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Article Cite This: J. Phys. Chem. C XXXX, XXX, XXX−XXX

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SnO2/PbOx (x = 1, 2) Core−Shell Nanowires and Their Growth on C‑Fiber Networks for Energy Storage Matthew Zervos,*,† Andreas Othonos,‡ Eugenia Tanasa,̆ ∥ Eugeniu Vasile,∥ and Epameinondas Leontidis§

J. Phys. Chem. C Downloaded from pubs.acs.org by KAOHSIUNG MEDICAL UNIV on 11/01/18. For personal use only.

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Nanostructured Materials and Devices Laboratory, School of Engineering, ‡Laboratory of Ultrafast Science, Department of Physics and §Laboratory of Physical Chemistry of Colloids and Interfaces, Department of Chemistry, University of Cyprus, P. O. Box 20537, Nicosia 1678, Cyprus ∥ Department of Science and Engineering of Oxides Materials and Nanomaterials, Politehnica University of Bucharest, 313 Splaiul Independentei, Bucharest 060042, Romania S Supporting Information *

ABSTRACT: SnO2 nanowires were grown on Si, fused SiO2, and C fibers by the vapor−liquid−solid mechanism at 800 °C and 10−1 mbar, and SnO2/PbO core−shell nanowires were obtained by the deposition of 50 nm Pb over the SnO2 nanowires followed by annealing between 100 and 200 °C. The SnO2/ PbO nanowires have diameters of 100−300 nm and lengths up to 100 μm and consist mainly of tetragonal rutile SnO2 and PbO. Higher temperatures between 300 and 500 °C resulted in the formation of Pb2O3 and Pb3O4 with monoclinic and orthorhombic crystal structures, but the SnO2/PbO and SnO2/Pb2O3 nanowires had low conductivities of 10−1 S/cm. In contrast, highly conductive SnO2/PbO2 nanowires were obtained by electrodeposition of PbO2 in 0.3 M HNO3 and 1 M Pb(NO3)2 (aq). PbO2 forms a straddlingtype heterojunction with SnO2, and the one-dimensional (1D) electron gas distribution is confined in the PbO2 shell for sufficiently thick shells, as shown by the self-consistent solution of the Poisson−Schrödinger equations in the effective mass approximation. The SnO2/PbO2 nanowires exhibit an open-circuit potential of 1.8 V versus C-fiber networks in 5 M H2SO4 (aq) and show symmetric cyclic voltammetry curves, suggesting a suppression of the redox reactions related to SnO2 and a high specific capacity of 206 mAh/g. We discuss the potential of both SnO2 and SnO2/PbO2 nanowires on C fibers for the attainment of even higher specific capacity in a Li-ion battery.

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INTRODUCTION Metal−oxide−semiconductor nanowires (NWs), such as ZnO,1 SnO2,22 In2O3,3 and Ga2O3,4 have been employed in the past for the realization of energy storage devices. These metal oxides are n-type semiconductors with large energy band gaps between 3 and 4 eV; thus, they have relatively low carrier densities on the order of ≈1016 cm−3 as we have shown previously for SnO2 NWs grown by the vapor−liquid−solid (VLS) mechanism at 800 °C and 10−1 mbar.5 Consequently, they require doping to improve their conductivity and the performance of devices. This has been achieved via the incorporation of Sb6 and Mo7 in SnO2, although recently, Ma et al.8 showed theoretically that substantial band gap tuning and a strain-controlled semiconductor-to-semimetal transition are possible through incorporation of Pb in SnO2. In contrast to the above, PbO2 is an n-type semiconductor with a small indirect band gap of ∼0.5 eV and metallic conductivity. The origin of the high conductivity and carrier density in PbO2 is attributed to oxygen vacancies, which form donor states in the conduction band (CB), whereas the dipole-forbidden gap combined with the large carrier density and Moss−Burstein shift results in large optical band gaps.9 PbO2 NWs have been © XXXX American Chemical Society

