Methodology for Studying Surface Chemistry and Evolution during the

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Cite This: J. Phys. Chem. C XXXX, XXX, XXX−XXX

Methodology for Studying Surface Chemistry and Evolution during the Nucleation Phase of Atomic Layer Deposition Using Scanning Tunneling Microscopy Dickson Thian,† Yonas T. Yemane,† Shicheng Xu,‡ and Fritz B. Prinz*,‡,¶ †

Department of Applied Physics, Stanford University, Stanford, California 94305, United States Department of Mechanical Engineering, Stanford University, Stanford, California 94305, United States ¶ Department of Materials Science and Engineering, Stanford University, Stanford, California 94305, United States ‡

S Supporting Information *

ABSTRACT: We study the nucleation stage and growth of atomic layer deposition (ALD) on hydrogen terminated silicon (Si:H) by in situ and ex situ scanning tunneling microscopy (STM). STM allows the in-depth study of surface chemistry and evolution during the ALD nucleation phase. Here, the ALD systems studied to demonstrate this technique are ZnO via diethyl zinc (DEZ) and TiO2 via titanium tetrachloride (TiCl4). In-situ STM revealed that DEZ does not discriminate between different surface sites, in contrast to TiCl4 which shows a strong preference toward dangling or OH bonds. Continued deposition showed distinct island growth for TiO2 deposition on Si:H, versus homogeneous growth for DEZ. ZnO ALD exhibited a delay of approximately 5 ALD cycles in transitioning from lateral to vertical growth and nominal physical film closure occurred after approximately 12−15 cycles. STM observations of these ALD chemistries demonstrated the strength of this technique in quantifying film closure and the effects of surface termination and defects on ALD growth mode. This technique can be applied to the study of a broad variety of ALD systems.

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INTRODUCTION

There are multiple growth modes for ALD. An example is the layer-by-layer growth mode, on which the ideal ALD is based.1 However, studies of the initial growth stages of ALD have shown that at least in the first cycles of ALD, the growth is not linear.29−31 It has also been shown in theory that because of steric hindrance due to the finite size of a precursor molecule, it is impossible to achieve full monolayer growth of a material per cycle.25−27 Finally, one would not expect linear growth initially because as ALD proceeds, the deposited material slowly changes the energetics of the reaction between the ALD precursors and the surface. The surface changes from only substrate atoms to a mix of substrate and deposited atoms, then to only deposited atoms. Unless the free energy change associated with the precursor’s adsorption on the substrate or deposited material are the same, the growth per cycle would not be constant in this transient stage.25−27 A more involved growth mode which takes these surface changes into account is the island growth mode,24 where discrete islands of material are deposited and grow before coalescing into a continuous film. Other pertinent issues in ALD nucleation that have yet to be fully addressed include the effects of surface termination on

Atomic layer deposition (ALD) is a widely used material deposition technique in the semiconductor industry.1,2 A subset of chemical vapor deposition, ALD has many advantages including precise control of material thickness1 and composition,3,4 conformality on high aspect ratio structures,5 and a low thermal budget for deposition.1 ALD has been used in varied applications such as solar cells,6−12 dielectrics,13−16 supercapacitors,17−19 and catalysis.20−23 With the increasing demand for the miniaturization of electronics, ALD has been utilized to make thinner and thinner material layers, exposing some fundamental limitations and nonideal aspects of ALD growth. One of the primary advantages of ALD is that deposited material thickness depends linearly on the number of ALD cycles. However, there is often an initial transient stage where growth is nonlinear and highly dependent on reaction energetics which, in turn, is determined by substrate termination and surface chemistry.24−27 This initial stage of ALD growth is not fully understood and needs to be investigated further if very thin, conformal, pinhole-free films of various materials are to be deposited by ALD. For example, this is an acute problem in the deposition of thin metallic films.28 © XXXX American Chemical Society

