Mirror-twin induced bicrystalline InAs nanoleaves

Nano Res

1

Mirror-twin induced bicrystalline InAs nanoleaves Mun Teng Soo1, Kun Zheng2,3 (*), Qiang Gao4, Hark Hoe Tan4, Chennupati Jagadish4, and Jin Zou1,2 (*) Nano Res., Just Accepted Manuscript • DOI: 10.1007/s12274-015-0955-z

http://www.thenanoresearch.com on November. 25, 2015

© Tsinghua University Press 2015

Just Accepted

This is a “Just Accepted” manuscript, which has been examined by the peer-review process and has been accepted for publication. A “Just Accepted” manuscript is published online shortly after its acceptance, which is prior to technical editing and formatting and author proofing. Tsinghua University Press (TUP) provides “Just Accepted” as an optional and free service which allows authors to make their results available to the research community as soon as possible after acceptance. After a manuscript has been technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Please note that technical editing may introduce minor changes to the manuscript text and/or graphics which may affect the content, and all legal disclaimers that apply to the journal pertain. In no event shall TUP be held responsible for errors or consequences arising from the use of any information contained in these “Just Accepted” manuscripts. To cite this manuscript please use its Digital Object Identifier (DOI®), which is identical for all formats of publication.

Nano Research DOI 10.1007/s12274-015-0955-z

Mirror-twin induced bicrystalline InAs nanoleaves

M. Teng Soo1, Kun Zheng2,3,*, Qiang Gao4, H. Hoe

Tan4, Chennupati Jagadish4 and Jin Zou1,2,*

1 Materials Engineering, 2 Centre for Microscopy and

Microanalysis, and 3 Australian Institute for

Bioengineering and Nanotechnology, The University

of Queensland, St. Lucia, QLD 4072, Australia 4 Department of Electronic Materials Engineering,

Research School of Physics and Engineering, The

Australian National University, Canberra, ACT 2601,

Australia

Biomimic of semiconductor nanoleaf induced by mirror-twin.


M. Teng Soo1, Kun Zheng2,3 (*), Qiang Gao4, H. Hoe Tan4, Chennupati Jagadish4 and Jin Zou1,2 (*)

1 Materials Engineering, The University of Queensland, St. Lucia, QLD 4072, Australia 2 Centre for Microscopy and Microanalysis, The University of Queensland, St. Lucia, QLD 4072, Australia 3 Australian Institute for Bioengineering and Nanotechnology, The University of Queensland, St. Lucia, QLD 4072, Australia 4 Department of Electronic Materials Engineering, Research School of Physics and Engineering, The Australian National University,

Canberra, ACT 2601, Australia

Received: day month year Revised: day month year Accepted: day month year (automatically inserted by the publisher)

© Tsinghua University Press and Springer-Verlag Berlin Heidelberg 2014

KEYWORDS InAs, nanoleaf, twin boundary, mirror twin

ABSTRACT In this study, leaf-like one-dimensional InAs nanostructures were grown by metal-organic chemical vapor deposition method. Detailed structural characterization suggests the nanoleaves contain relatively low-energy or mirror twins acting as their midribs and narrow sections connecting the nanoleaves and their underlying bases as petioles. Importantly, the mirror twins lead to identical lateral growth of the twined structures in terms of crystallography and polarity, which is essential for the formation of lateral symmetrical nanoleaves. It has been found that the formation of nanoleaves is driven by the catalyst energy minimization. This study provides a biomimic of semiconductor nanoleaf with that of real leaf in nature.

1. Introduction

In the thriving field of nanoscience, a major ambition is to synthesize nanoscale building

blocks of arbitrary dimensions, morphologies, and materials with increasing complexity [1].

Biomimicry is an interdisciplinary study which looks to nature and natural systems for inspiration in practical engineering design [2]. Discovering a synthetic pathway to artificial

analogs of materials in nature represents a fundamental milestone in the development of

Nano Research DOI (automatically inserted by the publisher)

Address correspondence to Jin Zou, [email protected]; Kun Zheng, [email protected]

Research Article

www.theNanoResearch.com∣www.Springer.com/journal/12274 | Nano Research

2 Nano Res.

nanomaterials. Studying and replicating these biologically inspired nanostructures may lead to new knowledge for fabrication and design of new and novel nano-scale devices, as well as providing valuable insight in to how such phenomena can be exploited [3, 4].

III-V semiconductor nanostructures are promising low-dimensional materials with a wide range of applications, including light-emission, logic, sensing and solar-cell technology [5-12]. As a main class of III-V semiconductors, InAs one-dimensional (1D) nanostructures have attracted significant research attention due to their direct and narrow bandgap, relatively high electron mobility and small electron effective mass [13, 14], which have made them a potential candidate for applications in resonant tunneling diodes, Josephson junctions and high-performance transistors [15-17].

Nanoleaf is a type of leaf-like 1D nanostructure, and it was reported mainly for metal oxides [18-23], and a few III-V materials [24, 25]. In this study, we demonstrated Au-catalyzed growth of InAs nanoleaves. The structure of these nanoleaves has been investigated carefully and we found that the axial twin boundary, which acts similar to the midrib of a real leaf in nature, is necessary in the formation of these nanoleaves. A new type of twins is identified and their nature is confirmed. The growth pattern of these nanoleaves is found similar to the growth of dorsiventral leaves in nature [26-28]. Therefore, we believe that by correlating both of the structure of InAs nanoleaves in this nanoworld and the structure of real dorsiventral leaves in our natural world, a bio-inspired nano-technology design can be motivated for future electronic or optoelectronic devices.

2. Experimental

InAs nanoleaves were grown via horizontal flow metal-organic chemical vapour deposition (MOCVD) using trimethylindium (TMIn) as the group III precursor and arsine (AsH3) as the group V precursor, and Au nanoparticles to drive nanowire growth. InAs substrates were treated with poly-L-lysine (PLL) solution followed by dispersion of a solution containing 50-nm colloidal Au nanoparticles. Nanoleaves

were grown at a pressure of 100 mbar and a total gas flow rate of 15 slm (slm = standard litre per minute). Prior to growth, each substrate deposited with the Au particles was annealed in situ at 700 °C under AsH3 ambient to desorb surface contaminants. It has been well documented that at high annealing temperature under AsH3 ambient, Au colloidal nanoparticles can first react with the InAs substrate to form Au-In alloy droplets [29-31]. AsH3 is typically introduced in the chamber during temperature ramp-up and annealing to counteract As sublimation and preserve the surface stoichiometry of the InAs substrate. After the substrate was cooled to the growth temperature of 500 °C, the AsH3 flow was adjusted and TMIn introduced to initiate nanowire growth. The growth time was 30 min. The absolute flow rate of TMIn and AsH3 was controlled at 1.2 x 10-5 mol/min and 3.4 x 10-5 mol/min, respectively, to attain a very low V/III ratio of 2.9 so that a growth environment in the excess of In can be maintained for promoting the growth of non-vertical 1D nanostructures [32].

Detailed morphology, structure and composition of the InAs nanoleaves were characterized by field-emission scanning electron microscopy (FE-SEM – JEOL JSM-7800F, operated at 5 kV) and transmission electron microscopy (TEM – FEI Tecnai F20, operated at 200 kV; and FEI Tecnai T12, operated at 120 kV). TEM specimens were prepared by dispersing the nanoleaves in ethanol using an ultrasonic bath for 15 min and then spreading the drops from the suspension onto holey carbon grids. FE-SEM analysis was used to identify the general morphology of the nanoleaves such as facet planes and size. Conventional bright-field and dark-field TEM imaging of the nanoleaves along the zone axis was used to identify the crystal structures and twin boundary; in combination with selected-area electron-diffraction (SAED) and lattice imaging. Nanoleaves were screened for twin boundary over their entire length. The composition of the Au-In alloy catalyst was studied by energy-dispersive X-ray spectroscopy (EDS) analysis. Convergent beam electron diffraction (CBED) was employed to determine the polarity of twin-sections of the nanoleaves through


3Nano Res.

correlating the CBED patterns simulated using the Bloch wave method in JEMS software [33].

