Increasing Polycrystalline Zinc Oxide Grain Size by Control of Film

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Increasing polycrystalline zinc oxide grain size by control of film preferential orientation Lorenzo Fanni, A. Brian Aebersold, Monica Morales Masis, Duncan T.L. Alexander, Aicha Hessler-Wyser, Sylvain Nicolay, Cécile Hébert, and Christophe Ballif Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.5b01299 • Publication Date (Web): 16 Oct 2015 Downloaded from http://pubs.acs.org on October 18, 2015

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Crystal Growth & Design

Increasing polycrystalline zinc oxide grain size by control of film preferential orientation Lorenzo Fannia, A. Brian Aebersoldb, Monica Morales-Masisa, Duncan T.L. Alexanderb, Aïcha Hessler-Wysera, Sylvain Nicolayc, Cécile Hébertb and Christophe Ballifc,a a

Ecole Polytechnique Fédérale de Lausanne (EPFL), Institute of Microengineering (IMT), Photovoltaics and Thin-Film Electronics Laboratory, rue de la Maladière 71B, Neuchâtel CH-2000, Switzerland b Ecole Polytechnique Fédérale de Lausanne (EPFL), Interdisciplinary Centre for Electron Microscopy (CIME), Station 12, CH-1015 Lausanne, Switzerland c CSEM, PV-center, rue Jacques-Droz 1, Neuchâtel CH-2002, Switzerland

Abstract We investigate the structural evolution of polycrystalline zinc oxide films grown by low pressure metal-organic chemical vapor deposition. The goal is to achieve larger grains – leading to higher charge carrier mobility from lower grain boundary density – by controlling the grain orientation during growth. The results are twofold. First we describe how the combination of deposition temperature and gas flow influences the nucleation and film thickening stages: low temperature and high gas flow favor a high nucleation density and the development of c-textured films, whereas high temperature and low gas flow lead to a lower nucleation density and a-textured films. Second we demonstrate how a fine control of the film preferential orientation at the different growth stages allows the fabrication of films with grains that are 25% larger, hence improving the carrier mobility with respect to the reference film.

1.

Introduction

The combination of wide band gap, low light absorption and high electrical conductivity makes zinc oxide (ZnO) thin films excellent transparent conductive electrodes for applications such as photovoltaic,1,2 lighting3,4 or display devices.5 Each particular application requires transparent conductive oxides (TCOs) with specific properties, such as a high carrier mobility, some of which can be tuned by changing the film microstructure. In order to tailor the film microstructure, a detailed knowledge of the mechanisms governing film growth is necessary.6–8 Growth of films by chemical vapor deposition techniques depends strongly on the chemical reactions happening at the film surface.9–12 These reactions and hence film growth are influenced by controllable parameters such as substrate temperature and gas precursor flows.13–15 Based on current growth models of low pressure metal-organic chemical vapor deposited (LPMOCVD) ZnO,16,17 these parameters influence how adatoms form new nuclei on the substrate and determine adatom sticking coefficients on the different crystal facets or exposed grain faces, which leads to a growth-velocity anisotropy. Grains with their fastest growth direction perpendicular to the film substrate overgrow grains oriented otherwise and hence dominate the late-stage texture and morphology of the film.18 ZnO in its hexagonal wurtzite phase presents a non-centrosymmetric unit cell, leading to anisotropic properties and growth along different crystallographic axes. The prismatic c-axis is polar,

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while the m- and a-axis (lying on the basal plane) are not. As a result mechanical and electrical properties can differ significantly depending on film texture and associated axis orientation and grain morphology.19–21 For these reasons, the film preferential orientation is a key parameter to tailor ZnO film properties. In this work, we investigate the influence of substrate temperature and gas precursor flow on the film microstructure (bulk property), surface morphology and film preferential crystalline orientation (texture) over a large deposition parameter range. First, we establish a relationship between these parameters and the mean free path of adatoms and show how this variable controls the film growth behavior during nucleation and film thickening. In the second part of the work, we apply this knowledge: by controlling the preferential orientation during film growth, we deposit a film with larger grains and thus improve the carrier mobility. 2.

