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Nano Res

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Enhanced thermoelectric properties of topological crystalline insulator PbSnTe nanowires grown by a vapor transport approach

Enzhi Xu1, Zhen Li1, Jaime Avilés Acosta1, Nan Li2, Brian Swartzentruber3, ShiJian Zheng2, Nikolai Sinitsyn4, H. Htoon2, Jian Wang5 and Shixiong Zhang1 (*) Nano Res., Just Accepted Manuscript • DOI: 10.1007/s12274-015-0961-1

http://www.thenanoresearch.com on December. 1, 2015

© Tsinghua University Press 2015

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Nano Research DOI 10.1007/s12274-015-0961-1

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

64 Nano Res.

Enhanced thermoelectric properties of topological

crystalline insulator PbSnTe nanowires grown by

a vapor transport approach

Enzhi Xu1, Zhen Li1, Jaime Avilés Acosta1, Nan Li2,

Brian Swartzentruber3, ShiJian Zheng2 , Nikolai

Sinitsyn4, H. Htoon2, Jian Wang5 and Shixiong

Zhang1*

1Department of Physics, Indiana University,

Bloomington, Indiana 47405, USA.

2Center for Integrated Nanotechnologies, Materials,

Physics and Applications Division, Los Alamos

National Laboratory, Los Alamos, New Mexico

87545, USA

3Center for Integrated Nanotechnologies, Sandia

National Laboratories, Albuquerque, New Mexico

87185, USA

4Theoretical Division, Los Alamos National

Laboratory, Los Alamos, New Mexico 87545, USA

5MST-8, Los Alamos National Laboratory, Los

Alamos, New Mexico 87545, USA

Single-crystalline PbSnTe nanowires were synthesized via a vapor

transport approach, and thermoelectric devices were fabricated to

measure electrical conductivity, thermopower and thermal conductivity

on the same individual nanowires. The PbSnTe nanowires show

enhanced thermoelectric figure of merit ZTs than bulk single crystals

(Orihashi et al. J. Phys. Chem. Solids 2000, 61, 919-923), owing to both

the improved thermopower and the suppressed thermal conductivity.

Enhanced thermoelectric properties of topological crystalline insulator PbSnTe nanowires grown by a vapor transport approach

Enzhi Xu1, Zhen Li1, Jaime Avilés Acosta1, Nan Li2, Brian Swartzentruber3, ShiJian Zheng2 , NikolaiSinitsyn4, H. Htoon2, Jian Wang5 and Shixiong Zhang1(*)

1 Department of Physics, Indiana University, Bloomington, Indiana 47405, USA. 2 Center for Integrated Nanotechnologies, Materials, Physics and Applications Division, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, USA 3 Center for Integrated Nanotechnologies, Sandia National Laboratories, Albuquerque, New Mexico 87185, USA 4. Theoretical Division, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, USA 5. MST-8, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, USA

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 Type keywords PbSnTe, thermoelectrics, topological crystalline insulator, nanowire

ABSTRACT Bulk lead telluride (PbTe) and its alloy compounds are well-known thermoelectric materials for electric power generation. Among all these alloys, PbSnTe has recently been demonstrated to host unique topological surface states that may have improved thermoelectric properties. Here we report on a vapor transport growth and thermoelectric study of high quality single-crystalline PbTe and PbSnTe nanowires. The nanowires were grown along <001> direction with dominant {100} facets and their chemical compositions are strongly dependent on the substrate position in the growth reactor. We measured the thermopower, electrical and thermal conductivities of the same individual nanowires to determine their thermoelectric figure of merit ZTs. In comparison with bulk samples, the PbSnTe nanowires show both improved thermopower and suppressed thermal conductivity, leading to enhanced ZTs of ~0.018 and ~0.035 at room temperature. The enhanced thermopower may arise from the unique topological surface states and the suppression of thermal conductivity is possibly due to the increased phonon-surface scattering. Compared with the PbTe nanowires, the PbSnTe has lower thermopower but significantly higher electrical conductivity. Our work highlights nanostructuring in combination with alloying as an important approach to enhancing thermoelectric properties.