obtained by electrodeposition (ELD) on bulk PbO2 and used for energy storage and the realization of a lead−acid battery (LAB), as shown by Moncada et al.10,11 Interestingly, the βform of PbO2 has a tetragonal rutile crystal structure with a = 0.491 nm and c = 0.3385 nm, which are close to the lattice constants of tetragonal rutile SnO2, i.e., a = 0.4737 nm and c = 0.318 nm; thus, they differ only by ≈3%, which in turn means that high-crystalline-quality PbO2 can be grown on SnO2. In fact, it has been shown theoretically by Butler et al.12 that the use of an ultrathin epitaxial layer of PbO2 on the surface of SnO2 can tune its work function and achieve energy levels commensurate with those of important technological materials. However, in the past, PbO2 has been deposited on SnO2 mainly for improving the performance of LABs, where the SnO2 acts as an interlayer between PbO2 and Ti.13,14 In addition, a PbO2/SnO2 nanoparticle composite has been prepared by ELD on Ti and used as a supercapacitor by Dan et al.,15 whereas PbO2 and SnO2 nanoparticles have also been Received: August 3, 2018 Revised: September 28, 2018 Published: October 22, 2018 A

DOI: 10.1021/acs.jpcc.8b07526 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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The Journal of Physical Chemistry C combined for energy conversion in a dye-sensitized solar cell.16 On the other hand, SnO2/PbO core−shell NWs have been investigated by Hyun et al.17 as a sensor; however, to the best of our knowledge, no study has investigated the properties of SnO2/PbO2 and SnO2/PbO core−shell NWs for energy storage. Consequently, we have grown SnO2 NWs on Si, fused SiO2 (f-SiO2), and also on C-fiber networks (CFNs) by the VLS mechanism, after which we obtained SnO2/PbO and SnO2/ Pb2O3 core−shell NWs by the deposition of 50 nm Pb over the SnO2 NWs and annealing between 100 and 500 °C. In contrast, highly conductive PbO2 was deposited by ELD on the SnO2 NWs. We describe in detail their structural, electrical, and optical properties and their potential for the realization of high-specific-capacity batteries.

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transmission spectrum using a PerkinElmer UV−vis spectrophotometer. In addition to the above, we have also grown SnO2 NWs on a 15 mm × 30 mm CFN using Au as a catalyst and the same growth conditions. Then, SnO2/PbO2 core−shell NWs were obtained by ELD of PbO2 on the SnO2 NWs, which was carried out at 25 °C in a solution of 0.3 M HNO3 and 1 M Pb(NO3)2 using a Pt counter electrode (CE) and applying a potential difference of ≈3 V for 1000 s with respect to a Ag/ AgCl reference electrode (RE). Finally, the open-circuit potential (OCP) and cyclic voltammograms (CV) of the SnO2/PbO2 core−shell NWs versus CFN in 5 M H2SO4 (aq) were measured using an ELD cell with a Digy-Ivy potentiostat, whereas the charge− discharge (CD) curves were measured using a Keithley 2635A and 2182A controlled by LabView.

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METHODS

RESULTS AND DISCUSSION Growth and Properties of SnO2 NWs on Si, f-SiO2, and C Fibers. The reaction of Sn with O2 at 800 °C and 10−1 mbar results into a high yield and uniform growth of SnO2 NWs on 10 mm × 10 mm Si and f-SiO2. The SnO2 NWs have diameters of ≈100 nm and lengths up to ≈100 μm and are grown by the VLS mechanism, whereby Sn enters the Au particles and forms liquid Au:Sn particles because the Au:Sn eutectic is liquid at 280 °C, even with a 20% Sn content. Upon saturation, solid SnO2 forms beneath the liquid Au:Sn particles, leading to one-dimensional growth.18 A schematic representation of the VLS growth mechanism is shown in Figure 1, and a typical SEM image of SnO2 NWs on Si is

