An Atomistic View of the Incipient Growth of Zinc ... - ACS Publications

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An Atomistic View of the Incipient Growth of Zinc Oxide Nanolayers Manh Hung Chu,*,†,⊥ Liang Tian,‡ Ahmad Chaker,‡ Valentina Cantelli,‡ Toufik Ouled,∇ Raphael̈ Boichot,§ Alexandre Crisci,§ Sabine Lay,§ Marie-Ingrid Richard,∇ Olivier Thomas,∇ Jean-Luc Deschanvres,‡ Hubert Renevier,‡ Dillon D. Fong,∥ and Gianluca Ciatto*,† †

Synchrotron SOLEIL - Beamline SIRIUS, L’Orme des Merisiers, Saint-Aubin, F-91192 Gif-sur-Yvette, France Université Grenoble Alpes, LMGP and CNRS, LMGP, F-38000 Grenoble, France ∇ Aix-Marseille Université, CNRS, Université de Toulon, IM2NP UMR 7334, 13397, Marseille, France § Université Grenoble Alpes, SIMAP and CNRS, SIMAP, F-38000 Grenoble, France ∥ Materials Science Division, Argonne National Laboratory, 9700 S. Cass Avenue, Argonne, Illinois 60439, United States ‡

S Supporting Information *

ABSTRACT: The growth of zinc oxide thin films by atomic layer deposition is believed to proceed through an embryonic step in which three-dimensional nanoislands form and then coalesce to trigger a layer-by-layer growth mode. This transient initial state is characterized by a poorly ordered atomic structure, which may be inaccessible by X-ray diffraction techniques. In this work, we apply X-ray absorption spectroscopy in situ to address the local structure of Zn after each atomic layer deposition cycle, using a custom-built reactor mounted at a synchrotron beamline, and we shed light on the atomistic mechanisms taking place during the first stages of the growth. We find that such mechanisms are surprisingly different for zinc oxide growth on amorphous (silica) and crystalline (sapphire) substrate. Ab initio simulations and quantitative data analysis allow the formulation of a comprehensive growth model, based on the different effects of surface atoms and grain boundaries in the nanoscale islands, and the consequent induced local disorder. From a comparison of these spectroscopy results with those from X-ray diffraction reported recently, we observe that the final structure of the zinc oxide nanolayers depends strongly on the mechanisms taking place during the initial stages of growth. The approach followed here for the case of zinc oxide will be of general interest for characterizing and optimizing the growth and properties of more complex nanostructures.

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INTRODUCTION Zinc oxide (ZnO) semiconductors have received great attention because of their desirable properties such as a wide band gap of 3.37 eV, high exciton binding energy at room temperature, good conductivity, and high transparency in the visible region.1,2 These features make ZnO thin films highly suitable for room temperature ultraviolet optoelectronic devices and transparent conductive electrodes in both classical and new dye-sensitized hybrid solar cells.3,4 To date, ZnO layers with thicknesses ranging from hundreds of nanometers to micrometers have been successfully synthesized for device applications by various techniques such as radio frequency sputtering,5 pulsed laser deposition,6 molecular beam epitaxy,7 and metal−organic chemical vapor deposition.8 For future micro- and nanodevice applications, however, it is important to be able to grow ZnO over large areas with nanometer scale thickness control. As a consequence, research into ZnO growth by atomic layer deposition (ALD), a technique that permits the monolayer-by-monolayer deposition of highly conformal films, has intensified.9−12 Unfortunately, the mechanisms governing the initial stages of the ALD growth process are not fully understood. Investigations into the initial nucleation and © XXXX American Chemical Society

growth behavior as well as the atomic structure of the heterointerface are crucial for optimizing the ALD process and understanding ZnO structure−property relationships in micro- and nanodevices. Since this behavior depends on the growth parameters in a complex way and is difficult to predict theoretically, in situ experimental characterization tools able to provide insight into growth dynamics have progressively gained in importance in recent years. In particular, the ALD growth mode, which proceeds through distinct and successive deposition cycles, is highly suitable for study by in situ X-raybased techniques since the growth can be paused after each cycle for a time sufficient to perform the characterization. Moreover, X-rays can easily penetrate through the ALD gas environment. The nanoscale dimensions of the deposits from the initial ALD cycles, as well as the need to use pauses as short as possible to preserve chemical stability of the structure, require intense X-ray sources, such as third generation synchrotron facilities. Received: June 3, 2016 Revised: August 1, 2016

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DOI: 10.1021/acs.cgd.6b00844 Cryst. Growth Des. XXXX, XXX, XXX−XXX