1 Introduction

Converting waste heat into electric power through the Seebeck effect is one of the most

promising approaches to harvesting renewable energy [1, 2]. The energy conversion efficiency is

quantified by the so-called figure of merit ZT of the material that is used in the thermoelectric

Nano Research DOI (automatically inserted by the publisher)

Address correspondence to [email protected]

Research Article


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generator. The ZT is dependent on the material properties, i.e. thermopower S, electrical conductivity s, and thermal conductivity κ in a relation of ZT=S2 σT/κ. The small bandgap semiconductor lead telluride (PbTe) with a rock-salt crystal structure is one of the most widely studied thermoelectric materials [3]. While the ZT of pure PbTe bulk samples is less than 1, it can be dramatically increased by alloying with other compounds or by doping with impurities [4-8]. Indeed, bulk Pb1-xSnxTe alloys exhibit a maximum ZT of ~ 1 [9, 10] and the doped PbTe1-xSex alloys have a peak ZT as high as ~ 1.8 [7, 8]. The enhanced thermoelectric properties are attributed to the fact that alloying not only reduces thermal conductivity by inducing additional phonon scatterings, but also influences thermal power and electrical conductivity favorably by modifying electronic band structures [3].

Another effective approach to enhancing ZT is to fabricate nanostructured materials [11, 12]. Quantum confinement effect in nanostructures gives rise to an increased thermopower, while surface/boundary scattering of phonons suppresses thermal conductivity [13, 14]. As an example, PbTe-based nanostructures, such as quantum dot superlattices [15] and dual-phase nanocomposites [16], have shown improved ZTs than their bulk counterparts. One-dimensional nanowires (NWs) are another format of nanostructures that have enhanced thermoelectric properties as demonstrated in many material systems [17-38]. The PbTe NWs have been fabricated by a variety of synthetic approaches, including chemical vapor transport [22, 27, 34, 39], template-directed electrodeposition [40] and hydrothermal process [41, 42]. Thermoelectric studies have indeed shown enhanced thermopower and/or reduced thermal conductivity of NWs than their bulk counterparts [22, 27, 34, 38, 39, 41-46].

In this work, we combined an alloying approach with the nanostructuring strategy to

study the thermoelectric properties of PbSnTe nanowires. As a prototypical PbTe-based alloy, Pb1-xSnxTe can form a complete solid solution (0≤x≤1) and its bulk bandgap can be tuned continuously from 0.3 eV to -0.18 eV by changing x from 0 to 1. The Pb1-xSnxTe alloy with a Sn concentration of x>0.37 has a negative bulk bandgap and has recently been demonstrated to possess unique topological surface states that are protected by crystalline mirror symmetry [47-49]. Recent theoretical studies have suggested that topological surface states may have enhanced thermoelectric properties than bulk states owing to their unique electronic band structures [50, 51]. Nanowires are anticipated to have more advantages than bulk samples in the sense that they have significantly higher surface-area-to-volume ratio which could magnify contributions from surface states.

Single-crystalline PbSnTe nanowires were grown by a vapor transport approach with Au nanoparticles as catalysts. The nanowires were grown along <001> direction of the rock-salt crystal structure and their chemical compositions are strongly dependent on the substrate position in the growth reactor. The thermopower, electrical and thermal conductivities were systematically measured on the same single-nanowire devices to accurately determine their figure of merit ZTs. In comparison with bulk samples of nearly the same chemical composition, PbSnTe nanowires show both enhanced thermopower and reduced thermal conductivity, leading to a significantly improved ZT. The same growth and thermoelectric studies were carried out on the undoped PbTe nanowires for comparison.