Received: July 2, 2017 Revised: November 8, 2017 Published: November 16, 2017 A

DOI: 10.1021/acs.jpcc.7b06491 J. Phys. Chem. C XXXX, XXX, XXX−XXX

Article

The Journal of Physical Chemistry C

hydrogen terminated and atomically flat silicon surface which facilitated STM observation.58−61 ALD was performed either in situ in a home-built combined STM−ALD tool62 or ex situ in a home-built hot-walled ALD reactor.63 For ZnO deposition, the precursors diethyl zinc (DEZ) and H2O were used. ZnO ALD was performed within the ALD temperature window at 150 °C.47 DEZ was pulsed into the chamber for 0.1 s, followed by a 60 s argon purge. H2O was then pulsed into the chamber for 0.1 s, followed by another 60 s purge. This is repeated for n number of cycles as desired. For TiO2 deposition, the precursors titanium tetrachloride (TiCl4) and H2O were used. The deposition parameters were identical to ZnO ALD. ALD was characterized using STM,62 scanning tunneling spectroscopy (STS) and XPS. XPS was performed on a PHI VersaProbe Scanning XPS Microprobe with Al(Kα) radiation at 1486 eV. For all STS data shown, 10 different spectra at similar locations were taken and averaged. For in situ STM observations, Si(111):H was prepared as previously described and immediately placed in the in situ STM-ALD tool. The system was then pumped to rough vacuum (10−3 Torr) and heated to 150 °C. A topographic scan of the surface was taken when temperatures stabilized. The STM tip was withdrawn from the surface approximately 500− 1000 nm during each injection of precursors. It is then reapproached to the surface for imaging. Taking an in situ scan before and after deposition allowed observation of the effect of surface bond termination on precursor reactivity without a vacuum break between deposition and characterization. Because the STM tip is only withdrawn within its Z-piezo range, it is able to reapproach the surface and track the same area with minimal thermal drift. The chemical adsorption enthalpies of precursor onto H- and OH-terminated silicon were calculated by density functional theory with the package QUANTUM ESPRESSO.64 Norm conserved Perdew−Burke−Ernzerhof (PBE) exchange-correlation functionals were used and the kinetic energy cutoff for wave functions was set to 35 Ry, which is sufficient for system energy convergence. The H-terminated silicon was modeled by a slab (4 × 7) cut along the silicon [111] plane, with surface Si terminated by H atoms. By substituting each H atom to an OH group, the slab was also used to model OH-terminated silicon. Dipole corrections65 were incorporated to compensate the artificial electric field across the asymmetrical slab. The zerotemperature chemical adsorption enthalpy was calculated by the difference between the energy when the DEZ molecule is placed far away from the slab and the energy when one ethyl group is protonated by taking the H atom from the surface. All calculated states had undergone ionic relaxations with energy convergence. STM image analysis was performed using the Gwyddion image processing software. Images were plane leveled, line leveled and scar corrected, in that order. For ALD surface coverage calculations, the uncovered substrate area is identified using the watershed algorithm in Gwyddion. Grain size analysis was performed using the ImageJ image processing software. One hundred grains were selected indiscriminately from each image and their areas were computed. Their diameters were obtained from their areas assuming each grain was circular. The average grain diameter and standard deviations for each cycle number were calculated.