3. Results and discussion

Figure 1(a) is a plan-view SEM image and shows the overview of Au-catalyzed InAs nanostructures grown on an InAs substrate, in which inclined leaf-shaped nanostructures with a density of over 1% can be seen (as arrowed). Figure 1(b) is a magnified SEM image of such a nanoleaf. Figures

Figure 1 (a) Top-view SEM image showing the overview of the Au-catalyzed InAs nanostructures, in which the red arrows point to the two different nanoleaves. (b) Top-view image of a single inclined InAs nanoleaf. Enlarged SEM images focused at the (c) tip (showing the asymmetrical catalyst as indicated by red arrow); and (d) base (showing the twin boundary located in the middle of nanoleaf starting from the base indicated by red arrow) of the nanoleaf. 1(c) and 1(d) are a pair of enlarged SEM images to show the tip and base of the nanoleaf, respectively. As can be seen from Fig. 1(c), a catalyst is associated with the tip of the nanoleaf, which is asymmetrically associated with the nanoleaf. On the other hand, at the base region, a line (as marked by the red arrow) is associated with the axial direction of the nanoleaf that divides the nanoleaf, starting from the base. In fact, this line acts as a midrib and makes the nanostructure to look like a plant leaf with a midrib across in the middle. Our extensive SEM investigations indicate that these nanoleaves all have asymmetrical catalysts on their tips and have straight lines dividing these leaves. By carefully examining these nanoleaves, unlike those inclined nanowires [32, 34, 35], no specific projected directions can be found.

To understand the structural characteristics of these nanoleaves, TEM investigations were performed, in which individual nanoleaves were prepared lying down on holey carbon films. In fact, these individual nanoleaves can be characterized by both SEM and TEM. Figures 2(a) and 2(b) are the typical SEM and TEM images of individual nanoleaves. Since they are lying flat on the holey carbon films, these SEM and TEM images suggest that the nanostructure has a leaf-shape indeed. In particular, as shown in Fig. 2(b), a crystal boundary (as marked by the red arrow) is seen to divide the nanoleaf from the middle and across the entire leaf from the tip to the bottom, acting as a midrib in a


4 Nano Res.

Figure 2 (a) SEM and (b) TEM images of individual InAs nanoleaves showing their flat nature. Inset in (a) showing the asymmetrical catalyst. Enlarged TEM images from (b) at the nanoleaf tip (c) showing the asymmetrical catalyst and at the region away from the tip (d), both have a vertical boundary in the middle of the nanoleaf. (e) and (f) are SAED patterns taken from left and right of both regions shown in (c) indicating that they have zinc-blende crystal structure. plant leaf. It is of interest to note that, in both SEM and TEM images, the nanoleaves are electron transparent, indicating that they are sufficiently thin. Figures 2(c) and 2(d) are magnified TEM images showing the tip and top regions of the nanoleaf. Interestingly, the crystal boundary can be clearly seen in both cases. Furthermore, the asymmetrical catalysts can be seen (refer to Fig. 2(c)), which is echoed by the inset in Fig. 2(a). Figures 2(e) and 2(f) are two SAED patterns taken from both sides of the crystal boundary near the catalyst. These SAED patterns are identical but relatively oriented, and can be respectively indexed as the and

zone-axes of the zinc-blende crystal structure, suggesting that the nanoleaf has the

surface. To determine the nature of the crystal

boundary of our nanoleaves, we employed high-resolution TEM (HRTEM) and SAED investigations to study a large number of nanoleaves. Through our extensive TEM investigations, we found that there are two different kinds of crystal boundaries. Figures 3(a) and 3(b) are HRTEM images taken from two different cases and their corresponding SAED patterns are shown in Figs. 3(c) and 3(d). Based on both HRTEM images and SAED patterns, twin structures can be clearly identified. By carefully analyzing the SAED pattern shown in Fig. 3(c), it is found that both and diffraction spots are the common diffraction spots for both

twinned

Figure 3 (a) and (b) are typical HRTEM lattice images at the coherent crystal boundary of two distinct InAs

nanoleaves viewed along the and zone-axes, with (c) and (d) corresponding SAED patterns (blue: left region; orange: right region). In (c), common spots showing

the mutual twin boundary of plane and growth

direction of ; and while in (d), common spots showing

the mutual twin boundary of plane.

structures. Furthermore, by carefully correlating the SAED pattern with the HRTEM image in Figs. 3(c) and 3(a), the diffraction spots are perpendicular to the twin boundary shown in Fig. 3(a), suggesting that the twin boundary is a

twin. In addition, since the is parallel to the twin boundary and the twin boundary parallel to the axial direction of the nanoleaf (refer to Fig. (2)), the axial direction of this nanoleaf can be determined to be along the

direction. Similarly, by correlating the SAED pattern with the HRTEM image shown in Figs. 3(d) and 3(b), twin boundary shown in Fig. 3(b) can be determined as the twin. Based on the fact that the SAED pattern is taken from

zone-axis for both twinned structures, the axial direction of this nanoleaf must be perpendicular to the and , which gives rise to the direction according to crystallography.

Since there exists polarity in the zinc-blende


5Nano Res.

structure (leading to ), it is necessary to determine the nature of the polarity and consequently the nature of the twin boundaries.

Figure 4 (a) TEM image of a typical nanoleaf. (b) and (c) are experimental CBED patterns taken from the left and

right twined structures. (d) Simulated CBED pattern of a 50-nm thick InAs. (e) Atomic model of InAs

zinc-blende structure viewed along the zone-axis. (f) HRTEM lattice image and (g) atomic model of bicrystalline

InAs nanoleaf at the growth front viewed along the

zone-axes, which consists of a twin boundary. Accordingly, CBED was performed on several nanoleaves. Figure 4 shows an example. Figure 4(a) is a bright-field TEM image showing a nanoleaf with a twin boundary. Figures 4(b) and 4(c) are a pair of zone-axis CBED patterns taken from both twinned structures, respectively. As can be seen from individual CBED pattern, there exists a mirror-symmetry, as indicated by their respective dashed lines; and the contrasts of the two diffraction disks are different, reflecting the polarity. To determine the absolute polarity, the theoretical CBED patterns were simulated using the JEMS software [33]. Figure 4(d) is a simulated InAs [011] zone-axis CBED pattern (where the following parameters were used for simulation: TEM specimen thickness is 50 nm; incident beam convergence half-angle is 6.707 mrad), in which

the contrasts in both diffraction disks are indeed different and there is a vertical mirror symmetry in the simulated CBED pattern. In fact, such a structural characteristic is demonstrated in Fig. 4(e) that present the atomic model of zinc-blende structured InAs. By correlating the

contrasts obtained from experiments (Figs. 4(b) and 4(c)) with the simulation results (Fig. 4(d)), both experimental diffraction disks can be unambiguously distinguished, as marked in Figs. 4(b) and 4(c). Accordingly, the polarities of the both twinned structures are determined. The comparison of Figs. 4(b) and 4(c) indicates that two CBED patterns are identical and symmetrical with the twin plane, suggesting that the observed twin belongs to the mirror twin, so that they have the identical growth direction, which is the essential component for the formation of the nanoleaf.

To understand the nature of interfaces between the catalyst and twin-structured nanoleaf, we carefully investigate their crystallographic relationships. Figure 4(f) is a

on-zone TEM image taken from the tip region of the nanoleaf shown in Fig. 4(a), in which sharp interfaces between the catalyst and two twin-structures can be clearly seen. By carefully analyzing the CBED pattern shown in Fig. 4(b) and correlating it with the TEM image shown in Fig. 4(f), the interface between the catalyst and the left twinned structure is determined to be . Using similar analysis, the interface between the catalyst and right twinned structure is determined to be . According to crystallography, both and

interfaces are A-type interfaces, i.e. In terminated interface. Figure 4(g) is the atomic model representing the crystallographic relationship between the catalyst and the two twinned structures, in which both interfaces are In terminated interfaces. Furthermore, according to the crystallographic requirement for the twinned nanoleaf, the angle between these two

interfaces is 141°, which matches well with our observation (refer to Fig. 4(f)).

