Spontaneous shape transition of thin films into ZnO nanowires with high structural and optical quality.

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nanowires with high structural and optical quality Sophie Guillemin,a,b Eirini Sarigiannidou,a Estelle Appert,a Fabrice Donatini,c Gilles Renou,d Georges Bremond,b and Vincent Consonni.a* a

b

Université Grenoble Alpes, CNRS, LMGP, F-38000 Grenoble, France.

Institut des Nanotechnologies de Lyon, Université de Lyon, UMR 5270 CNRS-ECL-CPE-INSA Lyon, 7 avenue Jean Capelle 69621 Villeurbanne, France. c

Université Grenoble Alpes, CNRS, Institut Néel, F-38042 Grenoble, France d

Université Grenoble Alpes, CNRS, SIMAP, F-38000 Grenoble, France

CORRESPONDING AUTHOR E-mail : [email protected]

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Published on 21 September 2015. Downloaded by University of Massachusetts - Boston on 26/09/2015 11:01:40.

Spontaneous shape transition of thin films into ZnO

Nanoscale

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ZnO nanowires are usually formed by physical and chemical deposition techniques following the

temperature. We demonstrate an original alternative approach for the formation of ZnO nanowire arrays with high structural and optical quality, which is based on the spontaneous transformation of a ZnO thin film deposited by sol-gel process following a simple annealing. The development of these ZnO nanowires occurs through successive shape transitions, including the intermediate formation of pyramid-shaped islands. Their nucleation under near-equilibrium conditions is expected to be governed by thermodynamic considerations via the total free energy minimization related to the nanowire shape. It is further strongly assisted by the drastic reordering of the matter and by recrystallization phenomena through the massive transport of zinc and oxygen atoms towards the localized growth areas. The spontaneous shape transition process thus combines the easiness and low-cost of sol-gel process and simple annealing with the assets of the vapor phase deposition techniques. These findings cast a light on the fundamental mechanisms driving the spontaneous formation of ZnO nanowires and, importantly, reveal the great technological potential of the spontaneous shape transition process as a promising alternative approach to the more usual bottom-up approach.

KEYWORDS: ZnO nanowires, nucleation, shape transition, sol-gel process, annealing.

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bottom-up approach consisting in supplying the reactants on a nucleation surface heated at a given

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Introduction Wide band gap semiconductors such as ZnO and GaN in the form of nanowires (NWs) are of great

benefit from their high aspect ratio at nanoscale dimensions, resulting in the occurrence of remarkable physical properties such as efficient light trapping and extraction,3 quantum and dielectric confinements,4,5 as well as stress relaxation at their lateral surfaces.6 Their integration into axial and radial NW heterostructures, including quantum wells and dots as well as semiconducting shells, can further provide additional functionalities in these nanoscale-devices.1,2 Until now, GaN and ZnO are well-known to spontaneously form high aspect ratio NWs along the polar c-axis via the bottom-up approach by a large number of vapor phase deposition techniques such as molecular beam epitaxy (MBE),7,8 physical vapor deposition,9-11 chemical vapor deposition,12 metalorganic and hybrid vapor phase epitaxy.13-16 The self-induced approach for the spontaneous formation of GaN and ZnO NWs can be achieved under relatively narrow growth conditions and does not require any catalyst or pre-patterning substrate using advanced lithography. N-rich conditions combined with a high growth temperature in the range from 700 to 850°C are typically used for the nucleation of GaN NWs by MBE.7,8,17 Low VI/II ratio and an intermediate temperature are favorable for the nucleation of ZnO NWs by metal-organic vapor phase epitaxy.18,19 The self-induced formation mechanisms of GaN and ZnO NWs are generally related to their strongly anisotropic wurtzite crystalline phase and are governed by thermodynamic considerations.17 The elongation of GaN and ZnO NWs is further driven by kinetic effects, in which the surface diffusion of gallium and zinc adatoms, respectively, plays an important role.20,21 Additionally, epitaxial strain effects and polarity effects originating from their structural relationship with the buffer or seed layers are also crucial to consider for elucidating these mechanisms.18,22-25 Alternatively, the use of wet chemistry as a low-cost, low-temperature, easily implemented process has received increasing interest for the formation of ZnO NWs.26,27 Nevertheless, their optical quality is still limited significantly and the difficulty to form optimized heterostructures