Crystal Growth & Design

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XAS was performed across the Zn K-edge at 9.66 keV in fluorescence mode, using a four-element silicon drift detector (SDD) mounted 30° above the surface plane. The incidence angle was set 1.3°. As the surface was nearly horizontal and the X-ray polarization was linear horizontal, the electric field is perpendicular to the c axis of the wurtzite structure. In these conditions, X-ray absorption near-edge structure spectroscopy (XANES) measurements were more sensitive to the orbitals oriented parallel to the surface plane.20 In the energy scans, we used a step size of 1 eV and integration times determined by the counting statistics. XRF and XAS data have been corrected for nonlinearity,21 self-absorption corrections were negligible due to the very low layer thickness. XAS Data Analysis. Ab initio simulations of the Zn K-edge XANES were performed using the FDMNES code22 in dipolar approximation within a non-muffin-tin finite difference method (FDM) approach.23 The calculations were carried out for a bulkwurtzite ZnO model (model A) considering a temperature identical to the experimental one (250 °C), and the resulting lattice parameters were a = 3.2536 Å and c = 5.2102 Å.24 The cluster used in the simulation consisted of 76 atoms around the central Zn with a dimension of 6.1 Å. Smaller clusters were analogously created for the other models discussed in the XANES section, 3.4 Å for model B, and 3.9 Å for model C. X-ray polarization in the simulations was set perpendicular to the c axis of the wurtzite structure, as in the experiment. The Zn K-edge extended X-ray absorption fine structure spectroscopy (EXAFS) data analysis was performed using the IFEFFIT package.25 After background subtraction, the χ(k) EXAFS oscillations were weighted by k3, multiplied by a Hanning apodization window and Fourier-transformed to real space in the k-range [0−7.8] Å−1. Theoretical backscattering amplitudes and phase shifts for all single and multiple scattering paths were modeled using FEFF code26 integrated in the ARTEMIS program.25,27 The calculations assumed lattice parameters identical to the ones used for the XANES simulations.24 The EXAFS fitting was performed to the first two neighbors shells around the central Zn atom in the R-range from 1.0 to 3.6 Å. The many-body amplitude reduction factor S20 was fixed to the value (0.9361) determined from an EXAFS measurement of a metallic Zn foil. In the most relevant fits, the variables used were the interatomic distances and Debye−Waller factors in the first and second atomic shells, and the edge energy shift, with the coordination numbers fixed to expected wurtzite values (four O atoms in the first shell, 12 Zn atoms in the second shell).

In situ X-ray absorption spectroscopy (XAS) has been used to investigate the early stage growth behavior of HfO213,14 and, more recently, TiO215 by ALD. Fong et al.,16 using in situ synchrotron X-ray scattering and fluorescence, also reported on the characterization of ZnO thin films grown on a Si substrate during the first ALD cycles.16 Their findings suggest that ZnO films initially grow as nanoscale islands that successively coalesce. However, the techniques used provided limited information on the local and crystal structure of these islands. Very recently, we studied the evolution of crystalline texture and strain during the initial growth of ZnO via grazing incidence X-ray diffraction (GIXRD), exploring two kinds of substrates: Si(001) with its native amorphous oxide (a-SiO2) and crystalline Al2O3(001) (c-Al2O3).17 The GIXRD results showed the development of tensile strain during coalescence and indicated that the different symmetries of the substrate surfaces lead to distinct preferred orientations of the grains in the nanofilms. However, GIXRD cannot provide details regarding the atomic bonding and relative positioning at the local scale; furthermore, it did not provide any information on the first ALD cycle for the a-SiO2 substrate, where no diffraction was observed. In this work, by coupling in situ XAS with ab initio simulations of the X-ray absorption cross section close to the Zn K-edge and quantitative analysis of the fine structure oscillations, we shed light on the atomistic mechanisms taking place during the very first stages of the ALD growth of ZnO, prior to the development of a well-defined long-range structure. This atomistic view of the incipient growth provides the key to understanding the crystal study reported in our previous paper and is here discussed in relation to those results.

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EXPERIMENTAL SECTION

Growth. The ZnO nanofilms were grown using an ALD reactor custom built by STIGMA18 and already described in a previous work.17 The zinc precursor was diethylzinc (DEZn), with Ar as the carrier gas, and N2O was the oxidant. The pulse times for DEZn and N2O were 1 and 10 s, respectively, and nitrogen was used to purge the chamber for 15 s in between pulses. Hence, the total time for one ALD cycle was 41 s. The chamber was maintained at a pressure of 10 mbar, and the substrates at a temperature of 250 °C, within the ALD growth window. ZnO nanofilms were grown on both SiO2/Si(100) (from Sil’Tronix ST) and Al2O3(0001) (from Kyocera) substrates. Received as wafers with epi-polished surfaces, the substrates were cleaned prior to deposition in a 1:4 mixture of 98% H2SO4/33% H2O2 (piranha solution) for 15 min to enhance surface stability. Finally, the substrates were annealed in situ at 250 °C for 1 h in argon prior to deposition. X-ray Fluorescence and XAS. The reactor was mounted on the six axis tower of a Newport diffractometer installed at the SIRIUS beamline of synchrotron SOLEIL19 for in situ characterization. The beamline features a HU36 apple 2 undulator source, a Si(111) monochromator, and Pt-covered mirrors for harmonics rejection, it provides a photon flux on the sample of about 3 × 1012 ph/s at 10 keV. For this study, X-ray fluorescence (XRF) data were acquired at 10 keV during the growth cycles and continued for an equivalent time just after the end of each cycle, prolonging the last nitrogen purge. XAS measurements, along with the XRD reported in our previous paper and X-ray reflectivity (XRR), were conducted in nitrogen after completion of the XRF acquisition. The total time per cycle dedicated to X-ray characterization was roughly 1 h; such characterization was performed for each of the first 10 growth cycles, on which this present work is focused. After the first 10 cycles were completed, another 190 cycles of growth were performed, and the final thin films were also measured by XAS, in order to have a reference to which the first cycles can be compared.