2 Experimental details

The PbTe and PbSnTe NWs were grown by a vapor transport approach that was widely employed for the controlled synthesis of single-crystalline nanostructures of SnTe [52-57] and its related compounds [58]. Figure 1 (a) illustrates a schematic


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of the quartz tube reactor where the growth took place. In contrast to the earlier chemical vapor transport growth of PbTe NWs in which the precursors are lead chloride and tellurium [22, 27, 34, 39], we used PbTe (Alfa Aesar, 99.999%) and SnTe (Alfa Aesar, 99.999%) powders as precursors. Argon gas with a flow rate of 15 sccm was introduced and the pressure inside the tube furnace was kept at ~ 33 Torr. Temperatures in the three zones were maintained at 820/750/600 ˚C for 30 minutes during the growth of PbTe NWs, and at 760/710/575 ˚C for 30 minutes during the growth of PbSnTe NWs. For the PbSnTe growth, the PbTe and SnTe precursors with an atomic ratio of ~1:1 were placed in the same boat and hence were at nearly the same temperature. Au nanoparticles of 20 nm diameter (Ted Pella) were deposited on the substrates and acted as metal catalysts for the vapor-liquid-solid growth. To avoid a possible coating of substrate material on nanowire surfaces as observed in previous chemical vapor deposition of PbTe [22], we used the relatively more stable tungsten foils as substrates. The nanowire morphology and chemical composition were studied by scanning electron microscope (SEM, Quanta FEI) and energy-dispersive X-ray spectrometer (Oxford INCA EDX and AMETEK EDAX), respectively. The nanowire crystalline structures were characterized using transmission electron microscope (FEI Tecnai F30).

Thermoelectric platforms that consist of a heater and four electrodes were fabricated on SiNx/SiO2/Si substrates via e-beam lithography and e-beam evaporation of 10 nm Ti & ~ 90 nm Au. Individual PbTe and PbSnTe nanowires were picked up from the growth substrates using a nanomanipulator and were then transferred onto the thermoelectric platforms. For the PbTe devices, a layer of ~120 nm Ni & ~200 nm Au was deposited on top of the four electrodes after e-beam lithography to wrap around the nanowires (nanowires were rinsed in buffered oxide etch solution for 20 s before metal deposition). For the PbSnTe devices, a layer of ~10 nm Ti & ~300 nm Au was deposited on top of the four electrodes after the nanowires were rinsed in 1% HCl for 20 s. The devices were then annealed at 300 °C for 20 s to promote good contacts. 3 Resutls and discussion

3.1 Nanowire growth and characterizations As shown in Figure 1 (b), a single vapor transport growth gives rise to a mixture of nanowires and micro-crystals. The dominant facets of both PbTe and PbSnTe are {100} planes of the cubic structure, as evidenced by the rectangular shape of each surface [Fig. 1 (b) for PbSnTe and Fig. S1 for PbTe]. This is consistent with recent density functional theory calculations which suggest that the {100} surfaces of PbTe have the lowest surface energies among all low-index planes [59]. While Au alloy particles are observed at the tip of some nanowires as shown in the lower left panels of Fig. 1 (b) and Fig. S1, some other NWs are free of particles [lower right panel of Fig. 1 (b) and Fig. S1]. We note that catalyst is not always required in the vapor phase growth of lead chalcogenide nanowires even though they have isotropic, cubic crystal structures (e.g. in the dislocation-driven growth process) [60, 61]. To test if our nanowire growth is indeed assisted by Au, we further performed a growth at the same condition but without Au nanoparticles. As shown in Fig. S2, the catalyst-free growth only leads to high density of micro-crystals, which suggests that the nanowires were indeed grown through an Au-catalyzed vapor-liquid-solid process. The absence of alloy particle on a nanowire may arise from the diffusion of Au atoms into the nanowire during the growth process [62].

The chemical composition of PbSnTe micro-crystals and nanowires has a strong dependence on the substrate position in the growth reactor. As shown in Figure 1 (a), three ~ 3-inch long substrates were placed in three unsealed quartz tubes in the growth reactor side by side from left to right. The substrates 1, 2, and 3 cover a range of x: 0 ~ 3 inches, 3.1 ~ 6.1 inches and 6.2 ~ 9.2 inches, respectively (the gap between two substrates was about 0.1 inch). As seen in Figure 1 (c), the Pb: (Pb+Sn) ratio increases from the left side of each substrate to the right side. However, it decreases dramatically