Based on our extensive TEM investigations, the surface of the nanoleaves has planes with their midribs being either or twin boundaries. These twin boundaries are


6 Nano Res.

perpendicular to the leaf surfaces and to their axial directions. Nevertheless, the interfaces between the catalysts and twinned structures of nanoleaves are all interfaces. These rigid crystallographic relationships require specific angles (θ) between

Figure 5 Schematic illustrations at the growth front of

nanoleaves with (a) twin boundary, with two

interfaces between catalyst and nanoleaf; (b)

twin boundary, with two interfaces between

catalyst and nanoleaf; and (c) twin boundary, with two interfaces between the catalyst and the right twinned

structure: one is and the other one is . the two adjacent interfaces. For twinned nanoleaves, (Fig. 5(a)); while for

twinned nanoleaves, (Fig. 5(b)). Since , for a spherical catalyst being cut by the two planes of a nanoleaf shown in Figs. 5(a) and 5(b), the surface area of the catalyst is larger in the case of twinned nanoleaf (Fig. 5(b)), suggesting a higher catalyst surface energy, which may result in the system being unstable. Interestingly, by carefully studying the catalysts in twinned nanoleaves, we found that the interfaces between the catalysts and nanoleaves are not simple, as illustrated in Fig. 5(b). As shown in Fig. 2(c), there are two interfaces between the catalyst and the right twinned section: one is and the other one is with the latter one being parallel to the twin plane of the nanoleaf. The resultant catalyst/nanoleaf relationship is illustrated in Fig. 5(c). The comparison of the surface energies of the catalysts in the configurations shown in Figs. 5(a) and 5(c) indicates that the surface energy becomes lowered for the twinned nanoleaves if the dimensions of the two catalysts are the same. Furthermore, the interface has been observed in most of the studied twinned

nanoleaves, suggesting that such a configuration is energetically stable. Thus the observed interface between the catalyst and a twinned structure must have a relatively low interfacial energy; otherwise, the configuration shown in Fig. 5(b) would be preferred. Based on our detailed TEM investigations, we anticipate that the complicated interfacial configurations found in the twinned nanoleaves is due to the required minimization of overall catalyst surface/interfacial energies. On the other hand, the fact that we did not find such a complicated interfacial configuration in twinned nanoleaf indicates that the interface might have a relatively higher interfacial energy when such a configuration is adopted to lower the overall catalyst energy. We found that most nanoleaves are twinned nanoleaves, which is coincidence with the prediction that

twins have relatively low energy than twins [36, 37]. On this basis, we anticipate

that most twinned nanoleaves are attributed to their relatively low energy required for their formation. It is of interest to note that no

twinned nanoleaf is observed, possibly due to the high-energy requirement for forming

mirror twins caused by the polarity reversal [38, 39]. Based on these analyses and discussion, we anticipate that the extraordinary twin boundaries of or found in the InAs nanoleaves are formed due to the required formation of interfaces between Au catalysts and their induced InAs nanostructures, in which two interfaces are formed (driven by the catalyst energy minimization). This is very different from the common interface found between Au catalysts and InAs nanowires, in which only one interface are formed between the catalysts and their underlying nanowires leading to the nanowire growth [40]. In our case, once the mirror twins are formed, the lateral growth of the twinned nanostructure leads to formation of leaf-shaped nanostructure. Therefore, the fact that asymmetrical catalytic particle at the tip of a nanoleaf produce a symmetrical nanoleaf is the rigid crystallographic requirement. On other hand, since the surfaces have a relatively


7Nano Res.

low energy for InAs (although higher than the {111} surface energy), the formation of all nanoleaves with flat surfaces is driven by crystallographic necessity, rather than the absolute energy minimization.

According the different growth mechanisms proposed in the literature, our nanoleaves (containing twins as midribs) should be grown synergistically by three growth mechanisms, namely via vapor-liquid solid (VLS) [41], twin-plane re-entrant-edge (TPRE) [42, 43], and vapor-solid (VS) mechanisms [44]. Both VLS and TPRE growth mechanisms lead to the axial growth of nanoleaves whereas the VS growth mechanism regulates the lateral growth at sidewalls of nanoleaves. It is important to note that the midribs of our nanoleaves are mirror twins, leading to growth fronts of sidewalls of both twined structures are identical in terms of crystallography and polarity. This fact in turn enables the identical lateral growth of both twined structures in a nanoleaf, thus forming the symmetrical structure [45, 46]. In general, the lateral growth of one-dimensional nanostructures, such as nanowires, leads to the formation of tapered nanostructures when these nanostructures are directly grown from the underlying substrates [45, 47]. In our case, all nanoleaves were developed from bases with narrow sections between them and these narrow sections act as “petiole” (refer to Fig. 1(d)), which is critically essential for the formation of nanoleaves. The physical resemblances between our unique nanoleaf and real leaf are presented in the Electronic Supplementary Material.

4. Conclusion

In summary, bicrystalline InAs nanoleaves with and mirror twin boundaries were,

for the first time, grown under an excessive In condition. Detailed electron microscopy investigations were performed, from which the growth behavior and fundamental reasons attributed to the growth of nanoleaves are clarified. Specifically, the necessary conditions for formation of nanoleaves are the formation of mirror twins and the existence of a narrow connecting section between the nanoleaves and their underlying bases. These biomimetic

nanoleaves are inter-related with natural plant leaves so that some valuable hints for a bio-inspired nano-technology design may be inspired.

Acknowledgements

This research was supported by the Australian Research Council. The Australian National Fabrication Facility and Australian Microscopy & Microanalysis Research Facility, both established under the Australian Government’s National Collaborative Research Infrastructure Strategy, are gratefully acknowledged for proving access to the facilities used in this work.

Electronic Supplementary Material: Physical resemblances of a real leaf and nanoleaf. This material is available in the online version of this article at http://dx.doi.org/10.1007/s12274-***-****-* (automatically inserted by the publisher). References

[1] Bierman, M. J.; Lau, Y. K. A.; Kvit, A. V.; Schmitt, A. L.; Jin, S. Dislocation-Driven Nanowire Growth and Eshelby Twist. Science 2008, 320, 1060-1063.

[2] Vincent, J. F. V.; Bogatyreva, O. A.; Bogatyrev, N. R.; Bowyer, A.; Pahl, A.-K. Biomimetics: its practice and theory. J. R. Soc. Interface 2006, 3, 471-482.

[3] Han, B.; Huang, Y.; Li, R.; Peng, Q.; Luo, J.; Pei, K.; Herczynski, A.; Kempa, K.; Ren, Z.; Gao, J. Bio-inspired networks for optoelectronic applications. Nat. Commun. 2014, 5.

[4] Buhl, K.; Roth, Z.; Srinivasan, P.; Rumpf, R.; Johnson, E. Biologically inspired optics: analog semiconductor model of the beetle exoskeleton. In Proc. SPIE 7057; International Society for Optics and Photonics, 2008; pp 705707-705707-705708.

[5] Duan, X.; Huang, Y.; Cui, Y.; Wang, J.; Lieber, C. M. Indium phosphide nanowires as building blocks for nanoscale electronic and optoelectronic devices. Nature 2001, 409, 66-69.

[6] Huang, Y.; Duan, X.; Cui, Y.; Lauhon, L. J.; Kim, K.-H.; Lieber, C. M. Logic Gates and Computation from Assembled Nanowire Building Blocks. Science 2001, 294, 1313-1317.

[7] Zhong, Z.; Wang, D.; Cui, Y.; Bockrath, M. W.; Lieber, C. M. Nanowire Crossbar Arrays as Address Decoders for Integrated Nanosystems. Science 2003, 302, 1377-1379.


8 Nano Res.

[8] Parkinson, P.; Lloyd-Hughes, J.; Gao, Q.; Tan, H. H.; Jagadish, C.; Johnston, M. B.; Herz, L. M. Transient Terahertz Conductivity of GaAs Nanowires. Nano Lett. 2007, 7, 2162-2165.