initial deposition, determining the growth modes of ALD for different chemistries, and when does film closure for a particular chemistry occur. Addressing these issues are especially relevant for selective area ALD,32 the growth of quantum dots by ALD33 and the deposition of mixed ALD films.2 To address these issues, simulations have been performed to check the reactivity of precursors to specific bonds, such as the Si−OH bonds common on silicon oxide, one of the most common substrates used for ALD. However, DFT calculations can be difficult to validate in practice because specialized in situ setups are often required to probe the reactivity of precursors to different substrate surface bonds. Among other techniques, STM,34,35 transmission electron microscopy (TEM),36−39 in situ ellipsometry,40 X-ray photoelectron spectroscopy (XPS),41 and atomic force microscopy (AFM)42,43 have also been used for nucleation studies. Each of these techniques has advantages and disadvantages. For example, XPS has limited sensitivity to minute amounts of ALD, and AFM, while relatively inexpensive, has limited lateral resolution which may be insufficient to detect deposition in the early cycles of ALD. Low energy ion scattering (LEIS)29,44,45 is a useful quantitative tool that has been used to study film closure for ALD, but few laboratories and commercial facilities worldwide have this capability. Though limited to conductive substrates, the STM has angstrom scale resolution that can be used to elucidate the surface reactions during the initial cycles of ALD. STM can also track the evolution of surface topology as ALD proceeds. STM has been previously used to study ZnS ALD growth on gold.46 However, gold as a substrate is not widely used in the semiconductor industry where silicon is the primary substrate. Ideally, our choice of a precursor-substrate system to study should be one that has been relatively well studied in literature as that would make it much easier for comparisons. The deposited material should also have multiple uses as a thin ALD film in industry. In this article, we describe both an in situ and an ex situ STM study of the nucleation stage of ALD on hydrogen terminated silicon. The example systems we have used to demonstrate the advantages of this technique are ZnO and TiO2 ALD using H2O and the precursors DEZ and TiCl4, respectively. ZnO is a well-known ALD system that has been reviewed extensively by Tynell et al.47 ZnO ALD has a relatively short nucleation stage and its nucleation has been studied ex situ by AFM on different substrates. That study found differences in the growth modes of ZnO ALD depending on the substrate (SiO2, GaN, or sapphire) used.48 TiO2 ALD1,38,49−53 via TiCl4 is known to be extremely selective to surface groups.54,55 TiO2 and ZnO can be used for a variety of applications.8,56,57 This work demonstrates the capability of in situ STM analysis to study the effects of surface termination on nucleation, mode of ALD growth, and film closure. The material systems studied can be interchanged for studying the nucleation stages of other materials such as other metal oxides or noble metals. Using a high resolution spatial mapping tool such as STM enables direct observation of surface evolution, with cycle by cycle resolution, as ALD proceeds. Here, we show a stark contrast in nucleation behavior between the two chemistries studied.

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METHODS Silicon (111) single side polished wafers were purchased and anisotropically etched in 40% NH4F to produce an oxide free, B


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Figure 1. (a−f) STM imaging of in situ TiO2 ALD deposition on freshly etched Si:H. (b) shows the sample after a single exposure to TiCl4. (c−f) Sample after the indicated number of full cycles of ALD. Scale bar = 50 nm. Images were taken at −1.5 V, 50 pA current set point. Blue arrows point to specific nucleation sites. Red circles mark out a few representative point defects. The green striped arrow in parts b and c show a dangling or OH bond in part b disappearing in part c, while the purple checkered arrow in part f marks a mobile TiO2 nucleate. Insets below parts d−f show the line profiles taken across a growing nucleate.

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RESULTS AND DISCUSSION Effect of Surface Termination on Initial TiO2 Deposition. To first qualify our methods, TiO2 deposition via TiCl4 and H2O was studied. This chemistry was chosen because DFT, XPS and other methods have shown high selectivity toward surface OH bonds.54,55 However, an in situ topographical approach to show that TiCl4 does indeed preferentially react with surface OH bonds had, to our knowledge, not been performed. Figure 1 shows the STM topographs during in situ deposition of TiO2 on Si:H, after 1, 5, 15, 20, and 25 cycles of ALD. The Si(111):H sample does not have a 100% hydrogen termination. As is evident from Figure 1a, four different types of sites can be seen. The blue arrows in Figure 1 point to brighter spots on the surface that are likely dangling bonds66−69 or −OH bonds.70,71 Defect sites, which are typically missing silicon atoms, are circled in red. Most of the surface that is not specifically highlighted is hydrogen terminated, represented by their flat profiles in the STM image. (Note that only a few of these dangling or OH and defect sites are highlighted.) The last type of site not highlighted but evident are the silicon step edges. Figure 1b shows the sample after a single exposure to TiCl4. Comparing Figure 1a and Figure 1b, it is evident that there is minimal reaction of TiCl4 on the Si:H surface. As the number of cycles increased, most of the surface remains inert, but at 15 cycles, some ALD growth of TiO2 begins at dangling bond or −OH sites. This is most apparent at the top right of Figure 1d−