Experimental Section

ZnO samples were deposited by low-pressure metal-organic chemical vapor deposition on 4x4 cm2 0.5-mm-thick borosilicate glass substrates (Schott AF32). Diethylzinc ((C2H5)2Zn; DEZ) and water vapor (H2O) were used as precursors for zinc and oxygen respectively. The DEZ and H2O are brought into the chamber by simple precursor evaporation, i.e. without a carrier gas. The gases were injected into the chamber through a showerhead facing the hotplate (30x30 cm2) where the substrates are placed The distance between hot-plate and shower is 20 cm. A detailed description of the deposition system was reported by Steinhauser.22 The effect of deposition parameters, namely hotplate temperature (T) and total gas flow (Φ, H2O + DEZ) was investigated in the ranges of 100–250 °C and 75–450 sccm; the H2O/DEZ ratio was kept at 1 for all depositions. In the following the terms “high” and “low” for substrate temperature and gas flow are considered relative to the conditions which are optimized for films used as front electrodes in thin film solar cells: 170 °C, 150 sccm.14 During deposition the total pressure in the chamber was kept constant at 0.35 mbar. The thickness of the films, varying between < 10 nm (non-continuous films) and 2 µm, was measured by a stylus-profilometer (Ambios XP-2). The deposition rate was calculated as the ratio between film thickness and deposition duration. All films were non-intentionally doped, i.e. deposited without the addition of extrinsic dopants. Film texture and crystalline quality (grain size, local strain/defects) were assessed by X-ray diffraction (XRD; Bruker D8 Discover); measurements were performed using Cu-Kα radiation (λ = 1.542 Å) in the 2θ range of 30–80° (0.05°-step, 1 s/step). The full width at half maximum (FWHM) due to instrumental set-up was assessed to be ~ 0.1°. The film surface morphology and crosssections were analyzed by scanning electron microscopy (SEM; JEOL JSM-7500 TFE); the nuclei density and average nucleus size were derived from atomic force microscopy (AFM; Bruker Dimension Icon) measurements. Automated crystal orientation mapping (ACOM; NanoMegas ASTAR23) of cross-section samples was carried out in a transmission electron microscope equipped with a field-emission gun (JEOL JEM-2200FS). The microscope was operated at 200 kV in nano-beam diffraction mode and the beam was scanned in a raster across the specimen with steps of 5 nm. Further details on sample preparation and microscope settings are described elsewhere.24 The TEM cross-section samples for ACOM were prepared by Ar+-ion beam milling (Gatan Inc. PIPS) according to the method described by Dieterle et


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al.. 25 The grain size was estimated as the mean transverse intercept length from TEM crosssection samples (stereologic method) using a proportionality constant equal to 1.26 Charge carrier mobility was measured by means of an Ecopia Hall effect measurement system using the Van der Pauw configuration. 3.

Results and discussion

Chemical vapor deposition processes can be classified according to the process limiting the growth rate of the film: either surface reactions or mass transport.27,28 The two regimes can be easily distinguished in the Arrhenius plot of Figure 1a, which shows the dependence of deposition rate on temperature and total gas flow. Starting from low temperature, all the curves are characterized by a linear increase of the deposition rate (surface reactions limited) followed by a constant trend (mass transport limited). The temperature at which the transition between the two regimes occurs increases with increasing gas flow. Indeed, at constant chamber pressure, a higher gas flow increases the gas velocity in the chamber so shifting the transition temperature to a higher value, as explained by Carlsson and Martin.29 In terms of deposition rate and process control, the ideal conditions are in the surface reactions limited regime, but at the cusp of the transition to the mass transport limited regime. Such conditions allow high deposition rates and fine control over the process throughout the temperature range, once the gas flow has been set.

Figure 1. (a) Arrhenius plot of the ZnO deposition rate for various gas flows. The stars indicate the transition temperature between surface reactions and mass transport limited regimes. (b) Combined effects of gas flow and temperature on film preferential orientation (assessed by XRD) and surface morphology (observed by SEM). Film thickness ~2 µm. The combination of temperature and gas flow influences not only the deposition rate but also the surface morphology and preferential orientation, as shown in Figure 1b for films in their “final stage” (we define the “final stage” as that when the film achieves a thickness of 2 µm). The different surface morphologies observed in SEM micrographs depend on the material growth velocity along its various crystalline orientations. The small c-textured grains (top-left images) are formed when the growth-velocity along the c-axis is larger than along the m- or a-axis. On the other hand, hexagonal flakes characterizing the bottom-right images



