CHAPTER V ZnO nanowire arrays
In this chapter, we compare two different routes of using the nanosphere lithography for the manufacturing of well-‐aligned, density-‐controlled ZnO nanowires by hydrothermal growth.
In addition to crystallographic and microstructural characterizations, we performed dye loading measurements in order to compare the surface area of the nanowires manufactured by both routes as regard to an unpatterned array.
Finally, the reversible surface wettability of
the samples was evaluated.
CHAPTER V -‐ ZnO nanowire arrays -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 147
1. Introduction -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 149
1.1 Synthesis procedure -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 149
1.2 Wetting properties -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 151
2. Experimental part -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 154
2.1 Synthesis of the samples -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 155
2.2 Morphological & crystallographic characterizations -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 156
2.3 Dye loading -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 156
2.4 Wetting properties -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 157
3. Results and discussion -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 157
3.1 Morphological and crystallographic properties -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 158
3.2 Surface area -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 161
3.3 Wetting properties -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 163
4. Conclusions -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 168
5. References -‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐-‐ 170
1. Introduction
ZnO nanostructures, particularly nanorods,
[1]nanowires,
[2]and nanotubes
[3]are attracting considerable interest because of their large effective surface area resulting from a high surface-‐to-‐volume ratio.
[4]One-‐dimensional (1D) ZnO nanostructures are promising candidates for applications in the fields of catalysis,
[5]field-‐emission devices,
[6]chemical sensors,
[7-‐9]nanogenerators,
[10]ultraviolet lasers,
[11]field-‐effect transistors
[12]and dye-‐sensitized solar cells.
[13-‐15]Most of these applications rely on ZnO characteristics such as the wide direct band gap (3.37 eV) or high electron mobility (100 cm
2V
-‐1s
-‐1).
Since the electrical and optical properties of ZnO nanostructures depend on their crystal structure, dimensions and morphology,
[16]shape and size of the ZnO nanostructures play a vital role for the performance of the devices.
1.1 Synthesis procedure
Several approaches have been developed for the growth of well-‐aligned 1D ZnO nanostructures. High-‐quality nanowires can be prepared by vapor-‐phase techniques such as vapor-‐liquid-‐solid growth (VLS),
[17]vapor-‐solid growth (VS)
[18]or chemical vapor deposition (CVD)
[19, 20]but these techniques require sophisticated equipment, fine tuning of the experimental conditions, single crystal substrates and high temperatures (up to 900 °C) which are not compatible with organic substrates and low production costs. Electrodeposition can be used to produce pure ZnO nanowires under low temperatures, but it is limited to conductive substrates.
[21]To date, the hydrothermal method stands out as the most attractive alternative for obtaining well-‐aligned nanostructures of ZnO under mild conditions with simple and cheap implementation. In the first step, the substrate is coated with ZnO nanoparticles (seeds). In the second step, each seed acts as a nucleation site for the formation of ZnO nanowires under hydrothermal growth conditions. Experimental parameters, such as Zn
2+concentration, organic additives or growth time and temperature determine the final nanowire dimension and quality.
[22]The nanowires density has also been reported to have a strong impact on the efficiency of related devices.
[23]Therefore, in order to control the size, position and arrangement of ZnO nanowires, various technologies have been used to force a pattern-‐arranged growth of nanowires.
They include conventional photolithography,
[24, 25]laser interference lithography
[26]and
e-‐beam lithography,
[27, 28]which can be used to manufacture nanopores with controlled
shape and spacing. However, the complexity of the preparation process combined with
high initial equipment costs makes those lithographic techniques unfavorable for many
researchers. Recently, Kang et al.
[29]reported the use of microcontact printing to directly
pattern ZnO nanoparticle seeds. However, although this process allows the rapid
replication of a similar pattern, it does not control the nanowire density inside the patterned area.
Nanosphere lithography appears therefore as a very promising approach, due to its rapid implementation and its compatibility with wafer-‐scale processes. It consists in two steps: the preparation of a colloidal crystal mask made of nanospheres and the deposition (often by sputtering) of the material of interest through the voids between the spheres. In a last step called lift-‐off, the mask is removed and the (sputtered) layer keeps the ordered patterning of the mask interstices. In recent years, nanosphere lithography has attracted growing interest due its potential to manufacture a wide variety of 1D, 2D or 3D nanostructures.