[9] Joyce, H. J.; Gao, Q.; Hoe Tan, H.; Jagadish, C.; Kim, Y.; Zou, J.; Smith, L. M.; Jackson, H. E.; Yarrison-Rice, J. M.; Parkinson, P.; Johnston, M. B. III–V semiconductor nanowires for optoelectronic device applications. Prog. Quant. Electron. 2011, 35, 23-75.

[10] Xia, H.; Lu, Z.-Y.; Li, T.-X.; Parkinson, P.; Liao, Z.-M.; Liu, F.-H.; Lu, W.; Hu, W.-D.; Chen, P.-P.; Xu, H.-Y.; Zou, J.; Jagadish, C. Distinct Photocurrent Response of Individual GaAs Nanowires Induced by n-Type Doping. ACS Nano 2012, 6, 6005-6013.

[11] Saxena, D.; Mokkapati, S.; Parkinson, P.; Jiang, N.; Gao, Q.; Tan, H. H.; Jagadish, C. Optically pumped room-temperature GaAs nanowire lasers. Nat. Photon. 2013, 7, 963-968.

[12] Chau, R.; Doyle, B.; Datta, S.; Kavalieros, J.; Zhang, K. Integrated nanoelectronics for the future. Nat. Mater. 2007, 6, 810-812.

[13] Milnes, A. G.; Polyakov, A. Y. Indium arsenide: a semiconductor for high speed and electro-optical devices. Mat. Sci. Eng. B-Solid 1993, 18, 237-259.

[14] Dayeh, S. A.; Aplin, D. P. R.; Zhou, X.; Yu, P. K. L.; Yu, E. T.; Wang, D. High Electron Mobility InAs Nanowire Field-Effect Transistors. Small 2007, 3, 326-332.

[15] Björk, M. T.; Ohlsson, B. J.; Thelander, C.; Persson, A. I.; Deppert, K.; Wallenberg, L. R.; Samuelson, L. Nanowire resonant tunneling diodes. Appl. Phys. Lett. 2002, 81, 4458-4460.

[16] Miao, J.; Hu, W.; Guo, N.; Lu, Z.; Zou, X.; Liao, L.; Shi, S.; Chen, P.; Fan, Z.; Ho, J. C.; Li, T.-X.; Chen, X. S.; Lu, W. Single InAs Nanowire Room-Temperature Near-Infrared Photodetectors. ACS Nano 2014, 8, 3628-3635.

[17] Doh, Y.-J.; van Dam, J. A.; Roest, A. L.; Bakkers, E. P. A. M.; Kouwenhoven, L. P.; De Franceschi, S. Tunable Supercurrent Through Semiconductor Nanowires. Science 2005, 309, 272-275.

[18] Chun-Zhi, Z.; Hong, G.; Di, Z.; Xi-Tian, Z. Local Homoepitaxial Growth and Optical Properties of ZnO Polar Nanoleaves. Chin. Phys. Lett. 2008, 25, 302.

[19] Yang, Y.; Liao, Q.; Qi, J.; Guo, W.; Zhang, Y. Synthesis and transverse electromechanical characterization of single crystalline ZnO nanoleaves. Phys. Chem. Chem. Phys. 2010, 12, 552-555.

[20] Xu, H.; Wang, W.; Zhu, W.; Zhou, L.; Ruan, M. Hierarchical-Oriented Attachment: From

One-Dimensional Cu(OH)2 Nanowires to Two-Dimensional CuO Nanoleaves. Cryst. Growth Des. 2007, 7, 2720-2724.

[21] Xu, X.; Zhang, M.; Feng, J.; Zhang, M. Shape-controlled synthesis of single-crystalline cupric oxide by microwave heating using an ionic liquid. Mater. Lett. 2008, 62, 2787-2790.

[22] He, Y. J.; Peng, J. F.; Chu, W.; Li, Y. Z.; Tong, D. G. Black mesoporous anatase TiO2 nanoleaves: a high capacity and high rate anode for aqueous Al-ion batteries. J. Mater. Chem. A 2014, 2, 1721-1731.

[23] Samanta, P. K.; Basak, S.; Chaudhuri, P. R. Fern leaves: The secret life of zinc oxide. Mater. Today 2011, 14, 295.

[24] Martelli, F.; Piccin, M.; Bais, G.; Jabeen, F.; Ambrosini, S.; Rubini, S.; Franciosi, A. Photoluminescence of Mn-catalyzed GaAs nanowires grown by molecular beam epitaxy. Nanotechnology 2007, 18, 125603.

[25] Li, J.; Liu, J.; Wang, L.-S.; Chang, R. P. Physical and electrical properties of chemical vapor grown GaN nano/microstructures. Inorg. Chem. 2008, 47, 10325-10329.

[26] Di Giacomo, E.; Iannelli, M.; Frugis, G. TALE and Shape: How to Make a Leaf Different. Plants 2013, 2, 317-342.

[27] Nakata, M.; Okada, K. The Leaf Adaxial-Abaxial Boundary and Lamina Growth. Plants 2013, 2, 174-202.

[28] Tsiantis, M.; Langdale, J. A. The formation of leaves. Curr. Opin. Plant Biol. 1998, 1, 43-48.

[29] Hiruma, K.; Yazawa, M.; Katsuyama, T.; Ogawa, K.; Haraguchi, K.; Koguchi, M.; Kakibayashi, H. Growth and optical properties of nanometer‐scale GaAs and InAs whiskers. J. Appl. Phys. 1995, 77, 447-462.

[30] Dick, K. A.; Deppert, K.; Mårtensson, T.; Mandl, B.; Samuelson, L.; Seifert, W. Failure of the Vapor−Liquid−Solid Mechanism in Au-Assisted MOVPE Growth of InAs Nanowires. Nano Lett. 2005, 5, 761-764.

[31] Dayeh, S. A.; Yu, E. T.; Wang, D. III−V Nanowire Growth Mechanism: V/III Ratio and Temperature Effects. Nano Lett. 2007, 7, 2486-2490.

[32] Zhang, Z.; Lu, Z.-Y.; Chen, P.-P.; Lu, W.; Zou, J. Controlling the crystal phase and structural quality of epitaxial InAs nanowires by tuning V/III ratio in molecular beam epitaxy. Acta Mater. 2015, 92, 25-32.

[33] http://cimewww.epfl.ch/people/stadelmann/jemsWebSite/jems.html [Online].


9Nano Res.

[34] Xu, H.; Wang, Y.; Guo, Y.; Liao, Z.; Gao, Q.; Tan, H. H.; Jagadish, C.; Zou, J. Defect-free< 110> zinc-blende structured InAs nanowires catalyzed by palladium. Nano Lett. 2012, 12, 5744-5749.

[35] Zhang, Z.; Zheng, K.; Lu, Z.-Y.; Chen, P.-P.; Lu, W.; Zou, J. Catalyst Orientation-Induced Growth of Defect-Free Zinc-Blende Structured <001̅> InAs Nanowires. Nano Lett. 2015, 15, 876-882.

[36] Wolf, D. Atomic-level geometry of crystalline interfaces. In Materials interfaces: atomic-level structure and properties. D. Wolf; S. Yip, Eds.; Chapman & Hall; Cambridge UK, 1992; pp 1-57.

[37] Wolf, D.; Merkle, K. L. Correlation between the structure and energy of grain boundaries in metals. In Materials interfaces: atomic-level structure and properties. D. Wolf; S. Yip, Eds.; Chapman & Hall; Cambridge UK, 1992; pp 87-150.

[38] Lee, B. T.; Lee, J. Y.; Bourret, E. D. Atomic structure of twins in GaAs. Appl. Phys. Lett. 1990, 57, 346-347.

[39] Jin, L.; Wang, J.; Cao, G.; Xu, Z.; Jia, S.; Choy, W. C. H.; Leung, Y. P.; Yuk, T. I. {113} Twinned ZnSe Bicrystal Nanobelts Filled with <111> Twinnings. J. Phys. Chem. C 2008, 112, 4903-4907.

[40] Zhang, Z.; Lu, Z.; Xu, H.; Chen, P.; Lu, W.; Zou, J. Structure and quality controlled growth of InAs nanowires through catalyst engineering. Nano Research 2014, 10.1007/s12274-014-0524-x, 1-10.