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potential for a broad variety of sensing, electronic, optoelectronic and photovoltaic devices.1,2 They

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including quantum wells for instance is detrimental. As a result, the self-induced formation of GaN and ZnO NWs with high structural and optical quality together with a high uniformity, reproducibility, and

previously, the bottom-up approach is based on the supply of reactants (i.e. gas species, chemical precursors…) on a nucleation surface heated at a given temperature in order to spontaneously form GaN and ZnO NWs. In this article, we present an original alternative approach to the more usual bottom-up approach consisting in supplying the reactants on a given surface heated at high temperature. It is based on the spontaneous transformation of a ZnO thin film deposited by sol-gel process into ZnO NW arrays with high structural and optical quality by a simple annealing. The spontaneous formation mechanisms of these ZnO NWs, including the intermediate nucleation of pyramid-shaped islands, are investigated by a wide variety of electron microscopy experiments. They further offer a deep insight into the driving forces involved in the spontaneous formation of ZnO NWs, which is achieved here from a ZnO thin film deposited by sol-gel process. Additionally, their high optical properties are thoroughly studied by lowtemperature photoluminescence (PL) and spatially-resolved cathodoluminescence (CL) measurements.

Experimental Section Fabrication of ZnO NWs by the spontaneous shape transition process The amorphous thin films as xerogel films were grown on top of (100) silicon substrates by sol-gel process using dip coating. The chemical precursor solution was composed of zinc acetate dihydrate (ZnAc2·2H2O) from Merck and monoethanolamine (MEA, C2H7NO) from JT Baker, which were dissolved with an equimolar concentration (i.e. typically 0.375M) in absolute ethanol. It was aged under continuous stirring for 24h, half of this time at 60°C. The amorphous thin films were then deposited by dip coating as xerogel films under controlled atmosphere (hygrometry zone axis of a ZnO NW nucleus 4. A typical low-magnification TEM image taken along the < ૚૚૛ with a diameter of 170 nm at its bottom and a length of 150 nm is presented in Fig. 2(a). The ZnO NW nucleus lies on a ZnO thin film with a thickness of about 26 nm, which is composed of very small ZnO grains with a mean diameter smaller than 8 nm, as shown by HRTEM, HAADF-STEM and ASTAR images of Fig. 2(c), 3(a) and 4, respectively. An amorphous phase also occurs between these ZnO grains as revealed in Fig. 2(c): this may be related to the native amorphous silicon oxide surface on the (100)

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pyramid-shaped islands, as revealed in Fig. 1(c). They present a mean diameter and length of about 81

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silicon substrate, on top of which ZnO thin film and nanostructures are nucleated. The ZnO NW nucleus is lightly tapered through the presence of slightly rough side facets with a tilt angle of about 75° with