[30, 31]In the ZnO research field, nanosphere lithography has so far mainly been used to create ordered arrays of metal nanodots that act as catalysts for nanowire growth by a vapor-‐
phase technique.
[32, 33]Some recent studies investigated with some success the possibility to combine nanosphere lithography and solution-‐based growth of ZnO.
Li et al.
[34]employed monolayers of polystyrene colloids to guide the growth of hexagonally patterned ZnO nanopillar arrays on zinc foils. However, the developed method was only optimized for an individual nanopillar growth at each interstice between neighboring colloidal spheres. Moreover, they faced difficulties to obtain perfect arrays of ZnO in a large area because of the defects in the mask.
Pyun et al.
[35]first reported the formation of ZnO nanotubes with the assistance of polystyrene colloids on ZnO seed layer prepared by metal-‐organic chemical vapor deposition (MOCVD), which is a rather expensive technique that fits poorly with the expected lowering of production costs. They observed that single-‐crystalline nanotubes were formed just below the nanopsheres, whereas solid nanorods were grown in the absence of PS colloids or at the apertures between three adjacent nanospheres. They didn’t evidence any significant effect on the luminescent properties of the ZnO nanotubes compared to those of the ZnO nanorods.
Hexagonally packed ZnO hemisphere-‐array films formed by growth-‐hindered nanowires were synthesized by Chang et al.
[36]Although, their nanostructures did not correspond to a templated growth of well-‐aligned ZnO nanowires with a high aspect ratio, they evidenced an increased surface area compared to unpatterned ZnO nanorod films.
However, the ZnO hemisphere-‐array and the unpatterned ZnO films were respectively synthesized by hydrothermal growth and by electrochemical deposition. According to us, this impedes any reliable comparison of their properties.
Finally, Fragala et al.
[37]attempted to pattern the seed layer by using polystyrene
microspheres for the selective ZnO nanorod hydrothermal growth. They observed a
bimodal ZnO nanorod growth as a result of a local nuclei concentration gradient in the
different regions of the patterned substrate.
In our study, we compare two different nanosphere lithography routes, which will be further referred to as the “templated growth” (TG) and the “templated seeding” (TS) procedures, to establish a low cost synthesis of patterned ZnO nanowire arrays over a large area on glass-‐FTO substrates. In addition to crystallographic and microstructural characterizations, we performed dye loading measurements in order to compare the surface area of the nanowires manufactured by both routes as regard to unpatterned array.
1.2 Wetting properties
Recently, there have been lot of interest in studying the wetting properties of metal oxides nanomaterials (mainly of ZnO
[38, 39], TiO
2[40, 41]and Al
2O
3[42]), which can be reversibly switched between superhydrophobicity and superhydrophilicity by alternation of ultraviolet (UV) irradiation and dark storage.
[38, 43]Indeed, as nanoscale devices aimed for chemical and biological sensing, surface wettability plays a very important role and surfaces with controllable and reversible wettability are highly desirable,
[44]particularly for the control of effective micro-‐ or nano-‐fluid motion.
[45]In nature, many surfaces (from plants or animals) are highly hydrophobic.
The best-‐known example of a hydrophobic surface is the leaves of the lotus plant.
[46]Numerous studies confirmed that the chemical composition (a wax) of the leaves combined with a specific micro/nano hierarchical surface structure (Figure V -‐ 1) provide the lotus plant its unique superhydrophobic
*and self-‐cleaning properties.
Figure V -‐ 1
Lotus leaf image and SEM micrographs of its upper side showing the hierarchical structure.
[47]Compared to organic materials, inorganic materials exhibit better light, thermal, and chemical stabilities.
By transferring the structure of selected plant surfaces to practical materials, superhydrophobic surfaces could be manufactured and hence have recently been the focus of considerable scientific interest.
[48-‐51]*
A surface is called superhydrophobic when water contact angle exceeds 150
°.
This is due to the fact that artificial superhydrophobic surfaces are promising candidates for a number of practical applications such as easy cleaning clothing or windows.
Indeed, surface cleaning of building materials like facades and glass panes generates considerable trouble, high consumption of energy and chemical detergents and, consequently, high costs. Non-‐wettable surfaces may also improve the ability to prevent frost from forming or adhering to the surface.