Initially, SnO2 NWs were grown on Si and f-SiO2 via the VLS mechanism using a 1″ hot wall, low-pressure chemical vapor deposition (LPCVD) reactor capable of reaching 1100 °C, which was fed via a microflow leak valve positioned on the upstream side just after the gas manifold consisting of four mass flow controllers. A chemically resistant, rotary pump capable of reaching 10−4 mbar was connected downstream. For the growth of the SnO2 NWs, 100 mg of Sn (2−14 mesh, 99.9%; Sigma-Aldrich) was weighed with an accuracy of ±1 mg. Square samples ≈10 mm × 10 mm of Si(001) and f-SiO2 were cleaned sequentially in trichloroethylene, methanol, acetone, and isopropanol; rinsed with de-ionized water; and dried with nitrogen. Then, 1 nm of Au was deposited over the Si and f-SiO2. The elemental Sn and the Si and f-SiO2 substrates were loaded in a quartz boat that was positioned inside the 1″ LPCVD reactor, which was pumped down to 10−4 mbar and then purged with 500 sccm of Ar for 10 min at 10−1 mbar. Afterwards, the temperature was ramped up to 800 °C at 30 °C/min using the same flow of Ar. Upon reaching 800 °C, a flow of 10 sccm O2 was added to the flow of Ar to grow the SnO2 NWs for 30 min, followed by cooling without O2. Care was taken to maintain a clean high-temperature zone for the growth of the SnO2 NWs. Subsequently, 50 nm Pb was deposited over the SnO2 NWs, which were annealed between 100 and 500 °C for 60 min using a ramp rate of 10 °C/min. For comparison, 50 nm Pb was deposited directly on 10 mm × 10 mm f-SiO2 and annealed along with the Pb that was deposited on the SnO2 NWs. The morphology and composition of the resultant SnO2, SnO2/PbO, and SnO2/Pb2O3 core−shell NWs were determined by scanning electron microscopy (SEM) and energydispersive X-ray analysis (EDX), whereas their crystal structure was determined by X-ray diffraction (XRD). High-resolution transmission electron microscopy (HRTEM) was carried out using a Tecnai F30 G2 S-Twin operated at 300 kV equipped with EDX. The electrical properties and conductivity of the SnO2, SnO2/PbO, and SnO2/Pb2O3 core−shell NWs obtained on f-SiO2 were measured by TCS, as described in detail elsewhere.5 The optical properties of the SnO2/PbO and SnO2/Pb2O3 core−shell NWs on Si were investigated by photoluminescence (PL) at room temperature using an excitation wavelength of λ = 260 nm. In addition, the optical band gaps of the PbO and Pb2O3 layers obtained on f-SiO2 were determined by measuring the steady-state absorption

Figure 1. (a) Schematic diagram of the VLS growth mechanism and (b) deposition of Pb over the SnO2 NWs. Formation of (c) SnO2/ PbO core−shell NWs and (d) SnO2/Pb2O3 core−shell NWs, by annealing between 100 and 500 °C.

shown in Figure 2a. In general, the thickness of the Au deposited on Si and f-SiO2 governs the diameter of the Au particles and the SnO2 NWs, which has to be as small as possible to obtain a maximum surface area for the realization of high-capacity energy storage devices. In addition to Au, we find that the deposition of 1 nm Ni on Si or f-SiO2 also leads to the growth of SnO2 NWs with diameters and lengths similar to those shown in Figure 2a. The VLS growth mechanism permits growth of SnO2 NWs on selected locations on Si and f-SiO2 as we always observed Au:Sn particles on their ends and that no growth took place on plain Si. However, we ought to point out that SnO2 NWs have been obtained by Luo et al.19 on Si B


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Figure 2. SEM images of SnO2 NWs on (a) Si, (b) CFNs, and (c) a strand of C fiber also showing a bare section with no SnO2 NWs. The top, end, and side views of the CFN without any SnO2 NWs are shown in (d), (e), and (f), respectively.