for the pyramid-shaped islands. Furthermore, it has a wurtzite crystalline phase and is oriented along the polar c-axis, as shown by the SAED pattern in the inset of Fig. 2(a). Owing to the hexagonal section of ZnO NW nuclei as revealed in Fig. 1(c), their sidewalls are therefore composed of non-polar planes. Interestingly, the ZnO NW nucleus is almost free of any extended defects, apart from the occurrence of few I2-type stacking faults (SFs) at its very bottom. The I2-type SF sequence is AaBbCcAaCc.30 Two of ഥ૙ Bragg filtered image (see Fig. 2(e)), which reveals a them are shown in Fig. 2(d) together with a ૚૙૚ ഥ૙) planes characterizing the basal SFs in the wurtzite crystalline phase. It should be shift of the (૚૙૚ noted that the absence of a c/2 component in the 0002 Bragg filtered images allowed us to distinguish these I2-type SFs from the I1-type SFs. The formation of these SFs is probably related to the shape transition process itself that may involve stress effects. The chemical analysis presented by EDS-STEM images in Fig. 3(b-d) indicates that zinc and oxygen atoms are homogeneously distributed in the ZnO NW nucleus and thin film. In contrast, silicon atoms are mainly located in the substrate, but a small proportion of these atoms is also measured in the center of the thin film composed of ZnO grains. In contrast, no silicon atoms are revealed in the ZnO nanostructures. It is thus very likely that these silicon atoms are not incorporated in the ZnO grains, but are instead located in the amorphous phase between the grains. Furthermore, the local orientation of the ZnO NW nucleus and of each ZnO grain composing the underneath thin film is presented in Fig. 4 by ASTAR images. The ACOM-TEM (or ASTAR if combined with a precession module) tool is similar to the electron backscattered diffraction (EBSD) technique used in SEM but for TEM. A small electron beam (typically around 1.5 nm in diameter) is scanned point-by-point with a given step over an area of interest, while the diffracting signal is collected, indexed and interpreted via the template matching strategy in order to determine the local information.31 The use of Bragg spot diffraction patterns in ACOM-TEM approach rather than Kikuchi lines as in EBSD makes the ACOM-TEM a much more sensitive tool for both orientation and phase 8



respect to the substrate surface (see Fig. 2(b)), which is larger than the mean tilt angle of about 60 to 70°

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mapping. The orientation of the ZnO NW nucleus along the polar c-axis is confirmed in red color in Fig. 4(a), although the majority of ZnO NW nuclei are twisted with each other as seen in Fig. 4(b). More

diameter, as presented in Fig. 4(a-b). This is in contrast to the initial highly textured ZnO thin film as seen in Fig. 1(a), suggesting that massive recrystallization phenomena also take place during annealing. In particular, no specific nucleation sites consisting in c-axis oriented ZnO grains for instance are revealed, in contrast to the typical formation mechanisms of ZnO NWs by CBD.32 In that sense, the ZnO NW nucleus has no structural relationship with the underneath ZnO thin film. The spontaneous shape transition process of a ZnO thin film deposited by sol-gel process into ZnO NW arrays with high structural quality is expected to be driven by several chemical and physical processes during annealing, which are mainly based on thermodynamic and kinetic considerations. First of all, the spontaneous shape transition process certainly occurs under near-equilibrium conditions since a high annealing temperature of 900 °C is used. As a result, the preferential shapes of ZnO nanostructures are expected to be governed by thermodynamic considerations, especially through the minimization of its total free energy. In the absence of any strain and interface energy (related to the absence of any structural relationship of ZnO nanostructures with the ZnO underneath thin film) and of any significant edge energy, the main contribution to the total free energy is reduced to surface energy.33 It is well-known that the ZnO non-polar planes have the lowest surface energy of 1.16 J/m2 in the wurtzite crystalline phase34-36 as for GaN non-polar planes.37 In contrast, the cleavage energy of the polar c-plane is about 4 J/m2.36 The vertical sidewalls composed of non-polar planes with a low surface energy thus develop at the expense of the top facet consisting of polar c-plane, minimizing in turn the total free energy. Therefore, the NW shape is favorable owing to surface energy minimization.38 As regards the intermediate formation of pyramid-shaped islands, thermodynamic considerations may still hold but the surface energies of more complex side facets are not known. It should further be noted that the occurrence of pyramid-shaped islands has been reported in the self-induced formation mechanisms of both GaN and ZnO NWs grown by several vapor phase deposition techniques. The nucleation and

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interestingly, the underneath thin film consists of randomly oriented ZnO grains with a very small