The adhesion of water, as well as contaminants is considerably reduced. Water droplets coming in contact with a superhydrophobic surface form nearly spherical beads. The contaminants, either inorganic or organic, on such surfaces are picked up by water droplets or adhere to the water droplet and are removed from the surface when the water droplets roll off (Figure V -‐ 2). The combination of low surface energy
†and micro-‐
and/or nano-‐structured features, which can certainly reduce the contact area between the surface and water droplets, form superhydrophobic surfaces.
Figure V -‐ 2
On a typical surface a drop of water slides across and leaves most dirt particles sticking to the object, while on a superhydrophobic surface the drop rolls across, picking up dirt and carrying it away.
[52]On most surfaces, the motion of the drop is opposed by energy barriers, which leads to an effect referred to as contact angle hysteresis. Indeed, there is usually a difference between the angle produced as the volume of the drop is increased (advancing angle) and that when the volume is reduced (receding angle). This hysteresis gives a measure of the surface stickiness.
†
The surface energy is a solid surface characteristic associated with the molecular forces of its interaction with another material.
Surface tension is the property of a liquid arising from unbalanced molecular forces at or near the surface. If it is higher than the surface energy of a material, the liquid tends to form droplets rather than spread out or “wet out” as some refer to it. Surface tension is normally measured in energy units called dynes/cm. A dyne is the force required to produce an acceleration of 1 cm/s
2on a mass of 1g.
A drop with a very low hysteresis easily rolls off an inclined surface. Figure V -‐ 3 presents two different methods for the measurement of contact angle hysteresis.
Figure V -‐ 3
Methods to measure the contact angle hysteresis (θ
advancing-‐ θ
receiding)
[53](a) The tilting plate method captures the contact angles measurements on both the left and right sides of a sessile drop while the solid surface is being inclined typically from 0°
to 90°.
(b) The add and remove volume method requires adding (or removing) liquid to the maximum volume permitted without inducing a motion of the three phase (3φ) contact line (liquid-‐vapor-‐solid).
The hysteresis characterizes the topography of the sample. Water droplets on rugged hydrophobic surfaces typically exhibit one of the following two states (Figure V -‐ 4): (a) the Wenzel state
[54]in which water droplets are in full contact with the rugged surface or (b) the Cassie state
[55]in which water droplets no longer penetrate, but rest on top of the roughness features.
Figure V -‐ 4
Two different wetting state for a drop on a hydrophobic textured surface. (a) Wenzel state and (b) Cassie state or “fakir” state.
[49]The Wenzel state is characterized by a huge hysteresis (from 50° to 100°) compared to the Cassie state (from 5° to 20°), and hence displays a more sticky behavior as contact area with the surface is bigger.
In summary, the wettability of solid surfaces is therefore a very important property, and may be governed by both the chemical composition and the geometrical structure of the surface.
[50, 56]!"#$ !"#$%&'(#)*+',+*-./' !%#$ 0//')%/'1+,.2+'3.#4,+'5+*-./!
!"$#$%&'#&()*%+((
!"#$ !%#$
Techniques to make superhydrophobic surfaces can be simply divided into two categories: making a rough surface from a low surface energy material and modifying a rough surface with a material of low surface energy.
In our study, we evaluated the impact of the templating of the ZnO nanowire arrays on the wetting properties, focusing on roughening of the material. We have also studied the reversibility of the surface wettability, which is crucial for real device applications.
2. Experimental part
The two preparation routes are schematically depicted in Figure V -‐ 5.
Figure V -‐ 5
Two nanosphere lithography routes proposed for the growth of ZnO nanowire arrays.
Templated growth process (TG) & Templated seeding process (TS)
In the “Templated-‐Growth” TG procedure (Figure V -‐ 5, left side), a continuous layer of ZnO seeds is obtained by spin coating a 1:1 solution of zinc acetate and ethanolamine in ethanol, followed by annealing at 350°C for 30 min to obtain oriented ZnO seeds (step 1).
!