Figure 1 (a) a schematic of the vapor transport reactor for the growth of PbTe and PbSnTe nanowires. The substrates were placed

in unsealed small quartz tubes. The position and the size of the substrates are not drawn to scale. (b) SEM images of PbSnTe

nanowires and nano/micro-crystals. The lower left panel shows a typical nanowire with an Au alloy particle at its tip, while the

lower right panel shows a nanowire free of particle. The scale bars are: 100 nm (left) and 200 nm (right). The chemical

composition ratio (c) Pb: (Pb+Sn) and (d) Te: (Pb+Sn) as a function of position. The error bar of the position x on each substrate is

~0.2 inches.

from one substrate to the next, i.e. from x ~ 3 inches to ~ 3.1 inches and from x ~ 6.1 inches to ~ 6.2 inches. The composition ratio covers a broad range of ~ 0.15 to ~ 0.55 although the ratio of PbTe: (PbTe+SnTe) precursors is ~ 0.5. In contrast, the Te composition does not show a clear position/substrate dependence [Te: (Pb+Sn) ~ 1 in Figure 1 (d)]. The position dependence of the Pb: (Pb+Sn) ratio can be qualitatively understood by comparing the melting points of the two precursors: PbTe has a slightly higher melting point than SnTe (924 oC versus 790 oC), so SnTe sublimates more rapidly than PbTe at the same temperature (~710 oC) and hence the vapor concentration of Pb is less than Sn near the precursors. As a

result, the microcrystals/nanowires deposited on the left side of the first substrate (i.e. near the precursors) have a Pb: (Pb+Sn) ratio < 0.5. Since the vapors were carried from left to right by the inert Ar gas, after less Pb (or more Sn) vapor was consumed at the left side, the vapor concentration of Pb: (Pb+Sn) increases towards the right side of the substrate, leading to an increased ratio in the deposited microcrystals/nanowires. On the left side of the second substrate, the vapor of high Pb concentration from the first quartz tube [pink arrow in Figure 1 (a)] mixes with the vapor of low Pb concentration from the outside of the quartz tube (purple arrow). As a result, the Pb: (Pb+Sn) ratio of the microcrystals/nanowires


5Nano Res.

deposited on the left side of the second substrate (x ~ 3.1 inches) is lower than that on the right side of the first substrate (x~3 inches). The same scenario is applicable to substrates 2 and 3.

Microstructural characterizations using transmission electron microscopy (TEM) suggest that both the PbTe and PbSnTe nanowires are single crystalline. Figure 2 (a) and (d) are conventional bright field strain contrast TEM images, in which no grain

boundary was observed. The selected area diffraction patterns in Fig. 2 (b) and (e) clearly show that the NWs are single crystalline and their growth direction is along <001>. The high crystalline quality is further confirmed by the high-resolution TEM images shown in Figure 2 (c) and (f). The interplanar spacing between the {200} planes were derived from the fast Fourier transform (FFT). It is ~ 6.4 Å for PbTe and ~ 6.3 Å for PbSnTe, in good agreement with bulk samples [63].

Figure 2 (a) A low magnification TEM image and (b) a selected area diffraction pattern of a typical PbTe nanowire; (c) a high

resolution TEM image of another typical PbTe nanowire; (d) a low magnification TEM image and (e) a selected area diffraction

pattern of a typical PbSnTe nanowire; (f) a high resolution TEM image of another PbSnTe nanowire. Fast Fourier Transform

filtering was applied to obtain the lattice-resolved TEM images.

3.2 Thermoelectric measurement of individual nanowires Figure 3 (a) shows a typical thermoelectric device used for a correlated measurement of S, σ and κ on the same nanowire. We measured three PbSnTe NWs (denoted as PbSnTe NW-1, -2 and -3) along with three undoped PbTe NWs (denoted as PbTe NW-1, -2, and -3) for comparison. The three PbSnTe NWs were taken from the right side of the third

substrate in the same growth so that their chemical composition ratios of Pb: (Pb+Sn) are close to ~ 0.5. The nanowire dimensions were summarized in table 1. Two-probe current-voltage measurements indicate good Ohmic contacts of the devices (Fig. S3 in the Supplementary Information). In the thermopower measurement, we applied an electric current through the resistive heater to create Joule heating, and then measure the thermal voltage ∆V


6 Nano Res.

and the temperature difference ∆T between the two electrodes B and C. Fig. S4 in the supplemental information shows the ∆V - ∆T curves taken on a typical nanowire device at different measurement temperatures from 25 to 300 K. The thermopower of a device is calculated as S=∆V/∆T, and the thermopower of the nanowire was determined by subtracting the contributions from the metal electrodes. Table 1 The dimensions of PbTe and PbSnTe nanowires

measured in this work.