[41] Wagner, R. S.; Ellis, W. C. Vapor‐Liquid‐Solid Mechanism of Single Crystal Growth. Appl. Phys. Lett. 1964, 4, 89-90.

[42] Faust Jr, J. W.; John, H. F. The growth of semiconductor crystals from solution using the twin-plane reentrant-edge mechanism. J. Phys. Chem. Solids 1964, 25, 1407-1415.

[43] Gamalski, A. D.; Voorhees, P. W.; Ducati, C.; Sharma, R.; Hofmann, S. Twin Plane Re-entrant Mechanism for Catalytic Nanowire Growth. Nano Lett. 2014, 14, 1288-1292.

[44] Brenner, S. S.; Sears, G. W. Mechanism of whisker growth — III nature of growth sites. Acta Metall. 1956, 4, 268-270.

[45] Zou, J.; Paladugu, M.; Wang, H.; Auchterlonie, G. J.; Guo, Y.-N.; Kim, Y.; Gao, Q.; Joyce, H. J.; Tan, H. H.; Jagadish, C. Growth Mechanism of Truncated Triangular III–V Nanowires. Small 2007, 3, 389-393.

[46] Paladugu, M.; Zou, J.; Guo, Y.-N.; Zhang, X.; Joyce, H. J.; Gao, Q.; Tan, H. H.; Jagadish, C.; Kim, Y. Formation of Hierarchical InAs Nanoring / GaAs Nanowire Heterostructures. Angew. Chem. Int. Ed. 2009, 48, 780-783.

[47] Paladugu, M.; Zou, J.; Guo, Y.-N.; Zhang, X.; Joyce, H. J.; Gao, Q.; Tan, H. H.; Jagadish, C.; Kim, Y. Polarity driven formation of InAs/GaAs hierarchical nanowire heterostructures. Appl. Phys. Lett. 2008, 93, 201908.

Electronic Supplementary Material


M. Teng Soo1, Kun Zheng2,3 (*), Qiang Gao4, H. Hoe Tan4, Chennupati Jagadish4 and Jin Zou1,2 (*)

1 Materials Engineering, The University of Queensland, St. Lucia, QLD 4072, Australia 2 Centre for Microscopy and Microanalysis, The University of Queensland, St. Lucia, QLD 4072, Australia 3 Australian Institute for Bioengineering and Nanotechnology, The University of Queensland, St. Lucia, QLD 4072, Australia 4 Department of Electronic Materials Engineering, Research School of Physics and Engineering, The Australian National University, Canberra, ACT 2601, Australia

Supporting information to DOI 10.1007/s12274-****-****-* (automatically inserted by the publisher)

10

Physical resemblances of a real leaf and nanoleaf

This nanoleaf resembles a real leaf in both the structure and growth mechanism. The structures of both are shown in Fig. S1. The parts of nanoleaf analogue of the real leaf are given in Table S1. In nature, leaves of vascular plant have mirror symmetry with the midrib located at the center along the axial axis and they have two flat surfaces, so-called blade. The margin is the side of the leaf. The petiole is the base of leaf, which attaches the leaf blade to the stem. Besides, simple leaves initiate as entire structures at the flanks of the shoot apical meristem [1] and result in a sharp apex. Thus, the sharp apex and petiole bound the shape of leaves. In our work, the catalyst and base that attached to the substrate are the boundary of InAs nanoleaves.

Figure S1. The structures of a real leaf (left) and InAs nanoleaf (right).

Table S1 Different parts of an InAs nanoleaf analogue of a real leaf.

True leaf InAs nanoleaf Apex Tip

Primary vein (midrib) Twin boundary Margin Symmetry sidewall

Lamina/Blade flat surface Petiole Base

References

[S1] Vlad, D.; Kierzkowski, D.; Rast, M. I.; Vuolo, F.; Dello Ioio, R.; Galinha, C.; Gan, X.; Hajheidari, M.; Hay, A.; Smith, R. S.; Huijser, P.; Bailey, C. D.; Tsiantis, M. Leaf Shape Evolution Through Duplication, Regulatory Diversification, and Loss of a Homeobox Gene. Science 2014, 343, 780-783.

Address correspondence to Jin Zou, [email protected]; Kun Zheng, [email protected]

11

Silver Nanowires with Semiconducting Ligands for Low Temperature Transparent Conductors

Brion Bob,1 Ariella Machness,1 Tze-Bin Song,1 Huanping Zhou,1 Choong-Heui Chung,2 and Yang Yang1,*

1 Department of Materials Science and Engineering and California NanoSystems Institute,

University of California Los Angeles, Los Angeles, CA 90025 (USA)

2 Department of Materials Science and Engineering, Hanbat National University, Daejeon

305-719, Korea

Abstract

Metal nanowire networks represent a promising candidate for the rapid fabrication of transparent electrodes with high transmission and low sheet resistance values at very low deposition temperatures. A commonly encountered obstacle in the formation of conductive nanowire electrodes is establishing high quality electronic contact between nanowires in order to facilitate long range current transport through the network. A new system of nanowire ligand removal and replacement with a semiconducting sol-gel tin oxide matrix has enabled the fabrication of high performance transparent electrodes at dramatically reduced temperatures with minimal need for post-deposition treatments of any kind.

Keywords: Silver Nanowires, Sol-Gel, Transparent Electrodes, Nanocomposites

12

1. Introduction. Silver nanowires (AgNWs) are long, thin, and possess conductivity values on the same order of magnitude as bulk silver

(Ag) [1]. Networks of overlapping nanowires allow light to easily pass through the many gaps and spaces between nanowires, while transporting current through the metallic conduction pathways offered by the wires themselves. The high aspect ratios achievable for solution-grown AgNWs has allowed for the fabrication of transparent conductors with very promising sheet resistance and transmission values, often approaching or even surpassing the performance of vacuum-processed materials such as indium tin oxide (ITO) [2-6].

Significant electrical resistance within the metallic nanowire network is encountered only when current is required to pass between nanowires, often forcing it to pass through layers of stabilizing ligands and insulating materials that are typically used to assist with the synthesis and suspension of the nanowires [7, 8]. The resistance introduced by the insulating junctions between nanowires can be reduced through various physical and chemical means, including burning off ligands and partially melting the wires via thermal annealing [9, 10], depositing additional materials on top of the nanowire network [11-14], applying mechanical forces to enhance network morphology [15-17], or using various other post-treatments to improve the contact between adjacent wires [18-21]. Any attempt to remove insulating materials the network must be weighed against the risk of damaging the wires or blocking transmitted light, and so many such treatments must be reined in from their full effectiveness to avoid endangering the performance of the completed electrode.

We report here a process for forming inks with dramatically enhanced electrical contact between AgNWs through the use of a semiconducting ligand system consisting of tin oxide (SnO2) nanoparticles. The polyvinylpyrrolidone (PVP) ligands introduced during AgNW synthesis in order to encourage one-dimensional growth are stripped from the wire surface using ammonium ions, and are replaced with substantially more conductive SnO2, which then fills the space between wires and enhances the contact geometry in the vicinity of wire/wire junctions. The resulting transparent electrodes are highly conductive immediately upon drying, and can be effectively processed in air at virtually any temperature below 300 °C. The capacity for producing high performance transparent electrodes at room temperature may be useful in the fabrication of devices that are damaged upon significant heating or upon the application of harsh chemical or mechanical post-treatments.

2. Results and Discussion

2.1. Ink Formulation and Characterization

Dispersed AgNWs synthesized using copper chloride seeds represent a particularly challenging material system for promoting wire/wire junction formation, and often require thermal annealing at temperatures near or above 200 °C to induce long range electrical conductivity within the deposited network [22, 23]. The difficulties that these wires present regarding junction formation is potentially due to their relatively large diameters compared to nanowires synthesized using other seeding materials, which has the capacity to enhance the thermal stability of individual wires according to the Gibbs-Thomson effect. We have chosen these wires as a demonstration of pre-deposition semiconducting ligand substitution in order to best illustrate the contrast between treated and untreated wires.