The colloidal crystal mask is then deposited (step 2) by spin coating a suspension of monodisperse polystyrene nanospheres (490 nm diameter). Next, we proceeded to the hydrothermal growth of the nanowires (step 3) in aqueous suspensions containing zinc nitrate hydrate (25 mM) and hexamethyltetramine (25 mM) at 90°C for 2.5 hours. We finally removed the polystyrene nanospheres (step 4) by calcination (350 °C for 30 min).
The samples before and after removal of the nanospheres will be further referred to as TG-‐NS and TG-‐calcined.
In the “Templated Seeding” TS procedure (Figure V -‐ 5, right side), the seeding and the colloidal crystal mask were realized with the same experimental conditions as in the TG procedure but in a different order. In the first step, a patterned layer of seeds was obtained by formation of the colloidal crystal mask on the substrate (step 1*), followed by the spin coating of the ZnO seeds over the nanospheres (step 2*). We then removed the nanospheres (step 3*) by calcination (TS-‐calcined) or by sonication+calcination (TS-‐
sonicated). The last step is the hydrothermal growth of the ZnO nanowires (step 4*).
2.1 Synthesis of the samples
All samples were manufactured on FTO conducting glass substrates (15Ω/sq) purchased from Dyesol. Previous to use, the substrates were washed by sonication 15 min in acetone and 15 min in ethanol and air-‐dried.
In both the templated seeding (TS) and templated growth (TG) routes, the seeding and the colloidal crystal masks were realized in the same way.
Monodisperse polystyrene nanospheres with a mean diameter of 490 nm were purchased from Bangs Laboratories as suspensions in water (concentration of about 10
% wt) in order to prepare single-‐layer colloidal crystal masks.
Before deposition, we diluted the aqueous nanospheres (150 μL) suspension in a surfactant Triton X-‐100/MeOH mixture 1:400 by volume (350 μL) and filtered through centrifugal filter units (porosity 0.65 µm) in order to eliminate aggregates of higher dimensions, which would disturb the formation of the monolayer. Then, we vortexed the suspensions during 2 minutes to ensure homogeneity.
We dispensed a drop of the suspension (50 μL) on the substrate with an Eppendorf pipette. All the samples were prepared with the same spin-‐coating parameters, characterized by a high rotation speed reached in a short time. The samples were accelerated to 1930 rpm (acceleration rate 643 rpm/s) for 2s and then spun at 2500 rpm (acceleration rate 190 rpm/s) for another 2s, followed by final spin at 6900 rpm (acceleration rate 1467 rpm/s) for 2s.
We seeded the samples with an oriented ZnO thin film by a sol-‐gel spin coating method.
We prepared the seeding solution by dissolving zinc acetate Zn(CH
3COO)
2.2H
2O (0.025
mol) in ethanol (100 mL) at room temperature. Ethanolamine (C
2H
7NO) was used as
stabilization agent and its molar ratio to zinc acetate was kept at 1:1. The resultant
solution was stirred at 60°C for 1h to yield a clear and homogeneous sol. Then, the mixed sol was aged at room temperature for 24h. The precursor solution was dropped onto the samples, which were then spinned at 3000 rpm for 20 s at room temperature in a constant humidity atmosphere. The relative humidity level was set to 35 ± 2 % and was measured with a digital hygrometer (Testo). To ensure complete sample coverage with ZnO seeds, the spin coating process was repeated nine times.
Seeded substrates were annealed at 350 °C for 30 min to obtain the oriented ZnO seed layers. In the TS-‐sonicated samples, we previously removed the nanospheres (step 3* in Figure V -‐ 5) by sonication before this calcination step while in the TS-‐calcined samples, the calcination played the dual role of removing the nanospheres and orienting the seeds.
The growth of the nanowires (step 3 or 4* in Figure V -‐ 5) was carried out by immersing seeded substrates upside down in aqueous solutions containing zinc nitrate hydrate (25 mM) and hexamethyltetramine (25 mM) at 90°C for 2.5 hours. The arrays were then rinsed with milliQ water and dried in oven in air at 60°C for 30 min. To remove the template of polystyrene nanospheres, we calcined the TG samples (step 4 in Figure V -‐
5) in air at 350°C during 30 min.
2.2 Morphological & crystallographic characterizations
The morphology of the samples was characterized by scanning electron microscopy (SEM) on a FEG-‐ESEM XL30 (FEI) with an accelerating voltage of 5 kV under high vacuum. All samples were gold-‐coated (60s) before observation.