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are due to a higher growth velocity along the basal plane (m- and a-) when compared to the caxis. At intermediate temperatures between these two cases, a-texture grains characterized by surface structures of wedges or pyramids are obtained. The change in film preferential orientation with temperature (c- to a- axis) has previously been related to the increase in adatom mobility enabling the growth of crystalline planes different from those with the lowest surface energy (c-plane16). The film lateral resistivity exhibits a minimum for a-textured films close to the transition to c-texture films,14 The absorptance of the films decreases monotonically with increasing temperature and decreasing gas flow due to the carrier concentration reduction. The optimal properties as electrodes in thin film silicon cells are those that give the best compromise between film absorption and film conductivity, and are therefore provided by a-textured films close to the transition with c-textured ones. It is noteworthy that Figure 1b shows how changes in film structure and crystalline orientation are achieved not only by modifying the temperature but also by adjusting the gas flow. In this section we set out the background of our work. In the next section, we clarify the link between temperature/flow and adatom mean free path, a key parameter to tune the film preferential orientation throughout the film evolution. 3.1 Adatom mean free path and film crystalline quality It is known that the formation of an ordered lattice requires that adatoms coming from the gas phase can travel long enough to find a good site where to attach. Hence, when the adatom mean free path is too short to find a low energy lattice site, the film will likely be amorphous.30 Conversely the degree of well-ordered crystalline material improves as the adatom mean free path increases. At the low temperatures studied here (< 200°C), adatom desorption is negligible. Consequently, the adatom motion on the film surface can be modeled as Brownian.31 If D is the adatom diffusion coefficient and τ the time before the adatom motion is stopped by the arrival of subsequent adatoms from the gas phase, the adatom mean free path can be expressed as:

〈〉 = √2 ∙

(1)

τ is related to the adatom flow from the gas phase ϕad and the lattice interatomic distance d0 by:

32,33

=

(2)

∙

Substituting (2) into (1) and considering that in first approximation the adatom flow is proportional to the gas flow ϕad ∝Φ gives the adatom mean free path as:

〈〉 =

∙

∝ ⁄

(3)

Seeing that the adatom mean free path is proportional to D/Φ, this ratio must therefore be maximized to achieve high quality crystalline material; it is noted that Brune et al. proposed a similar parameter (with deposition rate instead of gas flow) for diffusion of silver atoms on platinum.34 The adatom diffusion coefficient is a difficult parameter to measure; but it has


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been shown to monotonically increase with temperature;33 it can therefore be controlled by adjusting this parameter. We assessed the relative quality of the crystalline material by comparing the broadening of the XRD peaks. Note that TEM observations show that the films contain no detectable volume fraction of amorphous phase, other than the thin (sub-nm thick) disordered material along incoherent grain boundaries. Deviations from high crystalline material instead derive from: small grain sizes, associated with an increase in disordered grain boundary volume fraction; and from non-uniform local strain (microstrain), associated with an increase in the density of intragranular defects such as dislocations. These effects both translate into an increase in the measured FWHMmeas of an XRD peak, whose broadening can depend on contributions of crystallite size, FWHMsize, microstrain, FWHMstrain, and from instrumental setup FWHMinstr. For a Gaussian fitting peak, the following expression holds:35,36 = !" # $%!& # !&

%$(4)

Following from this, it is clear that, by removing the instrumental broadening, the resulting XRD peak width provides a comparative indication of the crystalline quality, here defined as increased grain size and/or lower microstrain/defect population, so corresponding to better overall material ordering.37,38

Figure 2. (a) Measured full width at half maximum (FWHM) of the (1 1 2( 0)-peak from XRD pattern for various deposition conditions. (b) Influence of temperature and gas flow on film crystalline quality: FWHM (only -size and -strain contributions). In the white areas no data were available since the (1 1 2( 0) peak was not present. SEM micrographs show the cross-section of the polycrystalline films grown under different conditions. Notice that the plot is based on 13 XRD-patterns, but for reasons of legibility only 5 curves are shown in part (a). In Figure 2, the crystalline quality is assessed by comparing the width of Gaussian functions fitted to the (1 1 2( 0)-peak, of different films with instrumental broadening removed. Despite the large differences between the (1 1 2( 0)-peak intensities at various deposition conditions, ACS Paragon Plus Environment