without a catalyst via the formation of Sn droplets and a selfcatalyzed VLS mechanism using SnO2 as the source of Sn and an Ar:H2 carrier gas but at a higher temperature of 1050 °C.19 The SnO2 NWs on Si and f-SiO2 have a tetragonal rutile crystal structure, and in the past, we have shown that they exhibit broad PL emission with a maximum at ≈550 nm or 2.5 eV; this is smaller than the energy band gap of SnO2, i.e., 3.7 eV, due to radiative recombination via surface states residing energetically in the energy band gap that are related to oxygen vacancies. More importantly, we have shown that SnO2 NWs grown under these conditions have a carrier density on the order of 1016 cm−3 and mobility of 70 cm2/(V s) as determined from THz conductivity spectroscopy.5 Hence, doping is required to increase their conductivity20 or alternatively one may deposit highly conductive oxides, such as PbO2, on SnO2 NWs. However, before elaborating further, it is important to point out that we have also grown SnO2 NWs on a CFN (Supporting Information, Section S1) because Si and f-SiO2 are certainly not suitable for making LABs. An SEM image of the SnO2 NWs on a 10 mm × 10 mm CFN is shown in Figure 2b. It is evident that the SnO2 NWs grow in a radial fashion, as shown by the SEM image of SnO2 NWs on a single strand of C depicted in Figure 2c. For completeness, SEM images of the CFN are shown in Figure 2d−f. The surface area of the SnO2 NWs grown on a 10 mm × 10 mm CFN is ≈500 cm2, which is attractive for the deposition of oxides like PbO, Pb2O3, and PbO2 and the realization of energy storage devices. Properties of SnO2/PbO and SnO2/Pb2O3 NWs on Si, f-SiO2, and C Fibers. We obtained SnO2/PbO core−shell NWs via the deposition of Pb over the SnO2 NWs and annealing between 100 and 200 °C. The SnO2 NWs remain one-dimensional after annealing, as shown by the SEM images in Figure 3, and the resultant SnO2/PbO core−shell NWs were light yellow and exhibited clear and well-resolved peaks in the XRD, as shown in Figure 4, corresponding mainly to the tetragonal rutile crystal structure of SnO2 and tetragonal PbO. The valence state of Pb in PbO is +2 and depends on the annealing temperature. Higher temperatures between 300 and 500 °C favored the formation of SnO2/Pb2O3 and SnO2/ Pb3O4 NWs, which were orange and consisted of Pb2O3 and Pb3O4 with monoclinic and orthorhombic crystal structures. Pb3O4 consists of PbO and Pb2O3 in which Pb is in +2 and +3, respectively. High-magnification TEM images of a PbO/SnO2

Figure 3. SEM images of SnO2/PbO core−shell NWs obtained on Si(001) and fused SiO2 at (a) 200 °C and (b) 300 °C and also of SnO2/Pb2O3 NWs at (c) 400 °C and (d) 500 °C. Note that the SnO2 NWs are not uniformly covered by Pb, so one can clearly observe bare segments of SnO2 NWs in (d).

Figure 4. XRD traces of the SnO2/PbO and SnO2/Pb2O3 core−shell NWs obtained on Si(001) at 200, 300, 400, and 500 °C.

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Figure 5. TEM and HRTEM images of the SnO2/PbO core−shell NWs obtained on Si(001) at 200 °C. One may observe (a) the (200) plane of SnO2 and the (110) as well as the (002) crystallographic planes of PbO and (b) the (001) and (111) crystallographic planes of Pb2O3. The corresponding EDX spectrum is shown for completeness from which one may observe the peaks belonging to Pb and Sn but also O.

Figure 6. TEM and HRTEM images of the SnO2/Pb2O3 core−shell NWs obtained on Si(001) at 500 °C. (a) TEM image showing the bare SnO2 NW segments and Pb2O3. (b) HRTEM image showing the (211) crystallographic planes of SnO2. (c) and (d) HRTEM images showing the (210) crystallographic planes of Pb2O3.

and annealed, oxygen moves into the metal and reacts to form a metal oxide, whereas the metal moves into the oxide, leading to the formation of a ternary oxide.22 Recently, Ma et al.8 and Ganose et al.23 showed theoretically that substantial band gap tuning and a strain-controlled semiconductor-to-semimetal transition are possible via the incorporation of Pb in SnO2. Both SnO 2 /PbO and SnO 2/Pb 2O 3 core−shell NWs exhibited PL at 500 nm or 2.5 eV, as shown in Figure 7. We do not observe any shift of the PL emission to the red or infrared, suggesting a reduction in the energy band gap that