f. Focusing on only one of these nucleation seeds and looking at its line profile, we can see that this growth progresses gradually from about 1 nm in height to about 3 nm in height from 15 cycles through 25 cycles. This evidence suggests that ALD growth is preferentially occurring on the nucleation seed instead of the hydrogen passivated silicon sites. Lateral broadening of the seed is most likely caused by growth on the edge sites of the seed itself. As ALD proceeds, new nucleation seeds are formed at other dangling bond or −OH sites. There is no specific growth at the defect sites circled in red, unlike certain other chemistries such as platinum,72 which initiates growth at defect sites when deposited on graphene. In parts b and c of Figure 1, a dangling or −OH bond (highlighted with the striped green arrow) that appeared in Figure 1b disappeared in Figure 1c. It is hypothesized that the dangling bond was unstable upon interaction with the STM tip during scanning and forms a bridge Si−Si bond with neighboring silicon atoms, explaining its disappearance. A large nucleate also appeared at 25 cycles (Figure 1f, marked by the checkered purple arrow) which was not there at 20 cycles (Figure 1e). Empirical observations indicate that larger nucleates (about 3 nm in height) can be dislodged by the scanning STM tip and pushed around the sample. This is because the tip was unable to get a tunneling current through a large TiO2 grain and thus tracks the substrate surface instead of the grain surface, resulting in the laterally displacement of the grain as the tip passed it. This does not happen for smaller particles as it is likely that the tip was able to get a tunneling C


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Similar to Figure 1, a representation of three different sites are circled in different colors. The sites circled in green show sites that are mostly terminated with hydrogen while those circled in red represent missing atoms/defect sites and blue represent areas that have a few dangling or −OH bonds. Lastly, silicon step edges are not highlighted but are evident. Comparing parts a and b of Figure 3, we highlight the same regions before and after DEZ exposure. The multiple particulates that coat the entire surface are believed to be reacted DEZ precursor (see Tynell et al.47 for multiple hypothesized DEZ reaction mechanisms on −OH surfaces). As is apparent, all three labeled sites reacted with DEZ. In the labeled green area (majority Si(111):H sites), particulates randomly cover the surface sites. For the sites labeled in red (defect sites), some of the defects remain depressions in the image, while one can be seen with a particle partially obscuring it. The blue sites (dangling bonds or −OH bonds) in the first image disappeared, indicating that they probably reacted with DEZ or were possibly chemically modified. There also appears to be no preferential deposition on the step edges of the Si(111):H surface. Comparison of Initial ALD Using TiCl4 and DEZ. In previous theoretical attempts to elucidate the mechanisms behind ZnO deposition via DEZ and H2O, Ren et al. showed that DEZ would react very favorably with −OH bonds on a silicon surface, while −H bonds remain intact.74 Thus far, theoretical studies focused on the reaction of DEZ with −OH terminated surfaces such as Si:OH or pre-existing Zn:OH.75,76 To explain these results, DFT calculations of the reaction enthalpy involved in the reaction of DEZ with fully terminated Si:H versus that of fully terminated Si:OH were performed. Calculations for DEZ on Si:H showed that there is a downhill enthalpy change of −0.833 eV, similar to the enthalpy change of −0.826 eV for DEZ reacting on Si:OH (see Supporting Information). This matches the in situ STM observation of the initial DEZ exposure which showed that experimentally, DEZ showed no discrimination between different surface sites, most likely because the reaction of DEZ with the different available sites are all energetically favorable reactions. Note that the calculations for ZnO only considered the overall enthalpy of reaction which indicated if a reaction was possible, but do not take into account the chemical kinetics and reaction pathway.