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the Gaussian functions fit the peaks well and are all characterized by a standard errors on FWHM-values below 2%, confirming the fairness of the comparison. The peak width narrows, and hence crystalline quality improves, as T increases and Φ decreases. This verifies the trend suggested by Eq. (3) that crystalline quality is directly related to the adatom mean free path Φ. A similar trend was observed for ZnO films synthesized by combustion chemical vapor deposition.39 The following observation points in the same direction: SEM crosssections of cleaved films shown in Figure 2 show that, when D/Φ increases, the fracture surfaces exhibit more clearly defined grains and grain boundaries. This is similar to reports for metallic thin films by Thornton40 and Barna et al.,41 and could correlate both with an increase in crystallite size, and with a decrease in microstrain and crystalline defects that leads to cleaner intergranular cleavage. 3.2 Nucleation & film thickening We have seen the effect of a D/Φ increase on the film final stage: a change in preferential orientation, larger grains, and crystalline quality increase. In this section we consider how the D/Φ ratio affects the evolution of the film. For the nucleation stage, we assess the influence of the growth parameters by means of three indicators: nucleus density, average nucleus size before coalescence (both measured by AFM); and grain orientations after coalescence (measured by XRD, film thickness around 15 nm). All the deposited samples exhibit island-like nucleation. Features of the ZnO-islands can be tuned by adjusting the deposition parameters as already shown by Baji et al. for atomic layer deposited ZnO.42 Table 1 shows that a low D/Φ leads to an increase in nucleus density and a decrease in nucleus height. Moreover, Table 2 shows that these particular conditions produce a c-texture as the film preferential orientation. Conversely increasing the D/Φ ratio reduces the density but increases the average height of the nuclei, with grains that are mostly oriented along the non-polar a- & m-axes. Table 1. Qualitative effect of deposition parameters on nucleus height and density (AFM) after 3” of deposition. Deposition conditions

Morphology

Average nucleus height

Nucleus density

~4 nm

High (~2000 µm-2)

~7 nm

Low (~50 µm-2)

low D/Φ (130°C, 170 sccm or 170°C, 435 sccm)

high D/Φ (170°C, 75 sccm)


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During the film thickening stage, depending on the relative growth velocities along the different directions, the preferential orientation set at nucleation is either preserved or overtaken by a faster one.18 A useful way to represent the evolution of the preferential orientation is by using texture coefficients.43,44 These indicators represent the relative intensity of a particular orientation in the XRD pattern with respect to the others for a given thickness. Table 2 shows the evolution of texture coefficients during film thickening for two different deposition conditions. In the deposition under a low D/Φ condition, the c-texture, set at nucleation, is preserved during film thickening. On the other hand, the deposition performed at high D/Φ favors the a-orientation which becomes preferential at the final stage.

Table 2. Evolution of the texture coefficients with film thickness under different deposition conditions (which were maintained constant during the process). Only the most prominent orientations are shown: c- (red); a- (green) m- (blue) and s- (violet).

low D/Φ (130°C, 170 sccm or 170°C, 435 sccm)

4

(210°C, 150 sccm or 170°C, 75 sccm)

PO at final stage

c-texture

c-texture

ill defined (likely non polar m- or aplanes)

a-texture

c-

3 2 1 0

s- m- a0

high D/Φ

PO at nucleation

Preferential orientation (PO) evolution texture coefficients

Deposition conditions

texture coefficients

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


1 film thickness (µm)

2

4 3

a-

2

c- m-

1

s-

0 0

1 film thickness (µm)

2

3.3 Increasing grain size by controlling nucleation As seen in the previous section, depending on the deposition conditions, the preferential orientation developed at nucleation is either preserved or overtaken by other orientations during film thickening. Numerical modeling based on competitive overgrowth has shown that when the preferential orientation at nucleation is biased away from the one at final stage, the ultimate average grain size is larger than when the nucleation preferential orientation is biased towards the one at final stage.45 Therefore a nucleation layer characterized by a low fraction of grains oriented along the fastest growth orientation should develop larger grains than a



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layer having a large fraction of grains oriented along the fastest orientation, as sketched in Figure 3a.