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RESULTS AND DISCUSSION Figure 1 shows the evolution of the Zn Kα XRF signal from the sample recorded during the first 10 cycles of ALD of ZnO on aSiO2 (red circles) and c-Al2O3 (blue squares) substrates. The points reported in the figure represent the XRF “plateau” values measured in the pause time after each cycle (see Experimental Section). The Figure 1 inset shows the XRF signal measured “live” during each cycle, followed by the measurement of the same signal in the pause time until stabilization. The vertical line segments separate the end of the cycles from the start of the pause time. The time used for XAS and other characterizations (roughly 1 h) has been omitted in the time scale. However, it is evident that the XRF intensity at the beginning of each cycle matches well with the last points of the pause times relative to the preceding cycle. This testifies to the good chemical stability of the structures under study during the characterization time and rules out Zn desorption. We note that the fluorescence intensity is still increasing slightly after the end of the cation and anion precursors injections, stabilizing during the successive pause (Figure 1 inset). This indicates that a small amount of the Zn precursor remains in the chamber after the end of the cycle and continues reacting with the oxidizer for a few seconds. Figure 1 shows that B



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(for cycles 1, 2, 5, and 10). The best fits (solid lines) obtained from the quantitative data analysis described in Experimental Section are superimposed with the corresponding experimental spectra. We note that the k-range was limited to 7.8 Å−1 due to the desire to minimize in situ characterization time for each cycle to preserve the stability of the intermediate structures deposited. Figure 2, panels b and d shows the magnitude of the Fourier transform (FT) of the corresponding EXAFS spectra reported in Figure 2, panels a and c, respectively. The first and second peaks in the FT are related to the EXAFS contributions of the first coordination shell (O nearest neighbors) and second coordination shell (Zn next nearest neighbors), respectively. The best fits in real (R)-space (solid lines) are superimposed with the experimental FTs. The EXAFS FTs show that, for c-Al2O3 substrate, there is a progressive increase in the amplitude of the first peak (at an apparent distance of R ≃ 1.3 Å) going from ALD cycle 1 to cycle 5. As for the a-SiO2 substrate, instead, the amplitude of the first FT peak is steady for all cycles. With regard to the second shell peak of the FT (at an apparent distance R ≃ 2.9 Å) for cycle 1 on a-SiO2, its amplitude is sensibly reduced with respect to the subsequent cycles; a similar decrease, but less evident, takes place also for the c-Al2O3. The decrease of the second shell peak amplitude is associated with the line shape change of the k3-weighted EXAFS oscillations observed after cycle 1, in particular, in Figure 2a. We complete the qualitative inspection of the FT by noting that, for both substrates, the spectra relative to the final ALD cycle 200 (not shown) are very similar to those of cycle 10. It is also worth remarking that the difference in the FT peak amplitude with respect to previous reports (the first shell peak is higher than the second one, while the opposite is usually reported in the literature for room temperature data on bulk ZnO29−31) is in part due to the limited k-range used here, which cuts off a region where the amplitude of the Zn backscattering function is stronger than that of O,32 and in part to the different thermal behavior for the different neighbors shells,33 considering that our data are taken at T = 250 °C. After performing quantitative analysis of the EXAFS data, we present the main results in Figure 3. Figure 3, panels b and a

Figure 1. Zn Kα XRF signal recorded during the first 10 cycles of the ALD of ZnO on a-SiO2 (red circle) and c-Al2O3 (blue square) as a function of the growth cycle. Error bars (calculated by Poisson’s statistics) are smaller than the point size. In the inset: XRF signal taken “live” as a function of time. The vertical segments indicate the end of each cycle and the start of the pause time. The time intervals used for XAS and other characterizations are omitted.

the dependence of XRF on cycle number is well approximated by a linear function for both growth on a-SiO2 and c-Al2O3 substrates after cycle 1. However, the line slope is different from 0 (no growth) to cycle 1, indicating a slower growth rate during cycle 1. This is compatible with a substrate-inhibited growth model in which ZnO islands nucleate during the first cycle and coalesce during the second one.28 We also note that the XRF count rate is always higher for deposition on c-Al2O3, which appears to yield more growth than on a-SiO2 at any cycle. To gain insight into the very first cycles of the ALD growth, EXAFS at the Zn K-edge was performed. The EXAFS k3χ(k) data (open symbols) are shown in Figure 2, panels a and c for a-SiO2 and c-Al2O3 substrates, respectively. For the sake of clarity, we show in the figure only a few representative spectra

Figure 2. (a) and (c) k3-weighted background subtracted Zn K-edge EXAFS (open symbols) for cycles 1, 2, 5, and 10 of the ZnO ALD growth on aSiO2 and c-Al2O3 substrate, respectively. (b) and (d) Fourier transformed spectra (open circles) of the same EXAFS spectra for a-SiO2 and c-Al2O3 substrate, respectively. The best fits (solid line) are superimposed to all experimental spectra in both k- and R-space. The spectra are shifted in vertical for clarity. C



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Figure 3. (a) and (b) Interatomic distance of the Zn next-nearest-neighbors and nearest-neighbors, respectively, obtained from the EXAFS fitting for the first 10 ALD cycles of ZnO on both a-SiO2 (circle) and c-Al2O3 (square) substrate. First shell distances (RZn−O) are represented by open symbols, second shell distances (RZn−Zn) by solid symbols. (c) and (d) DW factor of the second and first atomic shells, respectively, as a function of the ALD cycle number for a-SiO2 (triangle) and c-Al2O3 (diamond) substrate. First shell DW factors (σ2Zn−O) are represented by open symbols, second shell ones (σ2Zn−Zn) by solid symbols.