current through smaller particles to facilitate imaging. The particular large nucleate that appeared in Figure 1f is one such mobile nucleate that was moved by the STM tip into the image during a previous larger scan (not shown). Figure 2 shows XPS taken over two samples after 35 cycles of TiO2 ALD. XPS showed TiO2 deposition on the unetched

Figure 2. Ti XPS signals of two samples after 35 cycles of TiO2 ALD. The red line represents the Ti signal taken on an etched Si:H sample as imaged by STM in Figure 1. The blue line represents the Ti 2p peaks on an unetched SiO2 sample.

silicon oxide sample (SiO2), and no detectable deposition on the etched Si:H sample which was the same sample analyzed previously by STM in Figure 1. This reiterates not only the well-investigated selectivity of TiO2 ALD, but also demonstrates that compared to the STM, XPS is not sufficiently sensitive to study the nucleation stage of ALD when very little or selective deposition has occurred, given the lack of Ti signal when TiO2 islands were clearly seen in STM analysis. Effect of Surface Termination on Initial DEZ Deposition. For comparison, ALD of ZnO was performed with DEZ and H2O. Si(111):H before and after the single pulse of DEZ are shown in the STM images in Figure 3, parts a and b, respectively. The inset in Figure 3a shows a representative zoomed in image of the Si:H surface showing the 1 × 1 silicon atomic arrangement typical of a hydrogen terminated silicon (111) surface without surface reconstruction.73

Figure 3. STM images of sample taken at a bias of −2.0 V and 50 pA current set point. (a) Si:H surface showing areas with defects (red highlights), −OH or dangling bonds (blue highlights), or areas with mostly −H bonds (green highlights). Inset shows the well-ordered 1 × 1 Si:H atomic arrangement (inset scale bar = 1 nm). (b) Sample after a single exposure of DEZ showing random reaction with all surface sites without discrimination. Scale bar = 50 nm. D


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Figure 4. (a)-(f) STM imaging of ALD ZnO from 1 to 6 cycles taken at −2.0 V bias and 50 pA current set point. Scale bar =50 nm. (insets) Representative line profiles of their respective images at 4×, 5×, and 6× of ALD. Red dotted lines indicate the base, first, second, and third layer heights. Red circles in the figure highlights second or third layer growth. The blue arrows point to areas not yet covered by any ALD. (g) Typical ALD nucleation curve (adapted from Satta et al.29 with the permission of AIP Publishing.).

Investigating Nucleation Delay of ZnO. To provide further insight into the nucleation stage of ZnO, we continued to observe, topographically, the growth of ZnO on the silicon surface. This time, however, the deposition was done ex situ, and each sample was subsequently brought into the STM for surface analysis. Ex situ deposition was performed because it was found empirically that the electric field of the STM tip was strong enough to influence subsequent deposition if in situ STM and ALD were repeatedly done continuously and sequentially. This would alter the final amount of deposition within the analysis area. For TiO2 deposition, an in situ experiment was possible because of the extremely slow growth rate on Si:H, which allowed us to skip STM observation during certain cycles, mostly mitigating the effect of the electric field.

On the other hand, Figure 1 shows a distinctly different growth mode for TiO2 ALD as compared to the ZnO ALD in Figure 3. TiO2 ALD here follows the island growth mode described by Puurunen.27 Previous simulations54 for TiCl4 on silicon serve to explain this big difference in growth mode. The reaction of TiCl4 on Si:H was calculated to be thermodynamically unfavorable with an enthalpy change of +1.3 eV, as compared to a favorable enthalpy change of −0.16 eV when TiCl4 reacts with Si:OH. This suggests that reaction of TiCl4 on a hydrogen terminated surface is very unfavorable compared to reaction with a hydroxide surface.54 In our experiment, we showed in situ ALD growth observations corroborating that large difference in reactivity. E


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Figure 5. STM images of Si:H surfaces taken at a sample bias of 2.0 V and 50 pA current set point after ex situ ALD deposition of ZnO from (a−f) 7 cycles through 15 cycles. Scale bar = 100 nm.