The XRD patterns were recorded using a Bruker AXS D8 diffractometer in θ–2θ locked coupled mode (30-‐70 °2θ, step size 0.04°).
2.3 Dye loading
The dye loading was measured by UV-‐vis spectroscopy on EtOH/water (1/1 v/v) solutions with a Perkin Elmer UV-‐vis Spectrometer Lambda 14 P. The samples were soaked in N-‐719 (Solaronic) dye solutions (ethanolic solution, 3.0 10
-‐4M) during 16 h at room temperature.
After drying, the samples were desorbed in a known volume of KOH solution (10
-‐3M).
A calibration curve was used to calculate the experimental extinction coefficient (at 500 nm) of N-‐719 dye (12 500 (mol/L)
-‐1cm
-‐1).
The dye loading of the samples was then determined from the desorption solutions.
2.4 Wetting properties
Surface wettability was evaluated by ultra-‐pure water (MilliQ) contact angle (CA) measurements in the GRASP laboratory, using a CAM 200 Optical Contact Angle meter (KSV Instruments Ltd.) and the CAM 200 software provided with the instrument.
A 5 μL MilliQ water droplet was deposited on the surface of the as-‐manufactured samples using an automated drop dispensing system. Each CA measurement was repeated three times at various places on two different samples.
Light-‐induced hydrophilicity was evaluated by irradiating the samples during 2h inside a specially designed chamber equipped with six Eversun UVA fluorescent lamps (OSRAM, L40W/79K).
After each irradiation, a 5 μL water drop was placed on the sample and the corresponding CA was measured again.
In order to study the reversibility of the wettability transition, the samples were either stored in the dark for a week at room temperature or annealed during 20h at 50°C under O
2atmosphere before another CA measurement.
The contact angle hysteresis, defined as the difference between the advancing and receding angle, was also measured.
3. Results and discussion
A critical point in nanosphere lithography is that the experimental conditions of the colloidal crystal mask formation should ensure the presence of large covered areas as well as the accessibility of the voids, which is of crucial importance both for the growth of the nanowires in the TG procedure and for the deposition of the seeds onto the substrate in the TS procedure.
Electron micrographs of the colloidal crystal mask prior to ZnO hydrothermal growth are shown in Figure V -‐ 6. The low magnification micrograph proves that the nanospheres are present over several mm
2. The inset shows the hexagonal close packing (hcp) of the polystyrene nanosphere monolayer.
Figure V -‐ 6
SEM micrograph of a monolayer mask of polystyrene nanospheres (490 nm diameter),
prepared by spin coating with a high coverage rate. The inset is a high magnification view
showing the hexagonal packing of the monodisperse nanospheres.
3.1 Morphological and crystallographic properties
Scanning electron micrographs of the ZnO nanowires grown via the various synthetic routes are presented in Figure V -‐ 7. In the TG procedure, the ZnO nanowires grew in the voids between the nanospheres with hexagonal arrangement. The "side view"
micrograph (Figure V -‐ 7 (b)) highlights the role of the colloidal crystal mask in patterning the growth of the nanowires. After calcination at 350°C to remove the polystyrene mask, the hexagonal pattern was preserved (Figure V -‐ 7 (c)), as well as the orientation of the nanowires (Figure V -‐ 7 (d)).
In the TS route, the polystyrene masks were used to selectively deposit the seeds on the substrate and removed either by calcination or sonication, before the hydrothermal growth. At first sight (Figure V -‐ 7 (e) and Figure V -‐ 7 (g)), the coverage of the substrate by nanowires seems to be higher than in the TG route. However, the nanowires are not well-‐aligned anymore: the nanowires display a divergent, bush-‐like structure (see side-‐
view micrographs in Figure V -‐ 7 (f) and Figure V -‐ 7 (h)).
The dimensions of the nanowires manufactured are homogenous and similar for both procedures (TG – TS), with typical diameter of ≈ 50 nm and height ≈ 1 μm. This demonstrates the high reproducibility of the growth process and the high quality (coverage and ordering) of the colloidal masks.
In contrast, Fragala et al.
[37]observed a bimodal morphology of ZnO nanowires that was attributed to a lack of covering by polystyrene spheres (1 μm diameter). They reported that ZnO grown on a homogenous seed layer were larger (200 nm) than the nanowires in the patterned regions (80 nm).