Completed nanocomposite inks are formed by mixing AgNWs with SnO2 nanoparticles in the presence of a compound capable of stripping the ligands from the AgNW surface. In this work, we have found that ammonia or ammonium salts act as effective stripping agents that are able to remove the PVP layer from the AgNW surface and allow for a new stabilizing matrix to take its place. Figure 1 shows a schematic of the process, starting from the precursors used in nanowire and nanoparticle synthesis and ending with the deposition of a completed film. The SnO2 nanoparticle solution naturally contains enough ammonium ions from its own synthesis to effectively peel the insulating ligands from the AgNWs and allow the nanoparticles to replace them as a stabilizing agent. If not enough SnO2 nanoparticles are used in the mixture, then the wires will rapidly agglomerate and settle to the bottom as large clusters. Large amounts of SnO2 in the mixture gradually begin to increase the sheet resistance of the nanowire network upon deposition, but greatly enhance the uniformity, durability, and wetting properties of the resulting films. We have found that AgNW:SnO2 weight ratios ranging between 2:1 and 1:1 produce well dispersed inks that are still highly conductive when deposited as films.

The nanowires were synthesized using a polyol method that has been adapted from the recipe described by Lee et al. [22, 23] Silver nitrate dissolved in ethylene glycol via ultrasonication was used as a precursor in the presence of copper chloride and PVP to provide seeds and produce anisotropic morphologies in the reaction products. Synthetic details can be found in the experimental section. Distinct from previous recipes, we have found that repeating the synthesis two times without cooling down the reaction mixture generally produces significantly longer nanowires than a single reaction step. The lengths of nanowires produced using this method fall over a wide range from 15 to 65 microns, with diameters between 125 and 250 nm. This range of diameters is common for wires grown using copper chloride seeds, although the double reaction produces a number of wires with roughly twice their usual diameter. The morphology of the as-deposited AgNWs as determined via SEM is shown in Figure 2(a), higher magnification images are also provided in Figures 2(c) and 2(d).

13

The SnO2 nanoparticles were synthesized using a sol-gel method typical for multivalent metal oxide gelation reactions. A large excess of deionized water was added to SnCl4·5H2O dissolved in ethylene glycol along with tetramethylammonium chloride and ammonium acetate to act as surfactants. The reaction was then allowed to progress for at least one hour at near reflux conditions, after which the resulting nanoparticle dispersion can be collected, washed, and dispersed in a polar solvent of choice. The material properties of SnO2 nanoparticles formed using a similar synthesis method have been reported previously [24], although the present recipe uses excess water to ensure that the hydrolysis reaction proceeds nearly to completion.

After mixing with SnO2 nanoparticles, films deposited from AgNW/SnO2 composite inks show a largely continuous nanoparticle layer on the substrate surface with some nanowires partially buried and some sitting more or less on top of the film. Representative scanning electron microscopy (SEM) images of nanocomposite films are shown in Figure 2(b). Regardless of their position relative to the SnO2 film, all nanowires show a distinct shell on their outer surface that gives them a soft and slightly rough appearance, as is visible in the higher magnification images shown in Figure 2(e) and 2(f). The SnO2 nanoparticles do a particularly good job coating the regions near and around junctions between wires, and frequently appear in the SEM images as bulges wrapped around the wire/wire contact points.

The precise morphology of the SnO2 shell that effectively surrounded each AgNW was analyzed in more detail using transmission electron microscopy (TEM) imaging. Figures 3(a) to 3(c) show individual nanowires in the presence of different ligand systems: as-synthesized PVP in Figure 3(a), inactive SnO2 in Figure 3(b), and SnO2 activated with trace amounts of ammonium ions in Figure 3(c). The as-synthesized nanowires show sharp edges, and few surface features. In the presence of inactive SnO2, which is formed by repeatedly washing the SnO2 nanoparticles in ethanol until all traces of ammonium ions are removed, the nanowires coexist with somewhat randomly distributed nanoparticles that deposit all over the surface of the TEM grid. When AgNWs are mixed with activated SnO2, a thick and continuous SnO2 shell is formed along the nanowire surface. In when sufficiently dilute SnO2 solutions are used to form the nanocomposite ink, nearly all of the nanoparticles are consumed during shell formation and effectively no nanoparticles are left to randomly populate the rest of the image.

As the AgNWs acquire their metal oxide coatings in solution, the properties of the mixture change dramatically. Freshly synthesized AgNWs coated with residual PVP ligands slowly settle to the bottom of their vial or flask over a time period of several hours to one day, forming a dense layer at the bottom. The AgNWs with SnO2 shells do not settle to the bottom, but remain partially suspended even after many weeks at concentrations that are dependent on the amount of SnO2 present in the solution.

A comparison of the settling behavior of various AgNW and SnO2 mixtures after 24 hours is shown in Figures 3(d) and 3(e). The ratios 8:4, 8:16, and 8:8 indicate the concentrations of AgNWs and SnO2 (in mg/mL) present in each solution. The 8:8 uncoupled solution, in which the PVP is not removed from the AgNW surface with ammonia, produces a situation in which the nanowires and nanoparticles do not interact with one another, and instead the nanowires settle as in the isolated nanowire solution while the nanoparticles remain well-dispersed as in the solution of pure SnO2. The mixtures of nanowires and nanoparticles in which trace amounts of ammonia are present do not settle to the bottom, but instead concentrate themselves until repulsion between the semiconducting SnO2 clusters is able to prevent further settling.

Our current explanation for the settling behavior of the wire/particle mixtures is that the PVP coating on the surface of the as-synthesized wires is sufficient to prevent interaction with the nanoparticle solution. The addition of ammonia into the solution quickly strips off the PVP surface coating and allowing the nanoparticles to coordinate directly with the nanowire surface. This explanation is in agreement with the effects of ammonia has on a solution of pure AgNWs, which rapidly begin to agglomerate into clusters and sink to the bottom as soon as any significant quantity of ammonia is added to the ink.

We attribute the stripping ability of ammonia in these mixtures to the strong dative interactions that

occur via the lone pair on the nitrogen atom interacting with the partially filled d-orbitals of the Ag atoms

on the nanowire surface. These interactions are evidently strong enough to displace the existing

coordination of the five-membered rings and carbonyl groups contained in the original PVP ligands and

allow the ammonia to attach directly to the nanowire surface. Since ammonia is one of the original

surfactants used to stabilize the surface of the SnO2 nanoparticles, we consider it reasonable that ammonia

coordination on the nanowire surface would provide an appropriate environment for the nanoparticles to

adhere to the AgNWs.

14

Scanning Energy Dispersive X-ray (EDX) Spectroscopy was also conducted on nanoparticle-coated AgNWs in order to image the presence of Sn and Ag in the nanowire and shell layer. The line scan results are shown in Figure 3(f), having been normalized to better compare the widths of the two signals. The visible broadening of the Sn lineshape compared to that of Ag is indicative of a Sn layer along the outside of the wire. The increasing strength of the Sn signal toward the center of the AgNW is likely due to the enhanced interaction between the TEM’s electron beam and the dense AgNW, which then improves the signal originating from the SnO2 shell as well. It is also possible that there is some intermixing between the Ag and Sn x-ray signals, but we consider this to be less likely as the distance between their characteristic peaks should be larger than the detection system’s energy resolution.

2.2. Network Deposition and Device Applications

For the deposition of transparent conducting films, a weight ratio of 2:1 of AgNWs to SnO2 nanoparticles was chosen in order to obtain a balance between the dispersibility of the nanowires, the uniformity of coated films, and the sheet resistance of the resulting conductive networks. Nanocomposite films were deposited on glass by blade coating from an ethanolic solution using a scotch tape spacer, with deposited networks then being allowed to dry naturally in air over several minutes.

The as-dried nanocomposite films are highly conductive, and require only minimal thermal treatment to dry and harden the film. Without the use of activated SnO2 ligands, deposited nanowire networks are highly insulating, and become conductive only after annealing at above 200 °C. The sheet resistance values of representative films are shown in Figure 4(a). The capability to form transparent conductive networks in a single deposition step that remain useful over a wide range of processing temperatures provides a high degree of versatility for designing thin film device fabrication procedures.