Figure 3. (a) Schematic of the preferential orientation evolution for unseeded and c-seeded films. (b) Top view SEM micrographs of the unseeded and c-seeded films. ASTAR TEM measured cross-section of unseeded film and c-seeded film. The inverse pole figure orientation maps show the grain orientation perpendicular to the substrate; a-orientation (green), c-orientation (red) and m-orientation (blue). Each map has been overlaid with the corresponding reliability index map. In this experiment, we confirm this hypothesis by growing a ZnO film under conditions close to the ones optimized for the application as electrodes in solar cells (medium D/Φ ratio, 180°C, 150 sccm) which leads to the a-axis being the fastest growth direction. Additionaly we grown c-axis film on a-axis seed layer, without observing any variation in grain size with respect to unseeded film. In general we noticed that c-axis films grown at high D/Φ ratio are less influenced by the substrate than a-axis ones. Likely the “epitaxial effect” is less effective on the c-axis films due to the low mobility of adatoms. In order to minimize the fraction of aoriented grains at nucleation, we deposited a seed layer (~ 50 nm thick) with c-axis as preferential orientation (low D/Φ conditions, 150°C, 250 sccm). After having deposited the cseed layer, we increased the D/Φ ratio, such that the a-axis was now the fastest growth direction. The impact of the c-seeded layer on the grain size during growth is shown in Figure 3b, which presents the ACOM results of two cross-sections (unseeded, c-seeded). Both orientation maps have been overlaid with a greyscale reliability index, as defined by Rauch & Dupuy.46 This index provides a measure of the reliability or uniqueness of each found orientation during the ACOM template matching procedure. The index takes high values if there is a unique template that provides a good match to the experimental diffraction pattern and is low if multiple well-matching templates are found. Low values are typically found at regions where grains are overlapping grains or at grain boundaries, as multiple templates (one corresponding to each grain orientation) provide an equally good description of the diffraction


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data. Therefore, dark vertical lines can be considered as boundaries separating different grains, with each grain being colored according to its crystallographic orientation. Green grains are the ones growing with their a-plane parallel to the surface, while red ones have their c-plane and blue ones their m-plane parallel to the surface. The c-seeded layer is evident as red-colored grains in the first 100 nm of film at the film-substrate interface. It further has a marked influence on the subsequent average grain size. The difference in average size between the two films is below 5 nm at a film thickness of 0.5 µm, but increases as the film thickens up until reaching a 23% difference at a film thickness of 2 µm (129 vs. 105 nm, for c-seeded and unseeded films, respectively). A similar difference in grain size between the two samples (26 ± 6 %) was independently derived from SEM top view images, thus confirming the statistical significance of the grain size difference. In average, grains of the c-seeded film are tilted away from the direction perpendicular to the substrate surface by a larger angle when compared to the unseeded film, which qualitatively agrees with numerical simulations.45 The increased grain size significantly improves the film conductivity. The main contribution is provided by the enhanced carrier mobility from 6 cm2/(V·s) for the unseeded film to 18 cm2/(V·s) for the c-seeded, and to a lesser extent by the carrier concentration that increases slightly from 8·1018cm-3 to 1·1019cm-3 It has been shown that for non-intentionally doped films the lateral transport of carrier is limited by the barrier potential formed at grain boundaries.47 The larger mobility characterizing the c-seeded film confirms that the detrimental effect of boundaries on electrical transport is reduced by having larger grains and so fewer boundaries impeding electron mobility. 4.

Conclusions

For chemical vapor deposited zinc oxide films we observed how the combination of temperature and gas flow influences their surface morphology, preferential orientation and crystalline quality. We propose how the adatom diffusion coefficient over gas flow ratio takes into account the effects of both parameters on the film properties. By analyzing the influence of this ratio on nucleation and film thickening stages, we identify how to control the film preferential orientation throughout the entire film growth. Combining this knowledge with the ‘survival of the fastest’ theory, we demonstrate an increase in grain size and improved charge carrier mobility by pre-seeding film growth with a nucleation layer biased away from ultimate preferential orientation. This procedure, here demonstrated with ZnO grown by LPCVD, can furthermore be extended to a range of chemical vapor deposition techniques, including atomic layer deposition, and solution process growth. 5.

Acknowledgments

We would like to thank Dr. Thomas Lagrange for useful insight on XRD data interpretation and Dr. Quentin Jeangros for the careful review of the manuscript. This work was funded by the Swiss National Science Foundation (FNS) under the project ZONEM (grant no. 137833). 6.