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shape transition from truncated to full pyramid-shaped islands has been revealed prior to the formation of self-induced GaN NWs by MBE and has been attributed to the relief of the epitaxial strain generated

pyramid-shaped islands and NW nuclei are spontaneously formed on top of a ZnO underneath thin film and no structural relationship is revealed in Fig. 4. Also, the shape transition from full to truncated ZnO pyramid-shaped islands as shown in Fig. 1(b) is in contrast to the self-induced growth of GaN NWs by MBE.23 No pyramid-shaped islands have been revealed either in the self-induced growth of GaN NWs by MBE on an amorphous interlayer.39 In the present case, the faceting process is more likely related to surface energy minimization, as recently discussed for instance for the formation of selective area grown GaN NWs by MBE on a GaN buffer layer.40 It is worth noticing that the formation of selfinduced GaN and ZnO pyramid-shaped islands by MOVPE and PVD has been pointed out in Refs. 18, 24 and 41, and is usually strongly related to polarity effects with the nucleation surface. Interestingly, the formation of self-induced ZnO NWs by MOVPE has further been shown on top of ZnO pyramidshaped islands via the occurrence of inversion domain boundaries.18 However, their elongation does not proceed via the total resorption of the pyramidal base, as shown in the inset of Fig. 1(c), which may suggest, in the present case, the growth of steeper side facets toward the development of vertical sidewalls. Second, the spontaneous shape transition process is assisted by a drastic reordering of the matter, for which a very thin ZnO film (i.e. thickness < 10 nm) deposited by sol-gel process is considered in the first step of the process. This may occur through a massive transport of zinc and oxygen atoms, via their high surface diffusion at the annealing temperature of 900°C. It is well-known that the adatom surface diffusion, both on the substrate surface and on the sidewalls, plays a critical role in the spontaneous formation and elongation of GaN and ZnO NWs by vapor phase deposition.20,21 The typical diffusion lengths from several tens of nanometers to several micrometers has for instance been determined for GaN and GaAs NWs, respectively.21,42 It should further be noted that the massive transport of zinc and oxygen atoms is here a long range-process. The spontaneous formation of ZnO NWs takes place over

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from the AlN buffer layer:23 Nevertheless, epitaxial strain effects are here excluded since the ZnO

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some localized growth areas while bare surfaces specifically arise from the silicon substrate surface at a typical distance of several micrometers. In that sense, the spontaneous shape transition process exhibits

mechanisms for zinc and oxygen atoms may also be considered through successive desorption and adsorption processes between the ZnO nanostructures, as typically mentioned through collective effects for the elongation of semiconductor NWs by vapor phase deposition.43,44 Additionally, the spontaneous shape transition process described above is expected to depend on i) the deposition technique and properties of the initial ZnO thin film, on ii) the substrate used, as well as on iii) the annealing temperature. The starting point of the entire process is the formation of an amorphous thin film by sol-gel process using dip coating. The resulting xerogel film is mainly composed of zinc-oxoacetate nanoparticles (NPs) deposited at the surface of the native amorphous silicon oxide. It transforms, during the first step of the process, into a crystallized thin film consisting of ZnO grains. The complex chemical processes involved in the transformation may account for the continuous change of the morphology of the ZnO thin film (i.e. ZnO grain shape, size and orientation)) all along the four steps, which take place both underneath and between the different nanostructures. The initial thickness of the crystallized ZnO thin film is also crucial since the spontaneous shape transition process is driven by a massive transport of zinc and oxygen atoms. Varying the film thickness clearly affects the time scale of the process and the length of the resulting ZnO NWs is expected to be tunable. Furthermore, no structural relationship occurs with the silicon substrate, but only weak bonding are formed with the native amorphous silicon oxide surface, The structural role of the substrate can thus be neglected. Its chemical nature might nevertheless play a role, although the absence of incorporation process of silicon atoms during the process seems to exclude that. Eventually, while the high annealing temperature of 900 °C is used here, it should be noted that a threshold temperature above 800 °C is required to thermally activate the spontaneous shape transition process. Its dependence on the annealing temperature strongly affects the time scale and higher annealing temperatures can speed up the kinetic processes involved.