Besides, Pyun et al.
[35]underlined a difference in growth rates for ZnO nanowires grown with and without polystyrene colloids. They showed that ZnO nanowires enclosing the polystyrene colloids are much longer than nanowires formed on the area not covered with the nanospheres.
As we synthesized nanowires for short period (2h30) we didn’t observe this difference of length between templated and unpatterned nanowire arrays. For equivalent solution and growth time (2h30), our nanowires are longer (1 µm) than that obtained by Pyun et al.
[35](750 nm). The higher temperature (90°C) of hydrothermal synthesis could explain the enhanced growth rate along [001] direction of our ZnO nanowire arrays.
[22]Figure V -‐ 8 shows the X-‐ray diffraction (XRD) patterns of films prepared by the TG and TS synthetic routes. All reflections belong either to the wurtzite-‐type phase (ZnO – ICDD PDF4+ No. 04-‐003-‐02106) or to the fluorine-‐doped tin oxide layer of the substrate (SnO
2– ICDD PDF4+ No. 00-‐046-‐1088).
The prevailing intensity of the ZnO (002) peak evidences the c-‐axis texturation normal
to the substrate. This result confirms the role of the seed layer to promote the alignment
of ZnO nanowire arrays, as already reported for unpatterned ZnO nanowire arrays.
[22, 57]Figure V -‐ 7
SEM micrographs of ZnO nanowires manufactured via the “templated growth” (TG) and the
“template seeding” (TS) routes. (a) & (b) TG-‐NS : (a) As-‐grown nanowires highlighting a hexagonal symmetry. (b) The polystyrene nanospheres are still present between the nanowires in the side-‐view. (c) & (d) TG calcined : (c) The nanospheres have been removed by the annealing treatment and left a well-‐defined hole where they were sitting. (d) Side-‐
view of the nanowires with no trace of the nanospheres. Inset scale bars are 500 nm. Before the hydrothermal growth of the nanowires in the TS process, the nanospheres were removed either by calcination (TS -‐ calcined) at 350 °C ((e) & (f)) or by sonication (TS -‐ sonicated) in toluene ((g) & (h)). Tilted SEM micrographs ((f) & (h)) show the bush like orientation of the nanowires.
!
Figure V -‐ 8
XRD patterns of the samples grown via the “templated growth” (TG) or “templated seeding” (TS) routes.
In order to quantify the c-‐axis texturation, the ratios of (002) to (103)
‡peak net surface areas are given in Table V -‐ 1. The same pronounced c-‐axis texturation is found in the TG samples and in a reference unpatterned (=continuous) nanowire film. As predicted from the SEM micrographs, the c-‐axis texturation is far less pronounced in samples manufactured by the TS route, especially when the nanospheres were removed by sonication (TS-‐sonicated).
Table V -‐ 1
Comparison between XRD (002) textures of TG, TS and unpattenerd nanowire arrays
Sample Ratio of XRD peak
surface areas (002)/(103)
TG calcined 54
TG -‐ NS 51
TS calcined 12
TS sonicated 8
Unpaterned continuous
Nanowire Film 52
‡
In the hexagonal lattice (such as wurtzite structure), sometimes the four Bravais-‐Miller
indices (h k i l) are used instead of the three Miller indices (h k l). In the 4-‐index notation
the i index is related to the h and k indices by the relation h + k + i = 0.
The crystalline domain size calculated from the (002) ZnO peaks gave an average domain size of ≈ 45 nm, which agrees with our SEM observations of the nanowire diameter. This is expected since hydrothermal growth usually yields single crystalline nanorods.
[58]3.2 Surface area
The performance of such nanostructures for applications often depends critically on the accessible surface area. In order to assess this parameter, we performed dye loading measurements, after verifying that the N719 dye does not adsorb on the polystyrene nanospheres.
Dye loading results are listed in Table V -‐ 2 for the samples prepared by the TS and TG routes and for a reference unpatterned nanowire array synthesized in the same hydrothermal growth conditions.