Nanowire Width (nm) Thickness (nm)

PbTe NW-1 217 281

PbTe NW-2 222 124

PbTe NW-3 232 80

PbSnTe NW-1 459 170

PbSnTe NW-2 207 252

PbSnTe NW-3 156 97

The thermopowers of both PbTe and PbSnTe

nanowires are positive [Figure 3 (b)], indicating that the majority charge carriers are holes for all nanowires. We note that both positive and negative thermopowers were reported in undoped PbTe nanowires, depending on the growth methods [34, 38, 39, 41-43, 45, 46]. The p-type conductivity in our nanowires suggest the existence of native defects such as Pb/Sn vacancies that act as acceptors. The thermopowers of the three PbTe nanowires at 300 K are: 372 µV/K (PbTe NW-1), 354 µV/K (PbTe NW-2) and 399 µV/K (PbTe NW-3), which are about 40% higher than the bulk value (~265 µV/K) [64]. An enhancement of thermopower was also observed in the PbTe nanowires grown by hydrothermal process [41, 42, 46], stress-induced method [45] and lithographically patterned nanowire electrodeposition [43]. The room temperature thermopowers of PbSnTe NWs are: ~ 33 µV/K (PbSnTe NW-1), 26 µV/K (PbSnTe NW-2) and 24 µV/K (PbSnTe NW-3), about 2~3 times higher than that of the bulk Pb0.5Sn0.5Te single crystals (~10.4 µV/K) [9]. In contrast to PbTe NWs, the thermopower of PbSnTe shows a strong temperature dependence, i.e. it decreases

dramatically as the temperature decreases [Figure 3 (b)].

The electrical conductivities of PbTe NW-1 and PbSnTe NW-1 and -2 were determined by a standard four-probe resistance measurement (four-probe measurement was not performed on PbTe NW-2, -3 and PbSnTe NW-3 due to device failure). As shown in Figure 3 (c), the electrical conductivity of PbTe NW-1 at 300 K is ~1290 S/m, an order of magnitude lower than the bulk value (~25641 S/m) [64]. This difference suggests that the cation vacancy density is much lower in our nanowires than in the bulk samples. Indeed, the electrical conductivity decreases during cooling, indicating a lightly doped semiconducting behavior. In contrast, the PbSnTe NWs show a metallic behavior [Figure 3 (c)], manifesting a degenerately doped nature. The electrical conductivities at 300 K are on the same order of magnitude as the bulk Pb0.5Sn0.5Te single crystals [9], confirming their comparable charge carrier densities and chemical compositions. As a result, the much higher thermopower of PbSnTe NWs in comparison with bulk crystals should be attributed to other factors.

We note that undoped SnTe NWs also show enhanced thermopower than bulk samples [65]. Although the precise nature is not clear, the enhancement in NW samples may be related to the topological surface states possessed by both SnTe and PbSnTe. The thermopower of a material depends on its electronic band structure and can be written as S=-[∫v2τ(E-Ef)∂f0/∂ED(E)dE]/[eT∫v2τ∂f0/∂ED(E)dE], where v is the electron/hole velocity, τ is the relaxation time, Ef is the Fermi energy, f0 is the Fermi-Dirac distribution, and D(E) is the density of states (DOS). A large thermopower can be realized by creating a delta function in the DOS. The {100} surfaces of the SnTe-based topological crystalline insulators (TCIs) indeed have Van-Hove singularities in their DOS [47, 48]. In other words, the topological surface may have a significantly higher thermopower than the interior of the TCI. Since nanowires have a large surface-area-to-volume ratio than bulk samples, they could have enhanced thermopower. Future theoretical study of the thermopower of topological surfaces would be of great importance to understand the observed experiment results.