shows the extracted first shell (RZn−O) and second shell (RZn−Zn) interatomic distances, respectively, as a function of cycle number. Figure 3, panels d and c shows the first shell (σ2Zn−O) and second shell (σ2Zn−Zn) Debye−Waller (DW) factors. We remark on three significant observations: (1) RZn−O is larger for the c-Al2O3 substrate than for a-SiO2 for the first three ALD cycles. The values then stabilize around those for the 200-cycle samples (represented as dotted lines in Figure 3b). (2) σ2Zn−O for the c-Al2O3 substrate is large for the initial cycles and progressively decreases until around cycle 4. Conversely, for the a-SiO2 substrate, this DW factor remains virtually constant. This underlies the progressive increase of the amplitude of the EXAFS FT first shell peak for the c-Al2O3 substrate shown in Figure 2d. (3) σ2Zn−Zn for the a-SiO2 substrate is significantly larger after cycle 1 with respect to the successive cycles, while a similar behavior is not observed for the c-Al2O3 substrate. This is related to the small amplitude of the EXAFS FT second shell peak for cycle 1 on the a-SiO2 substrate (Figure 2b). As for the last observation, a similar decrease of the FT second shell peak (with the first shell peak less or not affected) with respect to reference samples has been reported in EXAFS studies of oxides and other materials and attributed to different origins, such as Zn vacancies and surface effects,30 size effects,34 amorphization or nanopores formation following ion implantation,35,36 or the formation of ultrathin near-interface layers.37 All of these cases involve a large amount of disorder in the nextnearest-neighbors distance distribution. In the case of amorphization, some increase in the first shell DW factor is also expected, which is not evident for the a-SiO2 substrate, as shown in Figure 3d. The formation of a ∼1 nm-thick interface layer cannot be discarded a priori; however, we are not sensitive to the Zn−Si correlation in these EXAFS measurements.

Moreover, the formation of a two-dimensional (2D) interface layer would be at odds with previous works demonstrating the formation of nanoscale islands16,17,28 and with the change of slope observed by XRF after cycle 1 (Figure 1), indicating island formation. The most likely explanation is bond length disorder coming from the substantial contribution of non-fully coordinated Zn atoms at (and close to) the nanoisland surface, as it will be detailed later. The progressive decrease of σ2Zn−O for the c-Al2O3 substrate during the first four cycles, along with a negligible variation of σ2Zn−Zn, may suggest a mechanism related to oxygen vacancies (VO). In order to quantify the possible presence of VO, we performed fits of the experimental EXAFS varying the Zn coordination number (CN) and fixing σ2Zn−O to the value determined for cycle 6 with our conventional analysis (see Experimental Section). The corresponding extracted CN as a function of cycle number is shown in the inset of Figure 3d (green symbols). This result would imply a huge V O concentration after cycle 1, since CN ≈ 2.8, instead of 4. However, this scenario is unlikely because 30% VO would produce a significant shift in the Zn absorption K-edge (as shown in Supporting Information Figure S1), which was not observed experimentally. Low average CN for Zn atoms could also originate from a significant contribution of non-fully coordinated Zn atoms at (and close to) the surface (surface VO). However, apart from the missing absorption edge energy shift, we did not observe any broadening of the second shell distance distribution coming as an effect of the large surface-tovolume ratio. Since the σ2Zn−Zn shows a weak dependence on cycle number, our results suggest that the surface-to-volume ratio is smaller for ZnO on c-Al2O3 than on the a-SiO2 substrate. The dependence of σ2Zn−O for c-Al2O3 is indeed not trivial, and the broadening occurring during the first cycles may originate from a bimodal distribution of the Zn−O distance, as explained below. D



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Figure 4. (a) Experimental (solid symbol) and simulated (open square) Zn K-edge XANES spectra for cycles 2−10 of the ALD growth of Zn on aSiO2 substrate. (b) Experimental (full circle) and simulated (open circle) Zn K-edge XANES spectra relative to the first cycle on a-SiO2. (c) Experimental (solid symbol) and simulated (open square) Zn K-edge XANES spectra for cycles 2−10 of the ALD growth of Zn on c-Al2O3 substrate. (d) Experimental (full circle) and simulated (open circle) Zn K-edge XANES spectra relative to the first cycle on c-Al2O3. The insets in panels (b) and (d) represent models B and C, used in the simulations of the first ALD cycle for a-SiO2 and c-Al2O3 substrate, respectively. The spectra were shifted vertically for clarity.

no change is visible for grains on the order of 10 nm or more.44 In our case, the formation of nanoislands with an internal structure that includes grain boundaries during the first ALD cycle of ZnO on c-Al2O3 could even contrast with the effect of free atoms at the surface, accounting for the results of Figure 3b. It is worth noting here that grain boundaries can generate different amounts of structural disorder at the nearest and nextnearest neighbors depending on grain alignment and the ordering of Zn atoms in the basal plane.29,43 Additional information can be extracted from analysis of the XANES region for the same spectra. Figure 4 shows the Zn Kedge XANES spectra taken after each of the first 10 ALD cycles of ZnO growth on the a-SiO2 (left panels) and c-Al2O3 (right panels) substrates along with ab initio XANES simulations performed as described in the Experimental Section. Panels (b) and (d) show the experimental spectra and simulations for the first cycle on a-SiO2 and c-Al2O3, respectively; panels (a) and (c) show the same for cycles from 2 to 10, with the experimental spectra for cycles 3−10 superimposed (top olive spectra). It is evident that for both substrates, starting from cycle 2, the XANES spectra are very similar. However, the XANES spectra obtained at the end of the first ALD cycle on both substrates are considerably different from the others. In particular, the cycle 1 spectrum for the a-SiO2 substrate is characterized by a single broad peak (instead of two) in the region 9700−9760 eV, smoothing of the peak at 9683 eV (which becomes a shoulder) and a height decrease and broadening of the main peak at 9673 eV. The cycle 1 spectrum for the c-Al2O3 substrate shows similar but less marked modifications. Ab initio simulations of the Zn K-edge XANES spectra (represented by open symbols at the bottom of each panel in Figure 4) assist us in understanding these experimental results. The features for the experimental spectra for cycles from 2 to 10 are well reproduced by a simulation based on a bulk wurtzite