The two samples obtained by the TS route (TS-‐calcined and TS-‐sonicated) have dye loading values significantly lower than the samples prepared by the TG route. This observation points to a lower density of nanowires (number of nanowires per area unit of substrate) due to a less efficient seeding in the TS route. The lowest value for the TS-‐
sonicated sample suggests that some seeds are removed from the substrate during the sonication step.
Table V -‐ 2
Dye loading values of TG, TS and unpattenerd nanowire arrays
Sample Dye loading N-‐719*
[mol/mm
2]
TG calcined 2.43 10
-‐10TG -‐ NS 1.88 10
-‐10TS calcined 1.60 10
-‐10TS sonicated 1.35 10
-‐10Unpaterned Nanowire Film 1.68 10
-‐10* Dye loading values are expressed in moles of N-‐719 dye per mm
2of soaked area.
Regarding the TG samples, the dye loading of the TG-‐calcined sample is higher than the
value for the TG-‐NS sample, because of the removal of the polystyrene nanospheres that
were "screening" the bottom of the nanowires. Moreover, the calcinations may also have
removed organic contaminants, therefore favoring the dye loading. The dye loading
value obtained for the TG-‐calcined sample also significantly exceeds the value for the
unpatterned nanowire array.
Since it is obviously unlikely that a patterned array may have a higher nanowire density than an unpatterned array, the reason for this difference in dye loading is to be found in a difference of accessible surface area. Unlike Chang et al.
[36], both the patterned and unpatterned ZnO samples were grown in the same conditions of seeding and hydrothermal growth. Therefore, we can unambiguously conclude that the templated growth (TG) route significantly contributed to an increase of the accessible surface area, which was the ultimate goal for using nanosphere lithography in this study.
In order to estimate the nanowire density of the arrays obtained by the TG route, we considered a perfect colloidal mask where each nanosphere is surrounded by six interstices (inset Figure V -‐ 6). In the case of only one nanowire per interstice, the density would be of 9.6 10
12nanowires/m
2for a colloidal mask with 490 nm diameter nanospheres. This scenario corresponds to samples manufactured by nanosphere lithography and vapor-‐liquid-‐solid (VLS) epitaxy mechanism catalyzed by Au nanodots.
[59]In our process, electron micrographs (Figure V -‐ 7) show that several nanowires grow between the nanospheres, leading to an increase in the total surface area. Based on geometric calculations,
[60]it is possible to evaluate the size of the triangular interstice in the colloidal mask (inset Figure V -‐ 6). Given the observed diameter of the nanowires, we estimate that four nanowires can take place in the voids from a colloidal mask with 490 nm diameter nanospheres. This would raise the density up to 3.8 10
13nanowires/m
2. Such a value is in the same range than reported values for track etched polymer templates designed for micro-‐ and nanofabrication.
[61]However, the pores in the polymer membranes are randomly distributed, which is not the case in the nanosphere lithography process.
By changing the size of the nanospheres, it is possible to tune the density of nanowires, which may influence the efficiency of the applications, as previously mentioned. For example, in the field of dye-‐sensitized solar cells, the fact that electron diffusion length is in the range of 20 nm suggests that the best well-‐aligned nanostructure should be composed of vertical nanostructures of about 20-‐40 nm diameter.
[15]However, the size of the interstice will decrease with the nanosphere diameter and a too small space could hinder the growth of the nanowires.
[35, 36]We estimate that the critical interstice size for the growth of one nanowire would be reached with ≈ 215 nm diameter nanospheres.
Therefore the adjustment of the interstice size is of major importance to have control
over the manufacturing of NSL-‐patterned ZnO nanowire arrays. The use of an oxygen
plasma etching treatment has been reported in literature for the reduction of the
polystyrene nanosphere diameters
[62]or correction of the deformation of soft
polystyrene nanospheres.
[63]This etching of the mask would be necessary if nanospheres with a diameter < 215nm are envisaged to increase the packing density of interstices and nanowires.
3.3 Wetting properties
To evaluate the patterning effect of the ZnO nanowire arrays on their wetting properties, we measured water contact angles on the as-‐prepared, unpatterned (Figure V -‐ 9 (a)) and TG (templated growth) ZnO nanowires (Figure V -‐ 9 (b)), which present higher surface area compared with TS samples (Table V -‐ 2).
Both unpatterned and patterned samples revealed a hydrophobic behavior with water contact angles (CA) higher than 90°.