NW are shown in Figure 5 along with HRTEM images, showing not only the (200) crystallographic planes of SnO2 but also the (110) and (200) planes of tetragonal PbO. For completeness, high-magnification TEM images of the SnO2/ Pb2O3 core−shell NWs are shown in Figure 6 along with HRTEM images, showing not only the (211) crystallographic planes of SnO2 but also the (210) planes of Pb2O3. The formation of the PbO/SnO2, SnO2/Pb2O3, and SnO2/Pb3O4 core−shell NWs is attributed to the reaction of Pb with O2.21 However, when a metal is deposited on the surface of an oxide D


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NW anode and Pb cathode in H2SO4 (aq) the anode reaction is Pb (s) + SO42− (aq) → PbSO4 (s) + 2e− with a standard reduction potential of −0.36 V. The standard reduction potential of the reaction PbO (s) + 2H+ (aq) + 2e− → Pb (s) + H2O (l) is −0.44 V, so a cell with the total reaction Pb (s) + PbO (s) + 2H+ (aq) + SO42− (aq) → PbSO4 (s) + Pb (s) + H2O (l) would have a standard total potential of only −0.08 V with a small correction from the solution activities. In other words, the open-circuit potential would be very close to 0. Consequently, the SnO2/PbO core−shell NWs on the CFNs must be converted first into SnO2/PbO2 core−shell NWs by charging up in H2SO4 (aq) before they are used as a LAB. Synthesis and Properties of SnO2/PbO2 Core−Shell NWs. Highly conductive SnO2/PbO2 core−shell NWs were obtained via ELD of PbO2 on the SnO2 NWs that were grown on CFNs. The SnO2 NWs on the CFN that constituted the working electrode (WE) were immersed in a solution of 0.3 M HNO3 and 1 M Pb(NO3)2, and a potential of 3 V was applied versus a Ag/AgCl reference electrode (RE) using Pt as a counter electrode (CE), as shown in Figure 9. The ELD of

Figure 7. Room-temperature PL spectra of the SnO2/PbO and SnO2/ Pb2O3 core−shell NWs obtained by annealing of Pb deposited on SnO2 NWs. The inset shows absorption and transmission spectra through a 50 nm layer of Pb deposited on fused SiO2 after annealing between 200 and 500 °C.

might be related to the incorporation of Pb into the SnO2. Note that the PbO, Pb2O3, and Pb3O4 obtained by thermal annealing of Pb deposited on f-SiO2 have large optical band gaps, as can be seen in the inset of Figure 7, which means that the PL is related to radiative transitions in SnO2 (Supporting Information, Section S2). In addition, we did not observe any change in the conductivity of the SnO2 NWs after the deposition of Pb and thermal annealing. All of the SnO2/PbO and SnO2/Pb2O3 core−shell NWs on f-SiO2 had a small conductivity on the order of 10−1 S/cm, as shown by TCS in Figure 8. Note that

Figure 9. Schematic diagram of the ELD cell that was used for the deposition of PbO2 on the SnO2 NWs grown on C and also directly on the C. The lower graph shows the periodic variation of current with time that occurred during the deposition of PbO2 on both SnO2 NWs (red) and CFN (blue); and the top graph shows the XRD of PbO2 on the CFN, where the dominant peaks correspond to tetragonal PbO2 and a few smaller peaks to residual cubic Pb(NO3)2.

Figure 8. Time evolution of the electric field through the SnO2, SnO2/PbO, and SnO2/Pb2O3 core−shell NWs obtained by annealing of Pb deposited on SnO2 NWs. The inset shows the conductivity extracted from the fast Fourier transform.

PbO2 occurs via the following reactions: (a) H2O → OH• + H+ + e−, (b) Pb2+ + OH• → Pb(OH)2+, (c) Pb(OH)2+ + H2O − e− → Pb(OH)2 2+ + H+, and (d) Pb(OH)22+ → PbO2 + 2H+. It should be noted that the morphology of plain SnO2 NWs did not change after immersion in 0.3 M HNO3 (aq); typical SEM images of the SnO2/PbO2 core−shell NWs obtained after 1000 s of ELD are shown in Figure 10. The deposition of PbO2 on the SnO2 NWs changed its color from white to black, and a typical SEM image of PbO2 deposited by ELD at 3 V for 1000 s directly on a CFN is also shown in Figure 10, from which one may observe the formation of PbO2 in the radial direction. The thickness of the PbO2 shell is ≈20 μm, i.e., 2 times larger than the diameter of the C fiber. All of the SnO2/PbO2 core−shell NWs exhibited clear and wellresolved peaks in the XRD, as shown in Figure 9, corresponding to the tetragonal rutile crystal structure of SnO2 and tetragonal crystal structure of PbO2. The SnO2/ PbO2 core−shell NWs have a high crystal quality and high