We now focus on the difference in the RZn−O bond length, which is longer for the c-Al2O3 than for the a-SiO2 substrate for the first three ALD cycles. We first estimated the propagation of the strain previously determined by in situ X-ray diffraction17 to the local scale, following a strategy already used by different groups for the case of zinc blende semiconductors.38,39 The extension of that approach to the hexagonal crystal system made clear that the effect of strain is very small and cannot explain the large differences observed in the experiment (see Supporting Information). A significant shortening of the bond lengths has been predicted in ZnO for atoms at and close to the surface: such shortening can be as large as ∼0.1 Å for the surface atoms in the (0001) direction and ∼0.02 Å between the surface and subsurface layer atoms in the (0001) plane.40 Also, VO are predicted to contract the interatomic distances,41 while in amorphous ZnO the Zn−O bond length is expected to be virtually identical to that of crystalline ZnO.42 For the first ALD cycles on a-SiO2 substrate there is no evidence of VO (neither significant shifts of the X-ray absorption edge energy, nor a decrease of the EXAFS FT first shell peak amplitude). Therefore, it is probable that the small values of RZn−O in Figure 3b stem from a non-negligible contribution of atoms located near the surface of the nanoislands. The importance of the atomic surface-to-volume ratio for the nanoislands grown on a-SiO2 substrate is in fact also suggested by the higher σ2Zn−Zn observed at the first cycle. Conversely, RZn−O for the first ALD cycles on the c-Al2O3 substrate never differs substantially from the “equilibrium” value after cycle 200 (the dotted line in Figure 3b). This may be due to the formation of larger islands on the c-Al2O3 substrate, which would reduce the contribution of surface atoms with respect to those within the island volume. It has been reported that in the case of nanostructured ZnO with a grain size ∼2 nm, the Zn−O bond length significantly elongates with respect to bulk ZnO due to larger interatomic spacing within the grain boundaries,43 while E



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wurtzite structure. The simulation reported in Figure 4b is an average of the simulations for each of these six positions and reproduces very well all of the features of the experimental spectrum for cycle 1 on a-SiO2. (We note that using Zn nextnearest neighbors shifted out of the basal plane along the c-axis results in worse agreement with the experimental data). As for the c-Al2O3 substrate, the best simulation is obtained by introducing four O nearest neighbors, two Zn next-nearest neighbors in the basal plane, and two O atoms in the third shell, which also permutate over the gray sites shown in Figure 4d inset. The latter simulation is similar to the one obtained by Cho et al.48 in 2-nm-thick ZnO films with very short-range order (SRO), which the authors called “amorphous”. For a true amorphous material, the order would be exclusively at the nearest neighbors. We also performed a simulation for amorphous ZnO, using a cluster with only four O nearest neighbors around the central Zn (also shown in Supporting Information, Figure S2): this model does not match the experimental XANES for the c-Al2O3 and a-SiO2 substrates. Ultimately, these XANES results suggest that, at the end of the first ALD cycle, the samples have developed an embryonal SRO, with the O nearest neighbors already at the expected distances and angles for the wurtzite structure and a weak order at the next-nearest neighbors located in the growth plane. This in-plane second shell order is weaker for ZnO on the a-SiO2 substrate as compared with c-Al2O3, which is in accordance with the smaller FT second shell peak shown in Figure 2. The XANES results show (more clearly than EXAFS in this case) that, even if very weak, some in-plane local order exists even after the first cycle. However, no out-of-plane SRO is developed beyond the first shell after cycle 1: this SRO starts after cycle 2 for both substrates and leads to recovery of the complete local wurtzite symmetry. We note that for the small clusters employed in models B and C, the use of a finite difference method (FDM) approach in the XANES simulations is fundamental.49 We used the XANES simulations to also investigate more carefully the effect of oxygen vacancies, which can alter the XANES spectra.50,51 We found that VO can induce minor changes to the Zn K-edge XANES, but high concentrations cause a shift in the absorption edge energy due to the modification of the average effective charge on the Zn cations.52 This edge energy shift becomes visible in the simulations (∼1 eV considering the resolution of our monochromator) only for VO concentrations of roughly 10% or more (Figure S1 in Supporting Information). Even though no shift is observable in our experimental data, this result indicates that a small percentage of VO may still be present. On the basis of all results shown in the previous sections, and considering also the XRD results previously presented,17 we construct an atomistic model of nanoscale island growth and coalescence for the two different substrates. Both the XANES and EXAFS results indicate a significant contribution of nonfully coordinated surface atoms after the first cycle, especially for growth on the a-SiO2 substrate (e.g., the different XANES spectrum and the EXAFS second shell peak amplitude reduction), demonstrating that the islands are relatively small. However, the lack of reduction in the Zn coordination number for the a-SiO2 substrate (which is identical to cycle 200 within the error bars) indicates that average island size after cycle 1 is larger than expected for canonical ALD growth (which should result in ∼0.2 nm per cycle53), in agreement with the XRR and XRF measurements. The reduced sensitivity to the surface in