mostly Si:H to mostly ZnO, there is still equal competition for the adsorption of DEZ, leading to a 2D to 3D transition. Two other observations from this sequence of images are that (1) the 3D growth phase starts without complete 100% filling of the 2D growth phase and that (2) the sizes of each new nucleation site remain approximately constant from cycle 1−4×. Only the number density of the nucleates increase. The first point illustrates a possible impact on the final surface roughness and quality of the deposited film. It is also possible that incomplete layer filling can cause the formation of bridges, where precursors form a bridge between the higher level growth, bypassing an incomplete lower level deposition. This suggests a mechanism for the formation of atomic defects in the form of a gap or vacancy in the film. Vacancies in the film ultimately impact the film density, which affects the film’s ability to act as moisture barriers for example. Densities of films deposited by ALD had been observed to change by changing ALD growth conditions (for example: growth temperature).77,78 Using this STM-ALD technique, we might be able to study how growth conditions change the fill factor in the nucleation stage which in turn affects vacancy formation and ultimately film properties. Film Closure for ZnO. Another issue in studying ultrathin ALD films is the question of when the ALD film closes. Measurements have been done to estimate the point of closure using LEIS.29,44,45 However, LEIS offers no information on the topography and state of the surface when the film closes. In Figure 5, we show STM images of ALD growth of ZnO from 7 through 15 cycles of ALD. Figure 5 presents a topographical view of the ZnO at the film closure stage (note that film closure is typically not at a single point but rather gradual). Using image processing software, we attempted to identify the amount of surface covered by the ZnO grains. The algorithm is described in the Methods, and the results are

Thirteen distinct Si(111):H samples were prepared as described above and each subjected to a distinct number of ALD cycles, 1−8×, 10×, 11×, 12×, and 15×. Figure 4 shows the STM images of ZnO ALD growth from 1× to 6×. In the first six cycles, the density of particulates deposited on the surface increases monotonically. As described in Satta et al.,29 the nucleation or transient stage of ALD proceeds in two phases, a 2D growth phase and a 3D growth phase before film closure occurs. These two different phases in the transient stage can be clearly distinguished in our observations of ZnO deposition. In Figure 4a through d (1 cycle to 4 cycles of ALD), the density of particles on the surface increases. However, each of the particles are uniformly about 2−3 Å in apparent height over the base substrate. This means that there is only 2D first layer growth across the entire sample. From the cross sectional profile seen in 5×, there is an appearance of taller particles of about 5 Å, denoting the start of 3D second layer growth. A height of 5 Å roughly matches what would be expected for a second atomic layer. At 6×, as shown in Figure 4f, the surface topography shows a few particles that are about 8 Å in height, likely representing the start of the third atomic layer growth. This sequence of observations indicates a transition from 2D to 3D growth during the nucleation phase of ALD on a silicon sample that would have been difficult to observe without the extended resolution of the STM as compared to a typical AFM. Such a transition is expected for ZnO. Afshar et al.75 calculated the enthalpy change for the half reaction of DEZ on ZnO to be about −15 kcal/mol. In comparison, the enthalpy change for the half reaction of DEZ on Si:H is approximately −19 kcal/mol, according to our own calculations. The reaction energy of DEZ with both types of substrate are similar and thermodynamically favorable. Thus, as the surface changes from F


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Figure 6. (a) Percentage of surface grain coverage with increasing number of ZnO ALD cycles. Theoretical 100% coverage is marked by the black dashed line. (b) Average ZnO grain diameter with increasing number of ZnO ALD cycles.