the PbO, Pb2O3, and Pb3O4 layers that were obtained by the deposition of Pb on f-SiO2 and thermal annealing had very low conductivity, which is consistent with the findings of Droessler et al.24 (Supporting Information, Section S2), and as such they did not result in a change in the conductivity of the SnO2 NWs. From the above discussion, it is clear that we do not observe a semiconductor-to-semimetal transition that might be related to the incorporation of Pb into the SnO2 NWs by thermal diffusion. We have also tried to obtain Pb-doped SnO2 NWs via the VLS mechanism, but Pb was not incorporated at all in the SnO2 NWs (Supporting Information, Section S3). Despite the low conductivity of the SnO2/PbO and SnO2/ Pb2O3 core−shell NWs, they were also grown on CFNs because PbO and higher oxides like Pb2O3 are used as starting materials to make LABs. However, it should be noted here that in the case of a LAB cell consisting of a SnO2/PbO core−shell E


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Figure 10. (a)−(f) SEM images of the PbO2 that was deposited on CFNs and (g)−(i) PbO2 deposited on SnO2 NWs on f-SiO2 by ELD at 3 V for 1000 s.

conductivity or small resistivity on the order of 10−4 Ω cm due to the large charge carrier densities in the PbO2. To obtain a better understanding of the electronic properties of the SnO2/PbO2 core−shell NWs, we have calculated the conduction band potential profile and one-dimensional electron gas (1DEG) charge distribution along the radial direction via the self-consistent solution of the Poisson− Schrödinger (SCPS) equations, which is described in detail elsewhere.25 The SCPS calculations were carried out by taking into account the effective mass and dielectric constant of SnO2, i.e., me* = 0.326,27 and εr = 13.5,28,29 respectively. Furthermore, we have taken into account not only the electron affinity of SnO2, i.e., χ = 4.8 eV,30 but also the work function of stoichiometric SnO2, i.e., ϕ = 5.5 eV, according to Rachut et al.,31 which is in agreement with that from Klein et al.32 Note that higher values of ϕ = 7.74 eV have also been reported for SnO233 but the work function of SnO2 prepared under reducing growth conditions is smaller and varies between ϕ = 4.1 and 4.48 eV. Consequently, the work function of SnO2 that was taken into account for the SCPS calculations, i.e., ϕ = 5.5 eV, lies in the middle of the range of values listed above. The Fermi level was fixed at ≈0.7 eV below the conduction band edge at the surface of SnO2 without PbO2 in accordance with Kar et al.34 In addition to SnO2, we have taken into account the effective mass of electrons in PbO2 from the calculations of Scanlon et al.,9 which give me* = 0.18mo in contrast to others who estimated that me* = 0.8mo.35 The dielectric constant of PbO2 is εr = 5.1,36 and the work function of PbO2 was taken to be 5.6−5.8 eV.24 This gives a straddling type of band alignment between SnO2 and PbO2 by taking into account that the energy gap of PbO2 is ≈0.5 eV, which is considerably smaller than the energy band gap of SnO2, that is, 3.7 eV. The energy band diagram of a SnO2/PbO2 core−shell NW is shown in the inset in Figure 11. The conduction band discontinuity at the SnO2/PbO2 heterojunction interface is ΔEC ≈ 0.6 eV, which results in the confinement of a 1DEG distribution inside the PbO2 shell. However, a smaller shell

Figure 11. SCPS CB potential profile relative to the Fermi level, i.e., EC − EF, and 1DEG charge distribution versus the distance along the radial direction in a SnO2 NW. The inset shows the potential profile and charge distribution of the SnO2/PbO2 core−shell NW.