ZnO model (model A) and consistent with those reported previously for bulk ZnO.45 This suggests that ZnO with the local symmetry of wurtzite is formed during the second growth cycle for both a-SiO2 and c-Al2O3 substrates. In order to gain insight into the modification of the local atomic structure of ZnO at the end of the first ALD cycle, we explored several other structural models. The first consideration was the effect of direct bonding with substrate atoms, i.e., the presence of Si or Al atoms in the second atomic shell around Zn; however, this model is unlikely based on the following considerations. X-ray reflectivity (XRR) measurements (see Supporting Information Figure S3) performed after the first ALD cycle of ZnO on cAl2O3 substrate give a ZnO thickness of 1.3 ± 0.5 nm. Although XRR measurements for the first cycle of ZnO on a-SiO2 did not allow the extraction of a reliable thickness, the slightly lower fluorescence intensity for cycle 1 on a-SiO2 substrate with respect to c-Al2O3 shown in Figure 1 suggests a ZnO thickness ∼1 nm. This means that the growth rate was faster during the first 10 cycles than from cycle 11 to 200, the final thickness of the film after 200 cycles being ∼20 nm as determined by transmission electron microscopy.17 We attribute this to DEZn vaporization system thermal issues.46 If the nanoscale islands16,17,28 formed during the first cycle were spheres or cubes of ∼1 nm dimension, we would estimate the percentage of Zn atoms in contact with the substrate to be at most 18%. In the case of partial surface coverage, the percentage could decrease further since this would mean that islands would be taller (XRR measuring the average thickness). Therefore, as the XANES measurement probes all the Zn atoms in the islands, the contribution of Zn atoms at the interface with the substrate is expected to be minor and not sufficient to account for the large differences observed in Figure 4b. Furthermore, the EXAFS data shown above fit models derived from bulk ZnO (see also Experimental Section), and the variation in distances values and distribution observed during the first cycles can be accounted for by introducing disorder in ZnO rather than allowing different bonding. Nevertheless, in order to better address this point, a simulation with Si atoms substituting for part of the Zn ones in the second shell around a Zn absorber has been performed, and the result is shown in Supporting Information, Figure S2: it is evident that the model does not reproduce the unique features observed in the XANES spectra for the first cycle. Even in the absence of a significant percentage of Si or Al in the second atomic shell of Zn, the presence of strong bonds between substrate atoms and oxygens at the interface could have an indirect effect on the XANES spectra, inducing a distortion of the local symmetry that can propagate into the inner part47 of the island, generating structural disorder. The disorder may originate, however, from other mechanisms such as the presence of defects, surface tension, or grain boundaries. In the models B and C used for the simulations reported in Figure 4, we introduced disorder in the second atomic shell around Zn. Since we cannot introduce here a DW parameter (and fixing the number of neighbors to the expected values for wurtzite) as we do in the EXAFS analysis, disorder is mimicked by reducing the number of next-nearest neighbors around Zn. In the schematics shown in the insets of panels b and d, the colored and gray atoms represent the number of atoms used in the fits for each model and the sites over which such atoms can permutate, respectively. In model B, the central Zn atom has four O nearest neighbors and only one Zn next-nearest neighbor that can occupy all six sites in the basal plane of the F



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Figure 5. Model of island size and distribution after completion of the first ALD cycle on a-SiO2 (a, c) and c-Al2O3 (b, d) substrates; red = O, gray = Zn. (a) and (b) are front-view magnifications of the nanoislands selected in the corresponding (c) and (d) top views. The c axis is oriented along the vertical direction in the front view, while the top views show the (001) sample surface.

average thickness of 1.3 ± 0.5 nm; in our model, we considered an average thickness of about 1.8 nm, assuming that islands cover roughly the half of the substrate area (as shown in the nanoisland top view in Figure 5d). Although the absence of a Bragg peak for ZnO on the a-SiO2 substrate after the first cycle means one cannot extract directly the domain size, since the XRF yield after cycle 1 is only slightly lower for a-SiO2 substrate as compared to c-Al2O3, the model shown in Figure 5 holds if nucleation density is higher for the case of a-SiO2 with respect to c-Al2O3) (roughly a factor 2). Therefore, after the first cycle of ZnO growth on a-SiO2, the smaller islands cover almost entirely the substrate surface, but remain unconnected (as shown in the top view in Figure 5c). The picture emerging from this interpretation of the results, i.e., the formation of wider 2D-like platelets with a lower nucleation density on c-Al2O3 substrate vs near-three-dimensional (3D) growth mode with a higher nucleation density on a-SiO2 substrate after the first ALD cycle is supported by highresolution TEM images that show a more granular structure immediately close to the substrate interface for the case of aSiO2 (Supporting Information Figure S4). The different incipient growth modes could be due to a poorer ZnO affinity for the silica surface as compared with alumina. The nucleation density depends in fact on the density of surface hydroxyl groups just prior the growth,54 which may be different even when an identical substrate preparation procedure is used, due to the different surface chemistries. In addition, for ZnO on cAl2O3, our ALD conditions are similar to those in which 2D growth has been already observed by other groups.55 As already noted, the Zn−O bond lengths are longer for growth on c-Al2O3 than for growth on a-SiO2; furthermore, the Zn−O bond length distribution shows a decreasing trend during (at least) the first three ALD cycles. The small fraction of missing nearest neighbors in the model shown in Figure 5b,d (with consequent small bond length distortion) does not allow us to account for the observed σ2Zn−O trend (Figure 3d) in the frame of a surface-to-volume ratio effect. However, we can