Figure 5(a) shows the sheet resistance and transmission of a number of nanocomposite films deposited from inks containing different nanowire concentrations. The deposited films show excellent conductivity at transmission values up to 85%, and then rapidly increase in sheet resistance as the network begins to reach its connectivity limit. The optimum performance of these networks at low to moderate transmission values is a consequence of the relatively large nanowire diameters, which scatter a noticeable amount of light even when the conditions required for current percolation are just barely met. Nonetheless, the sheet resistance and transmission of the completed nanocomposite networks place them within an acceptable range for applications in a variety of optoelectronic devices. Figure 5(b) shows the wavelength dependent transmission spectra of several nanowire networks, which transmit light well out into the infrared region. The presence of high transmission values out to wavelengths well above 1300 nm, where ITO or other conductive oxide layers would typically begin to show parasitic absorption, is due to the use of semiconducting SnO2 ligands, which is complimentary to the broad spectrum transmission of the silver nanowire network itself.

Avoiding the use of highly doped nanoparticles has the potential to provide optical advantages, but can create difficulties when attempting to make electrical contact to neighboring device layers. In order to investigate their functionality in thin film devices, we have incorporated AgNW/SnO2 nanocomposite films as electrodes in amorphous silicon (a-Si) solar cells. Two contact structures were used during fabrication: one with the nanocomposite film directly in contact with the p-i-n absorber structure and one with a 10 nm Al:ZnO (AZO) layer present to assist in forming Ohmic contact with the device. The I-V characteristics of the resulting devices are shown in Figure 6(a).

The thin AZO contact layers typically show sheet resistance values greater than 2.5 kΩ/⧠, and so cannot be responsible for long range lateral current transport within the electrode structure. However, their presence is clearly beneficial in improving contact between the nanocomposite electrode and the absorber material, as the SnO2 matrix material is evidently not conductive enough to form a high quality contact with the p-type side of the a-Si stack. We hope that future modifications to the AgNW/SnO2 composite, or perhaps the use of islands of high conductivity material such as a discontinuous layer of doped nanoparticles will allow for the deposition of completed electrode stacks that provide both rapid fabrication and good performance.

Figure 6(b) contains the top view image of a completed device. The enhanced viscosity of the nanowire/sol-gel composite inks allows for films to be blade coated onto substrates with a variety of surface properties without reductions in network uniformity. In contrast with traditional back electrodes deposited in vacuum environments, the nanocomposite can be blade coated into place in a single pass under atmospheric conditions and dried within moments. We anticipate that the use of sol-gel mixtures to enhance wetting and dispersibility may prove useful in the formulation of other varieties of semiconducting and metallic inks for deposition onto a variety of substrate structures.

3. Conclusions

In summary, we have successfully exchanged the insulating ligands that normally surround as-synthesized AgNWs with shells of substantially more conductive SnO2 nanoparticles. The exchange of one set of ligands for the other is mediated by

15

the presence of ammonia during the mixing process, which appears to be necessary for the effective removal of the PVP ligands that initially cover the nanowire surface. The resulting nanowire/nanoparticle mixtures allow for the deposition of nanocomposite films that require no annealing or other post-treatments to function as high quality transparent conductors with transmission and sheet resistance values of 85% and 10 Ω/⧠, respectively. Networks formed in this manner can be deposited quickly and easily in open air, and have been demonstrated as an effective n-type electrode in a-Si solar cells when a thin interfacial layer is deposited first to ensure good electronic contact with the rest of the device. The ligand management strategy described here could potentially be useful in any number of material systems that presently suffer from highly insulating materials that reside on the surface of otherwise high performance nano and microstructures.

4. Experimental Details

Tin oxide nanoparticle synthesis. Tin chloride pentahydrate was dissolved in ethylene glycol by

stirring for several hours at a concentration of 10 grams per 80 mL to serve as a stock solution. In a typical

synthesis reaction, 10 mL of the SnCl4·5H2O stock solution is added to a 100 mL flask and stirred at room

temperature. Still at room temperature, 250 mg ammonium acetate and 500 mg ammonium acetate were

added in powder form to regulate the solution pH and to serve as coordinating agents for the growing

oxide nanoparticles. 30 ml of water was then added, and the flask was heated to 90 °C for 1 to 2 hours in

an oil bath, during which the solution took on a cloudy white color. The gelled nanoparticles were then

washed twice in ethanol in order to keep trace amounts of ammonia present in the solution. Additional

washing cycles would deactivate the SnO2, and then require the addition of ammonia to coordinate with

as-synthesized AgNWs.

Silver nanowire synthesis. Copper(ii) chloride dihydrate was first dissolved in ethylene glycol at

1 mg/ml to serve as a stock solution for nanowire seed formation. 20 ml of ethylene glycol was then added

into a 100 ml flask, along with 200 µL of copper chloride solution. the mixture was then heated to 150 °C

while stirring at 325 rpm, and .35g of PVP (MW 55,000) was added. In a small separate flask, .25 grams of

silver nitrate was dissolved in 10 ml ethylene glycol by sonicating for approximately 2 minutes, similar to

the method described here.22 The silver nitrate solution was then injected into the larger flask over

approximately 15 minutes, and the reaction was allowed to progress for 2 hours. After the reaction had

reached completion, the various steps were repeated without cooling down. 200 µL of copper chloride

solution and .35g PVP were added in a similar manner to the first reaction cycle, and another .25g silver

nitrate were dissolved via ultrasonics and injected over 15 minutes. The second reaction cycle was allowed

to progress for another 2 hours, before the flask was cooled and the reaction products were collected and

washed three times in ethanol.

Nanocomposite ink formation. After the synthesis of the two types of nanostructures is complete,

16

the double washed SnO2 nanoparticles and triple-washed nanowires can be combined at a variety of weight

ratios to form the completed nanocomposite ink. The dispersibility of the mixture is improved when more

SnO2 is used, although the sheet resistance of the final networks will begin to increase if they contain

excessive SnO2. AgNW agglomeration during mixing is most easily avoided if the SnO2 and AgNW

solutions are first diluted to the range of 10 to 20 mg/ml in ethanol, with the SnO2 solution being added

first to an empty vial and the AgNW solution added afterwards. The dilute mixture was then be allowed to

settle overnight, and the excess solvent removed to concentrate the wires to a concentration that is

appropriate for blade coating.

Film and electrode deposition. The completed nanocomposite ink was deposited onto any desired

substrates using a razor blade and scotch tape spacer. The majority of the substrates used in this study were

Corning soda lime glass, but the combined inks also deposited well on silicon, SiO2, and any other

substrates tested. Electrode deposition onto a-Si substrates was accomplished by masking off the desired

cell area with tape, and then depositing over the entire region. The p-i-n a-Si stacks and 10 nm AZO

contact layers were deposited using PECVD and sputtering, respectively.

ACKNOWLEDGMENTS The authors would like to acknowledge the use of the Electron Imaging Center for Nanomachines

(EICN) located in the California NanoSystems Institute at UCLA.

REFERENCES [1] Sun, Y.; Gates, B.; Mayers, B.; Xia, Y., Crystalline silver nanowires by soft solution

processing. Nano Lett. 2002, 2, 165-168.

[2] Kim, T.; Kim, Y. W.; Lee, H. S.; Kim, H.; Yang, W. S.; Suh, K. S., Uniformly

interconnected silver-nanowire networks for transparent film heaters. Adv. Funct.

Mater. 2013, 23, 1250-1255.

[3] Hu, L.; Wu, H.; Cui, Y., Metal nanogrids, nanowires, and nanofibers for transparent

electrodes. MRS Bull. 2011, 36, 760-765.

17

[4] van de Groep, J.; Spinelli, P.; Polman, A., Transparent conducting silver nanowire

networks. Nano Lett. 2012, 12, 3138-3144.

[5] Yang, L.; Zhang, T.; Zhou, H.; Price, S. C.; Wiley, B. J.; You, W., Solution-processed

flexible polymer solar cells with silver nanowire electrodes. ACS Appl. Mater.

Interfaces 2011, 3, 4075-4084.