thickness and/or larger core radius results into a 1DEG confined in the core similar to that shown for SnO2 in Figure 11 despite the fact that the Fermi level is still close to the conduction band edge at the surface of the SnO2/PbO2 core− shell NW. In general, the deposition of a PbO2 shell with a thickness greater than 100 nm over SnO2 NWs with diameters of a few tens of nanometers and carrier densities on the order of 1016 cm−3 is expected to lead to a substantial increase of the conductivity and the confinement of carriers in the PbO2. For comparison, the energy band diagram and charge distribution of a SnO2 NW is shown in Figure 11. The CB edge potential of the SnO2 NWs is U-like, and the electric field is 0, i.e., a flat band condition exists at the core of the SnO2 NW. The 1DEG charge distribution has a maximum at the core and decays to 0 at a distance of 10 nm meaning that the depletion has a spatial extent of 15 nm. No sub-bands fall below the Fermi level, and the ground level sub-band is situated at 40 meV above the Fermi level, whereas the next one up is at 97 meV, giving a carrier density of 2 × 1016 cm−3. We obtain a 1DEG line density of 0.7 × 109 m−1, which translates F


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The Journal of Physical Chemistry C into a three-dimensional density of 2.2 × 1022 m−3 close to the value determined previously by TCS, suggesting that the SnO2 NWs are close to depletion.5 From the above discussion, it is evident that PbO2 can be deposited on SnO2 to tailor its properties such as the conductivity in accordance with the theoretical predictions of Butler et al.12 More importantly, the SnO2/PbO2 core−shell NWs on the CFNs may be combined with SnO2/Pb core− shell NWs or plain Pb to form a LAB using H2SO4 (aq). In this case, the following reactions occur in connection with PbO2 and Pb: (a) the cathode reaction is Pb (s) + SO42− (aq) → PbSO4 (s) + 2e− and its standard reduction potential at 25 °C is −0.36 V and (b) the anode reaction is PbO2 (s) + 4H+ (aq) + SO42− (aq) + 2e− → PbSO4 (s) + 2H2O (l); the standard reduction potential at 25 °C is +1.7 V. The overall reaction is Pb (s) + PbO2 (s) + 4H+ (aq) + 2SO42− (aq) → 2PbSO4 + 2H2O (l), giving a cell potential difference of 2.06 V. When the LAB is empty, both Pb and PbO2 have been converted into PbSO4. Upon charging up, the anode reaction is PbSO4 + 2e− → Pb (s) + SO42− (aq) and the cathode reaction is PbSO4 (s) + 2H2O (l) → PbO2 (s) + 4H+ (aq) + SO42− (aq). Therefore, the overall reaction is then 2PbSO4 + 2H2O (l) → Pb (s) + PbO2 (s) + 4H+ (aq) + 2SO42− (aq). After charging up, the PbSO4 is converted back to Pb and PbO2 at the negative and positive plates, respectively. The reaction of 3.87 mg of Pb and 4.47 mg of PbO2 (1.87 × 10−5 mol) with H2SO4 and their conversion to PbSO4 in a galvanic cell configuration will produce 1 mAh of charge. This gives a theoretical value of 224 mAh/g. The open-circuit potential of a half-cell consisting of PbO2/ SnO2 NW WE versus CFN CE and Ag/AgCl RE, which was measured versus time in the cell of Figure 9 using 5 M H2SO4 (aq), was ≈1.8 V, as shown in the inset in Figure 12. This is