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the strong asset of a selective growth regime over some localized areas. Alternative transport

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In order to investigate the optical properties of these ZnO NWs grown by the spontaneous shape transition process, 5 K CL and temperature-dependent PL measurements are gathered in Fig. 5. A

Fig. 5(a). The CL spectrum is mainly composed of an intense narrow near band edge (NBE) centered at 3.36 eV and of a much less intense broad green emission band centred at 2.4 eV. The intensity of the NBE is typically about 60 times larger than the intensity of the green emission band. This clearly reveals the high crystalline and optical quality of these ZnO NWs with a high aspect ratio, which are fairly comparable to that of ZnO NWs grown by vapor phase deposition techniques. This is remarkable since the annealing developed here can easily be implemented, similarly to the typical low-cost deposition techniques in wet chemistry that usually exhibits a much less intense NBE and a strong visible emission band composed of various lines (see for instance Ref. 45 and therein). The study of the NBE is further carried out by spatially-resolved 5 K CL measurements collected on three kinds of nanostructures by using a low acceleration voltage of 5kV together with a small spot size (i.e. smaller than 10 nm), as presented in Fig. 5(b): single ZnO NW nucleus (i.e. short NWs), single ZnO NW with a high aspect ratio (i.e. long NWs), and the underneath ZnO thin film. It is further supported by temperature-dependent PL measurements recorded on an ensemble of ZnO NWs with a high aspect ratio in Fig. 5(c). The spatially-resolved 5 K CL spectra as well as the 5 K CL and 12 K PL spectra collected on an ensemble of ZnO NWs with a high aspect ratio are all identical. The fact that the ZnO NWs and underneath ZnO thin film exhibits the same optical properties clearly supports again massive recrystallization phenomena during the spontaneous shape transition process as well as the absence of incorporation process of silicon atoms. The NBE observed here is also fairly similar to the NBE of selective area grown O-polar ZnO NWs by CBD as shown in Ref. 46. It is governed by a large number of radiative transitions in the range of 3.355 to 3.376 eV, mainly consisting of five lines centered at 3.376, 3.372, 3.365, 3.361 and 3.357 eV. The present radiative transitions are all expected to be excitonic and the evolution of their energy position with the temperature can indeed be fitted by the formula proposed by Pässler et al., as seen in Fig. 5(d).47 It is worth noticing that the alternative formula

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typical 5 K CL spectrum recorded on an ensemble of ZnO NWs with a high aspect ratio is presented in

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given by Varshni and Bose-Einstein give similar results and thus can also be indistinctly used in the range of 12 to 70 K.48,49 Also, their activation energy is determined by fitting their intensity as a

transverse free A-excitons that are not resolved here are involved in the 3.376 eV line, for which an activation energy of 58.9 ± 2.8 meV is found. The 3.372 eV line that is not resolved by PL measurements is associated with ionized donor bound excitons (D+XA) and might be attributed to the I0 and/or I1 lines, which are typically related to aluminum and gallium, respectively.51 The activation energy related to the 3.365 eV line is low (i.e. 9.1 ± 1.5 meV) and hence ionized donor bound excitons presumably involving the I2 and/or I3 lines as well as the X lines exclusively attributed to the O-polarity are mainly considered.46,51-53 The rest of the excitonic radiative transitions such as the 3.361 and 3.357 eV lines are related to neutral donor bound excitons (D°XA).54 The activation energy associated with the 3.361 eV line is higher (i.e. 11.8 ± 3.2 meV) and the I6 and/or I8 lines again related to aluminum and gallium may more specifically be involved, which would be in agreement with the occurrence of the I0 and/or I1 lines at 3.372 eV.51 The broad shoulder at 3.357 eV, for which the temperature behavior in PL was not possible to follow, may be attributed to the I9 line related to indium. The implication of aluminum, gallium, and indium donor impurities are very typical since they are easily incorporated into the ZnO lattice, accounting for the success of n-type ZnO (see for instance Ref. 55). At lower energies in the range of 3.210 to 3.340 eV, radiative transitions involving longitudinal optical (LO) phonon replicas are found. The broad shoulder centered at 3.315 eV is mainly related to the first LO phonon replica of free A-exciton transitions and to the two electron satellites (TES) involving aluminum, gallium and indium.51,54 The 3.237 and 3.217 eV lines correspond to the second LO phonon replica of free A-excitons transitions and neutral donor bound transitions, respectively. Importantly, no intense emission occurs between 3.32 and 3.33 eV in contrast to ZnO NWs typically grown by physical vapor deposition. This points out the very low density of extended defects in the center of these ZnO NWs and especially the absence of I1 type-SFs, as previously deduced from the TEM images in Fig. 2.56