References



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(1) (2) (3) (4) (5) (6) (7) (8) (9) (10) (11)

(12) (13) (14) (15) (16) (17) (18) (19) (20) (21) (22)

(23) (24) (25) (26) (27) (28)

Wenas, W. W.; Yamada, A.; Konagai, M.; Takahashi, K. Jpn. J. Appl. Phys. 1991, 30, L441. Rau, U.; Schock, H.-W. Appl. Phys. A 1999, 69, 131–147. Xu, W. Z.; Ye, Z. Z.; Zeng, Y. J.; Zhu, L. P.; Zhao, B. H.; Jiang, L.; Lu, J. G.; He, H. P.; Zhang, S. B. Appl. Phys. Lett. 2006, 88, 173506. Park, S.-H. K.; Hwang, C.-S.; Ryu, M.; Yang, S.; Byun, C.; Shin, J.; Lee, J.-I.; Lee, K.; Oh, M. S.; Im, S. Adv. Mater. 2009, 21, 678–682. Chen, M.; Pei, Z. L.; Sun, C.; Gong, J.; Huang, R. F.; Wen, L. S. Mater. Sci. Eng. B 2001, 85, 212–217. Thompson, C. V. Annu. Rev. Mater. Sci. 2000, 30, 159–190. Kajikawa, Y. J. Cryst. Growth 2006, 289, 387–394. Wang, M.; Jiang, L.; Wang, Y.; Kim, E. J.; Hahn, S. H. J. Am. Ceram. Soc. 2015, 1–7. Roth, A. P. J. Appl. Phys. 1981, 52, 6685–6692. Oda, S.; Tokunaga, H.; Kitajima, N.; Hanna, J.; Shimizu, I.; Kokado, H. Jpn. J. Appl. Phys. 1985, 24, 1607–1610. Pan, M.; Fenwick, W. E.; Strassburg, M.; Li, N.; Kang, H.; Kane, M. H.; Asghar, A.; Gupta, S.; Varatharajan, R.; Nause, J.; El-Zein, N.; Fabiano, P.; Steiner, T.; Ferguson, I. J. Cryst. Growth 2006, 287, 688–693. Ye, J. D.; Gu, S. L.; Liu, W.; Zhu, S. M.; Zhang, R.; Shi, Y.; Zheng, Y. D.; Sun, X. W.; Lo, G. Q.; Kwong, D. L. Appl. Phys. Lett. 2007, 90, 174107. Wenas, W. W.; Yamada, A.; Takahashi, K.; Yoshino, M.; Konagai, M. J. Appl. Phys. 1991, 70, 7119–7123. Faÿ, S.; Kroll, U.; Bucher, C.; Vallat-Sauvain, E.; Shah, A. Sol. Energy Mater. Sol. Cells 2005, 86, 385–397. Fanni, L.; Delaup, B.; Niesen, B.; Milstein, Y.; Shachal, D.; Morales-Masis, M.; Nicolay, S.; Ballif, C. Mater. Res. Express 2015, 2, 075006. Nicolay, S.; Faÿ, S.; Ballif, C. Cryst. Growth Des. 2009, 9, 4957–4962. Nicolay, S.; Benkhaira, M.; Ding, L.; Escarre, J.; Bugnon, G.; Meillaud, F.; Ballif, C. Sol. Energy Mater. Sol. Cells 2012, 105, 46–52. van der Drift, A. Philips Res. Rep. 1967, 22, 267–288. Ellmer, K.; Diesner, K.; Wendt, R.; Fiechter, S. Solid State Phenom. 1996, 51-52, 541– 546. Pung, S.-Y.; Choy, K.-L.; Hou, X.; Shan, C. Nanotechnology 2008, 19, 435609. Fanni, L.; Aebersold, B. A.; Alexander, D. T. L.; Ding, L.; Morales Masis, M.; Nicolay, S.; Ballif, C. Thin Solid Films 2014, 565, 1–6. Steinhauser, J. Low pressure chemical vapour deposited Zinc Oxide for thin film silicon solar cells. Optical and electrical properties. PhD Thesis, University of Neuchâtel, 2008. Rauch, E. F.; Véron, M.; Portillo, J.; Bultreys, D.; Maniette, Y.; Nicolopoulos, S. Microsc. Anal.-UK 2008, No. 128, S5–S8. Aebersold, A. B.; Alexander, D. T. L.; Hébert, C. Ultramicroscopy 2015, 159, 112– 123. Dieterle, L.; Butz, B.; Müller, E. Ultramicroscopy 2011, 111, 1636–1644. Han, J.-H.; Kim, D.-Y. Acta Metall. Mater. 1995, 43, 3185–3188. Ingle, N. K.; Theodoropolus, C.; Mountziaris, T. J.; Wexler, R. M.; Smith, F. T. J. J. Cryst. Growth 1996, 167, 543–556. Ohring, M. In Materials Science of Thin Films (Second Edition); Ohring, M., Ed.; Academic Press: San Diego, 2002; pp 277–355.