the case of the c-Al2O3 substrate (i.e., the smaller changes to the XANES spectrum and less damping of EXAFS second shell peak) can be accounted for within the framework of the nanoscale island model shown in Figure 5. In this model, islands on a-SiO2 substrate after the first ALD cycle (Figure 5a) have an aspect ratio close to 1: the height along the out-of-plane direction ∼0.8 nm, with a base size of ∼1.1 nm × 1.1 nm, while islands on c-Al2O3 substrate (Figure 5b) are larger and have a smaller aspect ratio (height ∼1.8 nm, base ∼4 × 4 nm). A rectangular cuboid shape has been used here only to simplify the calculations: a precise determination of the nanoisland shape is beyond the scope of this work. In the case of islands on the a-SiO2 substrate, bond counting gives an average number of O nearest neighbors = 3.53 and average number of Zn next-nearest neighbors = 7.46. This means that first shell coordination number is only 11% lower than the expected crystallographic value of 4, a deviation that is within the uncertainties of Figure 3d. However, the deviation in second shell occupation is larger (38%), leading to the reduction of the second shell peak in the EXAFS FT (Figure 2b) and the very large σ2Zn−Zn after the first cycle (Figure 3c), especially if we consider that the missing atoms also bring about additional disorder in the second shell distance distribution, which contributes to a smaller peak. As for the c-Al2O3 substrate, due to the larger island size, bond counting gives an average number of O nearest neighbors = 3.84 (4% lower than the crystallographic value) and average number of Zn next-nearest neighbors = 10.28 (14% lower). This accounts for the small differences in FT second shell amplitude and XANES line shape with respect to the bulk values. We remark that the relatively large base dimension of the islands and low aspect ratio are supported by the recent GIXRD performed around the (110) ZnO reflection after the first ALD cycle for the c-Al2O3 substrate,17 and the XRR measurement (Supporting Information Figure S3). From GI-XRD we estimated the in-plane domain size to be around 4 nm (corresponding to the nanoisland base in Figure 5a). The XRR result provides an G



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neighbors is sensitive to how the alignment of crystal planes takes place at the boundary between the nanodomains.29,43 After cycle 1 for the c-Al2O3 substrate, as represented in Figure 6, several smaller domains with the same in-plane orientation merge to form the bigger 4 × 4 × 2 nm3 islands. The small constituent domains tend to align in order to preserve the continuity of the crystallographic structure in the bigger 4 × 4 × 2 nm3 domains, adjusting the nearest neighbors distances to accommodate this alignment. This mechanism would be in some way similar to the one taking place in semiconductors doped with heterogeneous atoms, where deviation from the interatomic distances of the matrix is evident in the first shell around the dopant atom, but negligible in the second shell, due to the stiffness of the bonds.56 This would also remind one of atomic arrangements in ternary semiconductor alloys, where the unit cell lattice parameters follow the Vegard’s law, while the local interatomic distances deviate from the crystallographic average value to accommodate the different cation covalent radii.57 Our model accounts for a larger distribution of the first shell interatomic distances with respect to the second ones in ZnO/c-Al2O3. As soon as the domains grow and their populations stabilize, in parallel to establishing 2D growth with full substrate coverage, this grain boundary effect becomes less accentuated and disappears; the fact that the first shell distance “stabilization” is obtained only after cycle 4 (considering the error bars in Figure 3) indicates that growth equilibrium is not reached yet during the first three cycles. Notably, the first shell interatomic distances and DW factors reported in Figure 3 for the c-Al2O3 substrate stabilize roughly at the same cycle number where we observed initial relaxation of in-plane strain and stabilization of the domain size,17 i.e., when the crystalline substrate stops influencing the growth via epitaxial alignment. As for ZnO/a-SiO2, after cycle 1, the ∼1 nm islands are mostly unattached and no boundaries are formed, while after cycle 2 large domains form with boundaries aligned along the c axis, without any driving force for in-plane ordering. Since the X-ray polarization in our experiment was kept within the plane of the film, the XAS results are not very sensitive to bonds in the out-of-plane direction (see Experimental Section). This accounts for the lack of observation of effects related to the boundaries between the different nanoislands and contributes to the steadiness of first shell DW factor for all ALD cycles.

account for this result if the wider islands formed during the first cycle of ZnO growth on c-Al2O3 stem from the merging of smaller 3D domains (with lateral sides ∼1 nm, similar to the nanoislands formed on the a-SiO2 substrate), as illustrated in Figure 6. The figure shows a top view of the (001) surface of

Figure 6. Model of internal composition of the 4 × 4 × 2 nm3 islands shown in Figure 5b (c-Al2O3 substrate). The nanoisland is formed by smaller domains (dotted contours) which merge generating boundary regions characterized by virtually unchanged crystal plane alignment but with some distortions in the local distances. The atoms belonging to the boundary regions are highlighted in gold, and they constitute a non-negligible fraction of the total atoms.