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function of temperature with an Arrhenius-type law, as presented in Fig. 5(e-g).50 Longitudinal and

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The green emission band of ZnO NWs with a high aspect ratio is directly related to a zero-phonon line (ZPL) centered at 2.85 eV. The ZPL is followed by six LO phonon replica lines (PRLs) centered at

(i.e. the LO phonon energy in ZnO at low temperature).55 The green emission band can be attributed to copper impurities substituting for zinc sites57-60 or oxygen vacancies.61-65 However, it should be noted that the chemical precursors used in the dip coating process contains up to 5 ppm of copper and that the green emission band related to oxygen vacancies is typically not structured.66-69 The highly structured green emission band here is thus predominantly assigned to copper impurities. In particular, the high annealing temperature used is expected to activate the copper doping in the center of ZnO NWs. The corresponding Huang–Rhys factor (i.e. S factor) is deduced by fitting the intensity of each ZPL and PRL by a Poisson distribution and equals 4.9 ± 0.3, as presented in the inset of Fig. 5(a). The present value is significantly smaller than the typically reported values in the range of 6.05 to 6.9,57,60,66,70-72 which could be accounted for by the high aspect ratio of the ZnO NWs at nanoscale dimensions. For instance, Chen et al. has pointed out that the strength of the electron-phonon coupling in ZnO (and correlatively the S factor) is reduced in NWs with respect to bulk single crystals.70 Also, the presence of a non-structured green emission band associated with oxygen vacancies located in the vicinity of free surfaces66 and centered at 2.45 eV cannot be excluded here and may presumably alter the predominant highly structured green emission band related to copper impurities.73

Conclusion While the typical bottom-up approach is based on the supply of reactants on top of a given surface heated at high temperature to spontaneously form ZnO NWs, we demonstrate an original alternative approach based on the spontaneous transformation process of a ZnO thin film deposited by sol-gel process into ZnO NW arrays with high structural and optical quality by a simple annealing. The development of these ZnO NWs takes place via successive spontaneous shape transition, involving the

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2.78, 2.70, 2.63, 2.56, 2.50, and 2.43 eV, each PRL being separated from each other by about 70 meV

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intermediate formation of full and truncated pyramid-shaped islands. The formation mechanisms are expected to be governed by thermodynamic considerations, in which the NW shape can minimize the

well as edge energy. It is further assisted by the drastic reordering of the matter via the massive transport of zinc and oxygen atoms towards the growth localized areas, which is also combined with concomitant massive recrystallization phenomena. Eventually, the resulting ZnO NWs are almost free of any extended defects, apart from few I2-type SFs at their very bottom. Their optical quality is high and marked by the very predominant NBE over the green emission band associated with copper impurities centered at 2.4 eV. These findings cast a light on the spontaneous formation mechanisms of ZnO NWs and open the way to the technological development of a promising approach with a great potential (i.e. low-cost and easiness of sol-gel process and simple annealing) that is alternative to the more usual bottom-up approach.

Acknowledgements The authors would like to thank Florence Robaut, Grenoble INP, France, for the specimen preparation by FIB as well as Edgar Rauch and Muriel Veron, Grenoble INP, France, for fruitful discussions about the ACOM-TEM/ASTAR technique. This work was partly supported by the Carnot Institute Energies du Futur through the project CLAPE and by Grenoble INP via a Bonus Qualité Recherche grant through the project CELESTE. Funding from la Région Rhône-Alpes via the Research Cluster Micro–Nano is also acknowledged. Sophie Guillemin held a doctoral fellowship from la Région Rhône-Alpes. The authors would also like to thank the facilities, and the scientific and technical assistance of the CMTC characterization platform of Grenoble INP supported by the Centre of Excellence of Multifunctional Architectured Materials "CEMAM" n°AN-10-LABX-44-01 funded by the "Investments for the Future" Program.

