Ge nucleation is not observed on SiO2, although the impingement flux exceeds the desorption flux. 4. Conclusions We have gained phenomenological as well as quantitative understanding on Ge interactions with ultrathin chemical SiO2 films and thermal SiO2 films on Si. The high concentration (1.2 × 1022 cm-3) of OH appears to facilitate Ge diffusion through the 1.2-nm-thick chemical oxide layer, leading to Ge segregation at the SiO2/Si interface. Only in the presence of all three reactants (i.e., Ge, Si, and SiO2), the SiO2/Si interface becomes unstable, giving rise to the lateral consumption of chemical oxide and opening of windows that expose the underlying Si. On the basis of the parabolic relationship between the consumed amount of chemical SiO2 and Ge exposure time, the rate-limiting factor
in the oxide loss is determined to be the available reaction sites on the perimeter of the expanding windows in the SiO2 film. The thermally grown SiO2, which does not allow Ge diffusion, prevents the oxide loss and the formation of windows that expose the underlying Si. Decoupled from the oxide loss, we have determined that the weak VdW interaction between Ge adspecies and SiO2 is responsible for a low desorption activation energy (42 ( 3 kJ/mol) and a low diffusion barrier (∼13 kJ/ mol) for Ge adspecies on SiO2. At elevated temperatures, Ge adspecies readily desorb from SiO2 or diffuse away from SiO2 to segregate into exposed Si, provided that the interdistance between windows in the SiO2 template is smaller than the Ge diffusion length (>1 µm). Therefore, the selective growth of Ge on Si over SiO2 can be achieved without using selectivity controlling reagents. Our quantitative analysis on the interaction of Ge with SiO2 suggests that other template materials, such as Si3N4, Al2O3, and W, can be analyzed in a similar fashion to evaluate their selectivity in comparison to SiO2. In particular, the quantitative comparison in desorption activation energy and nucleation density should be able to provide an ideal design for the growth template, for which the interdistance between windows should be less than the diffusion length of stable Ge adspecies to achieve the selective epitaxial growth. Acknowledgment. We thank the National Science Foundation CAREER (DMR-0094145) for their financial support. The above material is also based upon work supported by, or in part by, the U.S. Army Research Laboratory and the U.S. Army Research Office under contract/grant number W911NF-05-10012. The National Science Foundation (CTS 98-71292 and National Nanotechnology Infrastructure Network) and the State of New Mexico support the TEM facility. References and Notes (1) Edgar, J. H.; Gao, Y.; Chaudhuri, J.; Cheema, S.; Casalnuovo, S. A.; Yip, P. W.; Sidorov, M. V. J. Appl. Phys. 1998, 84, 201. (2) Jin, G.; Liu, J. L.; Thomas, S. G.; Luo, Y. H.; Wang, K. L.; Nguyen, B. Y. Appl. Phys. Lett. 1999, 75, 2752. (3) Kamon, K.; Takagishi, S.; Mori, H. J. Cryst. Growth 1985, 73, 73. (4) Kitamura, S.; Hiramatsu, K.; Sawaki, N. Jpn. J. Appl. Phys., Part 2 1995, 34, L1184. (5) Li, Q. M.; Han, S. M.; Brueck, S. R. J.; Hersee, S.; Jiang, Y. B.; Xu, H. F. Appl. Phys. Lett. 2003, 83, 5032. (6) Okui, Y.; Jacob, C.; Ohshima, S.; Nishino, S. Silicon Carbide and Related Materials 2001, Pts 1 and 2, Proceedings 2002, 389-3, 331. (7) Sedgwick, T. O.; Berkenblit, M.; Kuan, T. S. Appl. Phys. Lett. 1989, 54, 2689. (8) Li, Q. M.; Jiang, Y. B.; Xu, H. F.; Hersee, S.; Han, S. M. Appl. Phys. Lett. 2004, 85, 1928. (9) Langdo, T. A.; Leitz, C. W.; Currie, M. T.; Fitzgerald, E. A.; Lochtefeld, A.; Antoniadis, D. A. Appl. Phys. Lett. 2000, 76, 3700. (10) Marchand, H.; Wu, X. H.; Ibbetson, J. P.; Fini, P. T.; Kozodoy, P.; Keller, S.; Speck, J. S.; DenBaars, S. P.; Mishra, U. K. Appl. Phys. Lett. 1998, 73, 747. (11) Fini, P.; Zhao, L.; Moran, B.; Hansen, M.; Marchand, H.; Ibbetson, J. P.; DenBaars, S. P.; Mishra, U. K.; Speck, J. S. Appl. Phys. Lett. 1999, 75, 1706. (12) Gale, R. P.; McClelland, R. W.; Fan, J. C. C.; Bozler, C. O. Appl. Phys. Lett. 1982, 41, 545. (13) Nam, O. H.; Zheleva, T. S.; Bremser, M. D.; Davis, R. F. J. Electron. Mater. 1998, 27, 233. (14) Park, J.; Grudowski, P. A.; Eiting, C. J.; Dupuis, R. D. Appl. Phys. Lett. 1998, 73, 333. (15) Li, Q.; Jiang, Y.-B.; Krauss, J. L.; Xu, H.; Brueck, S. R. J.; Hersee, S.; Han, S. M. Proc. SPIE 2005, 5734, 75. (16) Li, Q. M.; Pattada, B.; Brueck, S. R. J.; Hersee, S.; Han, S. M. J. Appl. Phys. 2005, 98, 073504. (17) Vescan, L. Mater. Sci. Eng., B 1994, 28, 1. (18) Vescan, L.; Dieker, C.; Hartmann, A.; Vanderhart, A. Semicond. Sci. Technol. 1994, 9, 387. (19) Aketagawa, K.; Tatsumi, T.; Hiroi, M.; Niino, T.; Sakai, J. Jpn. J. Appl. Phys., Part 1 1992, 31, 1432.
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