al.15 observed a peak in the CV at ∼0.2 V related to the redox of SnO2 probably due to the exposure of the latter to the H2SO4 (aq) and also a peak at 1.8 V due to the redox of PbO2. We do not observe the peak related to the redox of SnO2. Consequently, the PbO2 deposited over the SnO2 NWs leads to a suppression of the redox reaction related to SnO2 nanoparticles, although it has been suggested that the SnO2 will still undergo reduction and oxidation due to electron transfer between PbO2 and SnO2 according to PbO2 − SnO2 + 2H+ + 2e− ⇔ PbO2 − SnO + H2O.37 Despite the fact that the PbO2 is deposited in a conformal fashion all over the SnO2 NWs, it is important to mention that the morphology of the SnO2 NWs without any PbO2 did not change upon immersion in 5 M H2SO4 (aq), which is consistent with the fact that the solubility of SnO2 is 1.4 mg/L in 4−8 M H2SO4 (aq).38 This means that the SnO2 NWs on the CFNs may act as robust scaffolds in strongly acidic electrolytes and will not dissolve away even upon current flow or the application of a potential. A typical CD curve obtained from the SnO2/PbO2 core−shell NWs versus Pb that was deposited on the SnO2 NWs is shown in Figure 12 from which we find a specific capacity of 206 mAh/g; this is slightly higher than that obtained by Moncada et al.,10,11 who used short PbO2 NWs to make a highperformance battery with an energy density of 190 mAh/g. We obtained a smaller energy density of 182 mAh/g from the SnO2/PbO core−shell NWs versus Pb after charging up and converting the anode to SnO2/PbO2. However, both SnO2 and SnO2/PbO2 core−shell NWs on CFNs obtained here may be used to obtain even higher specific capacities in view of the fact that Sb:SnO2 nanoparticles deposited on C fibers39 in a Li-ion battery exhibited a very high capacity of 705 mAh/g and an outstanding high rate performance attributed to the Sb:SnO2 nanoparticles.

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CONCLUSIONS

SnO2 NWs have been grown on Si, f -SiO2, and CFNs by the VLS mechanism at 800 °C and 10−1 mbar, whereas SnO2/PbO and SnO2/Pb2O3 core−shell NWs were obtained by the deposition of 50 nm Pb over the SnO2 NWs followed by annealing between 100 and 500 °C due to the oxidation of Pb. Both SnO2/PbO and SnO2/Pb2O3 core−shell NWs had a low conductivity on the order of 10−1 S/cm as shown by TCS due to the high resistivity of SnO2 as well as that of PbO and Pb2O3. However, the SnO2/PbO core−shell NWs can be converted into SnO2/PbO2 by charging in H2SO4 (aq) as PbO and higher oxides like Pb2O3 are commonly used as starting materials in LABs. In contrast, highly conductive SnO2/PbO2 core−shell NWs were obtained directly by ELD of PbO2 on SnO2 NWs. We show that a straddling type of band alignment exists between SnO2 and PbO2 and the 1DEG charge distribution exists in the PbO2 when it is sufficiently thick. The SnO2/PbO2 core−shell NWs are ideal candidates for the realization of a LAB using H2SO4 (aq), exhibit an open-circuit potential of 1.8 V versus CFNs in 5 M H2SO4, and show symmetric CV curves, suggesting a suppression of the redox reactions related to SnO2 and a specific capacity of 206 mAh/g, which is close to the theoretical value of 224 mAh/g. The SnO2 NWs on CFNs are even more promising for the realization of higher-specific-capacity Li-ion battery cells.

Figure 12. CD curve from a half-cell consisting of SnO2/PbO2 NWs versus SnO2/Pb NWs on CFN in 5 M H2SO4 (aq). The lower inset shows the OCP versus time for the SnO2/PbO2 NWs versus a CFN and Ag/AgCl RE measured in the cell of Figure 9; and the top inset shows the CV obtained in the same configuration used to measure the OCP.

attributed to the PbO2 that was deposited by ELD over the SnO2 NWs but also on the back side of the CFNs that are not covered by the SnO2 NWs. The CFN itself has a high-metallike conductivity and does not react with the 5 M H2SO4 (aq); in other words, it does not contribute toward the open-circuit potential and merely acts as a current collector. We obtained symmetric CV curves, as shown in the inset in Figure 12, from the same half-cell configuration used to measure the opencircuit potential. This configuration is similar to that of Dan et al.15 who prepared a PbO2/SnO2 nanoparticle composite by ELD on Ti and used activated C as a CE. However, Dan et G


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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.8b07526. Detailed description of (a) the properties of the CFNs, (b) electrical and optical properties of PbO and Pb2O3 obtained on f-SiO2, and (c) Pb doping of SnO2 (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Matthew Zervos: 0000-0002-6321-233X Epameinondas Leontidis: 0000-0003-4427-0398 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors at the Polytechnic University of Bucharest acknowledge EU-funding POSCCE-A2-O2.2.1-2013-1/Axa Prioritara 2, Project No. 638/12.03.2014, Code SMIS-CSNR 48652.

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