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surface energy (i.e. total free energy) in the absence of any significant strain and interface energy as

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FIGURE CAPTIONS

process of an amorphous thin film deposited by sol-gel process into ZnO NW arrays with a high aspect ratio by a simple annealing. (a) Top-view FESEM image of a ZnO thin film composed of large grains formed during the first step of the process. (b) Cross-sectional FESEM image of pyramid-shaped islands formed during the second step of the process. The insets are cross-sectional FESEM images with a higher magnification, revealing the occurrence of full and truncated pyramid-shaped islands. (c) Crosssectional FESEM image of ZnO NW nuclei formed during the third step of the process. The insets are top-view and 20°-tilted view FESEM images with the same scale bar, showing the hexagonal section of ZnO NW nuclei as well as their formation and development on top of pyramid-shaped islands. (d) Cross-sectional FESEM images of ZnO NWs with a high aspect ratio formed during the fourth step of the process.

Fig. 2 (a) Low-magnification TEM image of a ZnO NW nucleus spontaneously formed on top of the ZnO thin film during the third step of the process. The inset is a SAED pattern of the ZnO NW nucleus. (b) HRTEM image of the slightly rough side facets of the ZnO NW nucleus (solid rectangle area in (a)). (c) HRTEM image of the ZnO grains composing the underneath ZnO thin film (dotted rectangle area in (a)). (d) HRTEM image of the ZnO NW nucleus at its bottom, showing the occurrence of I2-type SFs ഥ૙ Bragg filtered image of (d) revealing the shift close to the interface with the ZnO thin film. (e) ૚૙૚ ഥ૙) planes that is characteristic of I2-type SFs. All TEM images are collected along the along the (૚૙૚ ഥ૙ > zone axis. < ૚૚૛

Fig. 3 (a) HAADF-STEM image of a ZnO NW nucleus spontaneously formed on top of the ZnO thin film during the third step of the process. (b-d) Corresponding EDS-STEM elemental mapping of the silicon, zinc and oxygen elements, respectively. The scale bar is 100 nm. 21



Fig. 1 Evolution of the morphology of ZnO nanostructures, showing the spontaneous shape transition

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DOI: 10.1039/C5NR04394H

Fig. 4 ASTAR maps representing the crystal orientation (a) along the growth axis of a ZnO NW nucleus

the growth axis of the underneath ZnO thin film. The color scales are given in the insets.

Fig. 5 Optical properties of ZnO nanostructures. (a) 5K CL spectrum recorded using a 600 grooves/mm grating of an ensemble of ZnO NWs with a high aspect ratio, showing their high optical quality. The ZPL and six PRLs composing the green emission band centered at about 2.4 eV are fitted by using a multi-Gaussian model. The corresponding Huang-Rhys factor is deduced by fitting the intensity of the ZPL and PRLs by a Poisson distribution, as shown in the inset. (b) 5K spatially-resolved CL spectra recorded using a 1800 grooves/mm grating, a low acceleration voltage of 5 kV as well as a small spot size (i.e. smaller than 10 nm) of a single short ZnO NW, a single long ZnO NW and the underneath ZnO thin film. (c) Temperature-dependent PL spectra in the range of 12 to 300 K recorded using a 600 grooves/mm grating of an ensemble of ZnO NWs with a high aspect ratio. (d) Evolution of the energy position of the three main lines denoted in (b) in the NBE as a function of the temperature in the range of 12 to 70 K. The evolution of the energy position is fitted by the Pässler formula. (e-g) Evolution of the intensity of the three main lines denoted in (b) in the NBE as a function of the temperature in the range of (e) 12 to 100 K and of (f-g) 12 to 60 K. The evolution of the intensity is fitted by an Arrheniustype law.

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spontaneously formed during the third step of the process as well as (b) in the basal plane and (c) along


Figure 1


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