The patterning induced by the templated growth resulted in a significant increase of the contact angle, almost reaching the superhydrophobic limit (112° ± 2 à 132° ± 2). The CA of the unppatterned array agrees with recently reported studies
[45, 64]on ZnO nanowires/nanorods, keeping in mind that it is difficult to meet exactly the same synthesis conditions. Up to now, no CA measurement was ever reported on ZnO nanowire arrays manufactured with the templated growth process.
It is well known that the surface free energy and the surface roughness play a very important role in the wetting properties. Recently, it has also been reported that the wettability of ZnO depends on the surface crystal structure.
[65-‐67]Figure V -‐ 9
Reversible surface wettability transition of ZnO nanowire arrays grown manufactured by (a) conventional unpatterned hydrothermal growth and (b) templated growth route.
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We will consider all these concepts to understand the hydrophobic behavior of the as-‐
prepared samples and the increased hydrophobic character of the patterned sample.
The anisotropic nanowire growth evidenced in the XRD patterns (Figure V -‐ 8) is due to different surface free energies of the various growing crystallographic planes.
[68]A fast growing plane generally tends to disappear leaving behind slower growing planes with lower surface energy.
[38]For the anisotropic ZnO nanowire, the velocities of crystal growth in various directions were reported to be 100 > 101 > 001 ≈ 001 .
[69]Therefore, compared with a sample with a random orientation, the as-‐prepared, unpatterned and TG patterned samples would have the lowest surface free energy, which can be explained by a closer look at the crystal structure (Figure V -‐ 10).
Zinc oxide has wurtzite type symmetry and thus belongs to the space group C
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On the non-‐polar planes (e.g. (110) face), both oxygen and zinc ions are terminated in the same plane. On the other hand, there are two possibilities of surface-‐terminated ions on the polar (001) planes, that is, oxygen ions or zinc ions. Miyauchi et al.
[65]calculated that the total energy of the oxygen-‐terminated surface was higher than that of the zinc-‐
terminated surface. These results imply that the zinc ions terminated structure is more energetically stable than the oxygen ions termination. Under this stable structure, oxygen ions are not exposed at the surface, which is not the case in the non-‐polar planes.
As surface oxygen ions are considered to act as reactive sites for increasing OH species on the surface, it implies that the hydrophilicizing rate of ZnO nanowires with high proportion of nonpolar planes is faster than those with low proportion.
However, as we did not evidence any significant difference between the XRD patterns of both samples, the high contact angle measured on the TG patterned sample cannot be attributed to a lower surface energy.
Figure V -‐ 10
Schematic illustration of the atomic alignments on ZnO (110) and (001) crystal faces.
Inspired from ref [65].
We therefore suggest that the variation of the initial contact angle among samples should be attributed to the difference in their surface roughness.
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As a reminder, the contact angle hysteresis can give information about the topography of the sample and reflects the state of hydrophobicity (Wenzel or Cassie state Figure V -‐
11 (a)).
Both samples revealed high hysteresis behavior (Figure V -‐ 11 (b)), which is a characteristic feature of the Wenzel state.
According to the Wenzel model, the water contact angle θ
W, is given by the following equation
[54]:
cos !
!= ! cos ! (Equation V -‐ 1)
where r (roughness factor) is the ratio of the unfolded surface to the apparent surface under the droplet and θ is the contact angle on a flat surface of the same nature as the rough one. Since r is always greater than unity, this model predicts an increase of the contact angle with the surface roughness.
From Figure V -‐ 9, it seems evident that the TG patterning process increased the roughness of the ZnO nanowire array compared to the unpatterned process, which is denser. These results agree with the study performed by Das et al.
[45]on ZnO nanoneedles and nanorods manufactured by metal organic chemical vapor deposition.
Due to the large size of the samples, we measured the roughness with a portable surface roughness tester and evidenced a subsequent increase (factor 2-‐3) of the surface roughness due to the TG patterning process. However, a precise knowledge of the substrate roughness is difficult to measure due to the finite size (radius and aspect ratio) of the measuring tips.
Figure V -‐ 11
Contact angle hysteresis
(a) Characteristic shape of receiding drop in Wenzel and Cassie states.
[70](b) Contact angle hysteresis measurements on unpatterned and TG patterned samples !
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