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(29) Carlsson, J.-O.; Martin, P. M. In Handbook of Deposition Technologies for Films and Coatings (Third Edition); Martin, P. M., Ed.; William Andrew Publishing: Boston, 2010; pp 314–363. (30) Ohring, M. In Materials Science of Thin Films (Second Edition); Ohring, M., Ed.; Academic Press: San Diego, 2002; pp 495–558. (31) Mirica, E.; Kowach, G.; Evans, P.; Du, H. Cryst. Growth Des. 2004, 4, 147–156. (32) Lewis, B.; Anderson, J. C. Nucleation and Growth of Thin Films; Academic Press: New Yor, San Francisco, London, 1979; Vol. 4. (33) Reichelt, K. Vacuum 1988, 38, 1083–1099. (34) Brune, H.; Bales, G. S.; Jacobsen, J.; Boragno, C.; Kern, K. Phys. Rev. B 1999, 60, 5991–6006. (35) Bish, D. L.; Post, J. E. Modern Powder Diffraction; Reviews in Mineralogy; Mineralogic Society of America: Washington DC, USA, 1989; Vol. 20. (36) Karen, P.; Woodward, P. M. J. Solid State Chem. 1998, 141, 78–88. (37) Patterson, A. L. Phys. Rev. 1939, 56, 978. (38) Cullity, B. D. Elements Of X Ray Diffraction, 3rd ed.; Prentice-Hall, 2001. (39) Polley, T. A.; Carter, W. B. Thin Solid Films 2001, 384, 177–184. (40) Thornton, J. A. J. Vac. Sci. Technol. 1974, 11, 666–670. (41) Barna, P. B.; Adamik, M. Thin Solid Films 1998, 317, 27–33. (42) Baji, Z.; Lábadi, Z.; Horváth, Z. E.; Molnár, G.; Volk, J.; Bársony, I.; Barna, P. Cryst. Growth Des. 2012, 12, 5615–5620. (43) Harris, G. B. Philos Mag Ser 1952, 7, 336. (44) Moutinho, H. R. J. Vac. Sci. Technol. Vac. Surf. Films 1995, 13, 2877. (45) Ophus, C.; Luber, E. J.; Mitlin, D. Phys. Rev. E 2010, 81. (46) Rauch, E. F.; Dupuy, L. Arch. Metall. Mater. 2005, 50, 87–99. (47) Steinhauser, J.; Faÿ, S.; Oliveira, N.; Vallat-Sauvain, E.; Ballif, C. Appl. Phys. Lett. 2007, 90, 142107.

FOR TABLE OF CONTENTS USE ONLY Increasing polycrystalline zinc oxide grain size by control of film preferential orientation Lorenzo Fanni, A. Brian Aebersold, Monica Morales-Masis, Duncan T.L. Alexander, Aïcha HesslerWyser, Sylvain Nicolay, Cécile Hébert and Christophe Ballif



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Our work demonstrates enhanced carrier mobility by controlling the film preferential orientation during growth. Firstly, we determine the influence of the temperature and gas flow on film preferential orientation during CVD growth. Thereafter, by exploiting the survival of the fastest principle, we demonstrate grains 25% larger than the reference film leading to enhanced carrier mobility.


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Figure 3. (a) Schematic of the preferential orientation evolution for unseeded and c-seeded films. (b) Top view SEM micrographs of the unseeded and c-seeded films. ASTAR TEM measured cross-section of unseeded film and c-seeded film. The inverse pole figure orientation maps show the grain orientation perpendicular to the substrate; a-orientation (green), c-orientation (red) and m-orientation (blue). Each map has been overlaid with the corresponding reliability index map. 354x175mm (96 x 96 DPI)


Increasing Polycrystalline Zinc Oxide Grain Size by Control of Film

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