Figure 3 (a) An SEM image of a typical thermoelectric device. Temperature dependent (b) thermopower S, (c) electrical

conductivity σ, and (d) power factor S2σ of PbTe and PbSnTe NWs. The corresponding data of bulk samples reported in reference

[64] (Heremans et al.) and reference [9] (Orihashi et al.) are shown for comparison.

The power factor of PbTe NW is limited by its

low electrical conductivity, while that of PbSnTe NWs is magnified by its enhanced thermopower. As shown in Figure 3 (d), the room temperature power factor of PbTe NW is only ~ 179 µWm-1K-2, an order of magnitude lower than the bulk value (~1800 µWm-1K-2) [64]. This result agrees qualitatively with the previous studies on lead chalcogenide (e.g. PbTe and PbSe) nanowires [23, 41, 43, 45]. The lower electrical conductivity in nanowires than in bulk samples may arise from a reduced charge carrier density and/or a suppressed carrier mobility. As discussed earlier, the lower charge carrier density may be due to a smaller concentration of Pb vacancies (p-type defects) in our high quality single crystalline nanowires. The

suppressed carrier mobility is likely due to the enhanced scattering of charge carriers by surface trap states in nanowires [23]. In contrast, the power factors of the two PbSnTe NWs at room temperature are 154 µWm-1K-2 (NW-1) and 104 µW m-1K-2 (NW-2), both of which are much higher than that of the bulk single crystal (~ 41 µW m-1K-2) [9]. The difference in power factors between the two PbSnTe NWs highlights the importance of measuring thermopower, electrical and thermal conductivities on the same nanowire in order to accurately determine the figure of merit ZT.

We further measured the thermal conductivity using a self-heating method [66], in which the nanowire was heated up by passing a high electric


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current I. Thermal conductivity was extracted by measuring the mean temperature increase ∆TM of the nanowire as a function of the Joule heating power I2R based on the relation: κ=(I2Rl)/(12∆TMA) in the limit of ∆R<<R, where A, l, R and ∆R are the cross sectional area of the nanowire, the channel length, the electrical resistance of nanowire, and the change in resistance due to Joule heating, respectively [37, 65, 66]. We note that a high bias voltage for self-heating could generate non-equilibrium between the charge carriers and the phonons, which will complicate the thermal conductivity measurement. To ensure that the non-equilibrium effect is not significant in our case, we roughly estimated the temperature difference between the charge carriers and the lattice. With a maximum bias voltage of only 0.01~0.03 V and a channel length of roughly 3 µm, the estimated temperature difference between the charge carriers and lattice is only a few percent of the lattice temperature [67, 68], indicating that the charge carriers and the phonons are in nearly equilibrium. As shown in Figure 4 (a), the I2R versus ∆TM of a typical PbSnTe device is linear, further justifying the validity of the measurement. Clear non-linearity was observed in the measurement of the PbTe device, which may be due to the highly resistive nature of the device. As a result, here we only extracted the total thermal conductivity κtot of PbSnTe NWs based on the slope of the I2R - ∆TM curve. As shown in Figure 4 (b), κtot of PbSnTe NW-1 and NW-2 at 300 K are ~1.3 W m-1 K-1 and ~1.7 W m-1 K-1, respectively. These numbers are only about half of the value of bulk single crystal sample [9], suggesting enhanced surface boundary scattering of phonons. Upon cooling, κtot increases slightly and reaches its maximum value at ~ 150 K, after which it decreases rapidly. A similar temperature dependence was observed in single-crystalline PbTe NWs, in which thermal conductivity was suppressed by surface scattering of phonons [27].

Figure 4 (a) A linear fitting to a typical I2R versus ∆TM curve for thermal conductivity measurement. (b) Thermal conductivity κtot and (c) lattice thermal conductivity κlatt of PbSnTe NWs as a function of temperature. The corresponding data of bulk single crystal samples reported in reference [9] (Orihashi et al.) are shown for comparison.