terminated sites, −H terminated sites and edge sites. As a result, ZnO ALD exhibited more homogeneous growth. Subsequent deposition of ZnO ALD was performed ex situ and characterized. The exact point of transition from 2D to 3D growth within the nucleation stage was observed to be around 5×/6× for ZnO. This transition occurred without the 2D growth fully filling the first layer deposition. Film closure for ZnO ALD is estimated to be between 12× and 15×, but the exact point of physical film closure is difficult to determine using topographical measurements, despite the utilization of image analysis. STM could be a useful technique to provide surface topology information during film closure for a complementary tool such as the LEIS. The observations here serve to probe the nucleation stage of ALD which can be difficult to study using less sensitive methods such as XPS. An in situ approach allows the observation of the influence of specific surface sites on different ALD chemistries, as our comparison between TiO2 and ZnO ALD demonstrated. This technique can play an important role in characterizing the effects of chemical functionalization of surfaces on ALD and investigating the role of defects for area selective ALD. In situ observations can also help guide simulations in matching thermodynamic calculations with experimental observations. Our findings can help improve the deposition of thinner layers using ALD which is especially useful where nominal physical film closure is sufficient for a particular application, such as adhesion or blocking layers. The general technique of STM−ALD observation can be applied to a variety of material systems. The findings presented here is specific to ZnO and TiO2 ALD chemistry using DEZ and TiCl4, respectively, but the same technique can be applied to study the nucleation stage in a wide variety of chemistries from pure metals ALD for diffusion barriers to metal sulfides ALD for solar cell applications.

plotted in Figure 6a. As expected, the percentage of surface grain coverage increases as the number of cycles increases. However, it reaches somewhat of a plateau as full surface filling is achieved and multistack grains start to form. The surface coverage stagnates at about 90%. Topologically, it is impossible to differentiate between the substrate and the deposited grains, and the image detection algorithm then identifies divots between grains as “substrate” even if the grains are already touching. (This detection overestimate is illustrated in the Supporting Information.) Figure 6b shows the average diameter of ZnO grains with the number of cycles. The grain diameter remains fairly constant as the number of ALD cycles increase, suggesting that for the nucleation stage of the ALD process at a fixed temperature, deposited grains remained at a fixed size. The number (and in turn, the density and coverage) of the particles increases instead. (7 cycles of ALD, Figure 5a, was not included in these calculations because of streaking during STM imaging, which would cause artifacts in this analysis.) Figure 6a shows that surface coverage asymptotes at 12×, indicating that further deposition does not increase the calculated surface coverage of ZnO. This means that any further ZnO deposition occurs on a closed ZnO film. Together with the physical topology shown in Figure 5, parts e and f, we can deduce that film closure likely occurs between 12× and 15×. Note that physical film closure does not equate to electrical film closure which would require conductivity measurements to determine.

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CONCLUSIONS

We demonstrated the capability of the STM to gain valuable insight into the nucleation stage of ALD both in situ and ex situ. The STM’s subnanometer resolution allowed us to study the very first few cycles of ALD where other current techniques for studying ALD nucleation are inadequate. Here the nucleation stage of ZnO and TiO2 ALD were studied as example systems. TiO2 ALD was shown to exhibit island mode growth on Si:H, with TiCl4 preferentially reacting with dangling or OH bonds on the silicon surface and the rest of Si:H remaining passive. In contrast, DEZ was observed to react with all available surface sites at random, including defect sites, −OH

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.7b06491. Description of the overestimation in using image analysis for calculating surface coverage, XPS data for the ZnO G


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deposition, models used in density functional calculations, and description of empty states versus filled states imaging in STM (PDF)

AUTHOR INFORMATION

Corresponding Author

*(F.B.P.) E-mail: [email protected]. ORCID

Dickson Thian: 0000-0003-2979-8629 Yonas T. Yemane: 0000-0001-6625-7990 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Part of this work was performed at the Stanford Nano Shared Facilities (SNSF), supported by the National Science Foundation under Award ECCS-1542152. D.T. gratefully acknowledges fellowship support from the Agency for Science, Technology, and Research. Y.T.Y. gratefully acknowledges fellowship support from the Ford Foundation and the Stanford DARE program. We gratefully acknowledge support from the Center on Nanostructuring for Efficient Energy Conversion (CNEEC) at Stanford University, an Energy Frontier Research Center funded by the U.S. Department of Energy (DOE), Office of Science, Office of Basic Energy Sciences, under Award No. DE-SC0001060.

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