one ∼4 × 4 nm island, where the colored dotted lines represent the contours of the domains; the atoms belonging to the regions where different domains superimpose are highlighted in gold. We assume that crystallographic planes have already starting aligning at the boundaries between the domains so that continuity of the long-range crystal structure is obtained (as shown in Figure 6): this is in agreement with the relatively large in-plane domain size measured from GIXRD. At the same time, we suppose that some “memory” of the original domains is retained at the local scale, where an elongation of the Zn−O bond length in the boundary region takes place, according to previous studies.43 Such elongation would induce a bimodal distribution of the interatomic distances (“red” atoms inside the domains and “gold” atoms at the boundaries), which causes a reduction of the FT first shell peak amplitude and an increase in the first shell DW factor in the EXAFS measurements, as observed experimentally. Furthermore, the longer average Zn− O bond lengths in Zn/c-Al2O3 with respect to Zn/a-SiO2 can be accounted for by the contribution of atomic pairs at the domain boundaries along with the lower surface-to-volume atomic ratio. Now, why is a similar trend in ZnO/c-Al2O3 not observed for the second shell distance distribution, which remains virtually constant until cycles 8−10? We know from previous XRD analysis that, while ZnO/a-SiO2 (for cycles ≥2) exhibits (001) fiber texture with grain boundaries largely perpendicular to the plane of the film and without any in-plane order, ZnO does not form randomly oriented in-plane grains on c-Al2O3. Conversely, ZnO/c-Al2O3 shows three main orientations, for which the relative populations are still varying during the first three to four cycles. The local order at the nearest and next-nearest

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CONCLUSIONS We have performed the first complete and quantitative study on the evolution of the local structure of ZnO during the early stages of atomic layer deposition, using in situ synchrotron radiation X-ray fluorescence, XANES and EXAFS spectroscopy. The short-range sensitivity of these techniques allows us access to the transient structure obtained after completion of the first cycle, which, due to poor long-range order, may be inaccessible by means of X-ray diffraction. We were able to shed light on the mechanisms taking place during the first stages of the growth on two different substrates: amorphous SiO2 and crystalline Al2O3. The present results, in combination with those from in situ X-ray diffraction recently published on the same samples, allow us to form a comprehensive atomistic view of the incipient growth behavior. Quantitative analysis of the EXAFS oscillations allows us to spot surprising differences in the growth mode for the two substrates during the first three cycles. In the case of the amorphous SiO2 substrate, variations in the interatomic H



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and software, respectively. Financial support for this work by ANR Moon (ANR-11-NANO-0014) is gratefully acknowledged. V.C. was supported by the Nanosciences Foundation (FCSN-2011-03CE), and D.D.F. was supported by an award from the Nanosciences Foundation and the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, Materials Sciences and Engineering Division. The experiment at the SIRIUS beamline benefited from the SOLEIL beam time allocation No. 20131343.

distance mean values and distributions are understandable by considering the surface-to-volume ratio of the nanoscale islands. The islands after the first cycle have roughly 1 nm size and an aspect ratio close to 1. They are densely distributed over the substrate even but (mostly) noncoalesced, following a 3D growth mode. Conversely, for the crystalline Al2O3 substrate, the first cycle of growth leads to larger 2D-like nanoplatelets of lower aspect ratio (4 × 4 × 2 nm3) and lower nucleation density, reducing the surface-to-volume ratio. These nanoplatelets are characterized by a well-defined crystal structure and in-plane orientation, driven by the epitaxial response to the crystalline substrate, despite the low coherence of the interface. The difference in growth mode can be due to a poorer ZnO affinity for silica surface compared to alumina. The clear trends of Zn−O bond lengths and DW factors with the cycle number observed for ZnO/Al2O3 suggest that the nanoplatelets originate from merging of smaller (∼1 nm) domains with similar orientation. The Zn−O bond lengths elongate at the interfaces between the constituent domains, and the two different bond lengths (in the inner domains and at the boundary) induce a characteristic broadening of the Zn−O distance distribution, along with an increase of the average distance compared to the case of ZnO/SiO2. The changes and amount of disorder in the interatomic distances progressively lessen with cycle number and become negligible after cycle 4, in parallel with the gradual stabilization of domain orientation and size and the establishment of full 2D growth behavior. The relevant features of the final ZnO thin films, such as its texture, roughness, and 2D character are connected to the atomistic mechanisms at play during the first ALD cycles. The new insight into the nucleation process provided by the in situ characterization approach used in this paper is essential for the optimization of the ALD of ZnO and other, more complex oxide nanostructures to be used for future applications in advanced optoelectronic and photovoltaic devices.

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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.cgd.6b00844. XANES simulations for O vacancies, XANES simulations for alternative models (bonding to substrate atoms, amorphous ZnO), XRR characterization, TEM images, and an analytic derivation of the formula predicting the effect of strain on the bond lengths for hexagonal crystals (PDF)

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REFERENCES

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. Present Address ⊥

(M.H.C.) International Training Institute for Materials Science, Hanoi University of Science and Technology, No 1, Dai Co Viet, Hai Ba Trung, Hanoi, Vietnam.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Synchrotron SOLEIL for general facilities placed at our disposal and in particular N. Aubert and P. Fontaine (SIRIUS beamline) for the help with the experimental setup I



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