[6] Scardaci, V.; Coull, R.; Lyons, P. E.; Rickard, D.; Coleman, J. N., Spray deposition of

highly transparent, low-resistance networks of silver nanowires over large areas. Small

2011, 7, 2621-2628.

[7] Wiley, B.; Sun, Y.; Xia, Y., Synthesis of silver nanostructures with controlled shapes

and properties. Acc. Chem. Res. 2007, 40, 1067-1076.

[8] Korte, K. E.; Skrabalak, S. E.; Xia, Y., Rapid synthesis of silver nanowires through a

cucl- or cucl2-mediated polyol process. J. Mater. Chem. 2008, 18, 437-441.

[9] Anuj, R. M.; Akshay, K.; Chongwu, Z., Large scale, highly conductive and patterned

transparent films of silver nanowires on arbitrary substrates and their application in

touch screens. Nanotechnology 2011, 22, 245201.

[10] Lee, J.-Y.; Connor, S. T.; Cui, Y.; Peumans, P., Solution-processed metal nanowire

mesh transparent electrodes. Nano Lett. 2008, 8, 689-692.

[11] Zhu, R.; Chung, C.-H.; Cha, K. C.; Yang, W.; Zheng, Y. B.; Zhou, H.; Song, T.-B.;

Chen, C.-C.; Weiss, P. S.; Li, G.; Yang, Y., Fused silver nanowires with metal oxide

nanoparticles and organic polymers for highly transparent conductors. ACS Nano 2011,

5, 9877-9882.

[12] Chung, C.-H.; Song, T.-B.; Bob, B.; Zhu, R.; Duan, H.-S.; Yang, Y., Silver nanowire

composite window layers for fully solution-deposited thin-film photovoltaic devices.

Adv. Mater. 2012, 24, 5499-5504.

18

[13] Kim, A.; Won, Y.; Woo, K.; Kim, C.-H.; Moon, J., Highly transparent low resistance

zno/ag nanowire/zno composite electrode for thin film solar cells. ACS Nano 2013, 7,

1081-1091.

[14] Ajuria, J.; Ugarte, I.; Cambarau, W.; Etxebarria, I.; Tena-Zaera, R. n.; Pacios, R.,

Insights on the working principles of flexible and efficient ito-free organic solar cells

based on solution processed ag nanowire electrodes. Sol. Energy Mater. Sol. Cells

2012, 102, 148-152.

[15] Tokuno, T.; Nogi, M.; Karakawa, M.; Jiu, J.; Nge, T.; Aso, Y.; Suganuma, K.,

Fabrication of silver nanowire transparent electrodes at room temperature. Nano Res.

2011, 4, 1215-1222.

[16] Lim, J.-W.; Cho, D.-Y.; Jihoon, K.; Na, S.-I.; Kim, H.-K., Simple brush-painting of

flexible and transparent ag nanowire network electrodes as an alternative ito anode for

cost-efficient flexible organic solar cells. Sol. Energy Mater. Sol. Cells 2012, 107,

348-354.

[17] De, S.; Higgins, T. M.; Lyons, P. E.; Doherty, E. M.; Nirmalraj, P. N.; Blau, W. J.;

Boland, J. J.; Coleman, J. N., Silver nanowire networks as flexible, transparent,

conducting films: Extremely high dc to optical conductivity ratios. ACS Nano 2009, 3,

1767-1774.

[18] Hu, L.; Kim, H. S.; Lee, J.-Y.; Peumans, P.; Cui, Y., Scalable coating and properties of

transparent, flexible, silver nanowire electrodes. ACS Nano 2010, 4, 2955-2963.

[19] Garnett, E. C.; Cai, W.; Cha, J. J.; Mahmood, F.; Connor, S. T.; Greyson Christoforo,

M.; Cui, Y.; McGehee, M. D.; Brongersma, M. L., Self-limited plasmonic welding of

silver nanowire junctions. Nat. Mater. 2012, 11, 241-249.

19

[20] Yu, Z.; Zhang, Q.; Li, L.; Chen, Q.; Niu, X.; Liu, J.; Pei, Q., Highly flexible silver

nanowire electrodes for shape-memory polymer light-emitting diodes. Adv. Mater.

2011, 23, 664-668.

[21] Song, T.-B.; Chen, Y.; Chung, C.-H.; Yang, Y.; Bob, B.; Duan, H.-S.; Li, G.; Tu,

K.-N.; Huang, Y., Nanoscale joule heating and electromigration enhanced ripening of

silver nanowire contacts. ACS Nano 2014, 8, 2804-2811.

[22] Lee, P.; Lee, J.; Lee, H.; Yeo, J.; Hong, S.; Nam, K. H.; Lee, D.; Lee, S. S.; Ko, S. H.,

Highly stretchable and highly conductive metal electrode by very long metal nanowire

percolation network. Adv. Mater. 2012, 24, 3326-3332.

[23] Lee, J. H.; Lee, P.; Lee, D.; Lee, S. S.; Ko, S. H., Large-scale synthesis and

characterization of very long silver nanowires via successive multistep growth. Cryst.

Growth Des. 2012, 12, 5598-5605.

[24] Bob, B.; Song, T.-B.; Chen, C.-C.; Xu, Z.; Yang, Y., Nanoscale dispersions of gelled

Sno2: Material properties and device applications. Chem. Mater. 2013, 25, 4725-4730.

20

Figure 1. Process flow diagram showing the synthesis of AgNWs and SnO2 nanoparticles followed

by stirring in the presence of ammonium salts to create the final nanocomposite ink. Transparent

conducting films were produced by blade coating the completed inks onto the desired substrate.

21

Figure 2. (a,c,d) SEM images of as-synthesized AgNWs at various magnifications. (b,e,f) SEM

images of nanocomposite films, showing the tendency of the SnO2 nanoparticles to coat the entire

outer surface of the AgNWs, increasing their apparent diameter and giving them a soft appearance.

22

Figure 3. Schematic diagrams and TEM images of (a) a single untreated AgNW, (b) an AgNW in the

presence of uncoupled SnO2 (all ammonium ions removed), and (c) an AgNW with a coordinating

SnO2 shell. Scale bars in images (a), (b), and (c) are 300 nm, 400 nm, and 600 nm, respectively. (d,e)

Optical images of AgNW and SnO2 nanoparticle dispersions mixed in varying amounts (d) before and

(e) after settling for 24 hours. The numbers associated with each solution represent the AgNW:SnO2

concentrations in mg/ml. The uncoupled solution contains AgNWs and non-coordinating SnO2

nanoparticles, and shows settling behavior similar to the pure AgNW and pure SnO2 solutions. (f)

Normalized Ag and Sn EDX signal mapped across the diameter of a single nanowire, with the inset

showing the scanning path across an isolated wire.

23

Figure 4. Sheet resistance versus temperature for films deposited using (red) AgNWs that have been

washed three times in ethanol and (blue) mixtures of AgNW and SnO2 with weight ratio of 2:1. The

annealing time at each temperature value was approximately 10 minutes. The large sheet resistance

values of the bare AgNWs when annealed below 200 °C is typical for nanowires fabricated using

copper chloride seeds, which clearly illustrate the impact of SnO2 coordination at low treatment

temperatures.

24

Figure 5. (a) Sheet resistance and transmission data for samples deposited from solutions of varying

nanostructure concentration. Each of these samples were fabricated starting from the same

nanocomposite ink, which was then diluted to a range of concentrations while maintaining the same

AgNW to SnO2 weight ratio. (b) Transmission spectra of several transparent conducting networks

chosen from the plot in plot (a).

25

Figure 6. (a) I-V characteristics of devices made with AgNW/SnO2 rear electrodes with (blue) and

without (red) a 10 nm AZO contact layer. The dramatic double diode effect is likely a result of a

significant barrier to charge injection at the electrode/a-Si interface. (b) Top view SEM image of the

AgNW/SnO2 composite films on top of the textured a-Si absorber. (c) Schematic cross section of the

a-Si device architecture used in solar cell fabrication. The thickness of the thin AZO contact layer is

exaggerated for clarity.

Documents

Mirror-twin induced bicrystalline InAs nanoleaves