We decomposed the total thermal conductivity into the electronic part κe and the lattice part κlatt, i.e. κtot=κe+κlatt. According to the Wiedemann-Franz law,


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the electronic component can be written as a function of electrical conductivity: κe=LσT, where L is the Lorentz number. Given the degenerately doped nature of the PbSnTe NWs, we used L=2.44×10-8 V2K-2 [9] to calculate κe and then determined the lattice thermal conductivity κlatt=κtot-κe. As shown in Figure 4 (c), κlatt at 300 K are ~ 0.28 W m-1 K-1 for PbSnTe NW-1 and ~ 0.59 W m-1

K-1 for NW-2, lower than the bulk value of ~ 1.1 W m-1 K-1 [9]. This further suggests that the suppression of thermal conductivity is primarily due to the enhanced surface scattering of phonons in nanowire samples.

We finally calculated the figure of merit ZT of the PbSnTe nanowires. The ZT increases as a function of temperature and it reaches ~ 0.035 for NW-1 and 0.018 for NW-2 at 300 K [Figure 5]. The temperature dependence of ZT is mainly attributed to the thermopower which increases dramatically with temperature [Figure 3 (b)]. The maximum ZT of PbSnTe NWs at 300 K is nearly an order of magnitude higher than the bulk value [9], owing to both the enhanced thermopower and the reduced thermal conductivity. We also note that the n-type PbTe NWs grown by a similar vapor transport approach (using different precursors) have a ZT of ~0.0054 at 300 K [34]. Therefore, the ZTs of the PbSnTe NWs are several times higher than that of the PbTe NWs [34]. This enhancement is attributed to a significantly higher electrical conductivity in the PbSnTe NWs.

4 Conclusions

In summary, we performed a vapor transport growth and thermoelectric studies of high quality single-crystalline PbTe and PbSnTe nanowires. The NWs were grown along the <001> direction with dominant {100} facets. The chemical composition ratio of Pb: (Pb+Sn) in the PbSnTe NWs has a strong dependence on the substrate position in the growth reactor. Correlated measurements of S, σ, and κ were performed on the same single-nanowire devices to accurately determine the figure of merit ZT.

Figure 5 Temperature dependent figure of merit ZT of PbSnTe NWs. The ZT of bulk single crystal sample calculated based on S, σ, κ at ~300 K in reference [9] (Orihashi et al.) is shown for comparison. Comparing with the bulk single crystal samples of nearly the same chemical composition, PbSnTe nanowires show both enhanced thermopower and reduced thermal conductivity, leading to a maximum improvement of ZT by a factor of nearly ~10. We suggest that the suppression of thermal conductivity may be due to the increased phonon-surface scattering, while the enhanced thermopower could be related to the unique topological surface states. In comparison with the PbTe nanowires grown using the same method, the PbSnTe has a lower thermopower but a significantly higher electrical conductivity. The ZTs of the PbSnTe NWs are also several times higher than that of the n-type PbTe NWs reported in the literature [34]. Our work highlights nanostructuring in combination with alloying as an important approach to enhancing thermoelectric properties. We note that the PbSnTe bulk crystals show high ZTs at high temperatures [9], so future measurements on nanowires at those temperatures will be important to test if the thermoelectric enhancement observed here persists well above room temperature.

Acknowledgements

We thank Dr. Julio Martinez, John Nogan, Anthony R. James, Douglas V. Pete, Denise B. Webb and


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Renjie Chen for experimental assistances. S.X.Z., H.H., and N.S. acknowledge support from the Laboratory Directed Research & Development program at Los Alamos National Laboratory. This work was performed, in part, at the Center for Integrated Nanotechnologies, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science by Los Alamos National Laboratory (Contract DE-AC52-06NA25396) and Sandia National Laboratories (Contract DE-AC04-94AL85000). We also thank the Indiana University Nanoscale Characterization Facility for access to the instrumentation. Electronic Supplementary Material: Supplementary material (SEM images of PbTe nanowires & microcrystals and PbSnTe microcrystals; two probe I-V characteristics of PbTe and PbSnTe nanowires; and detailed information about thermopower measurement) is available in the online version of this article at http://dx.doi.org/10.1007/s12274-***-****-* References [1] Bell, L. E. Cooling, heating, generating power, and

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14

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

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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).

16

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.

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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

18

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,

19

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.

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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.

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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.

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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.

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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.

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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).

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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.

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