Letter pubs.acs.org/NanoLett
Electron-Rich Driven Electrochemical Solid-State Amorphization in Li−Si Alloys Zhiguo Wang,*,†,‡ Meng Gu,‡ Yungang Zhou,‡ Xiaotao Zu,† Justin G. Connell,§ Jie Xiao,‡ Daniel Perea,‡ Lincoln J. Lauhon,§ Junhyeok Bang,∥ Shengbai Zhang,∥ Chongmin Wang,*,‡ and Fei Gao*,‡ †
Department of Applied Physics, University of Electronic Science and Technology of China, Chengdu, 610054, P. R. China Pacific Northwest National Laboratory, P.O. Box 999, Richland, Washington 99352, United States § Materials Science and Engineering, Northwestern University, Evanston, Illinois 60208, United States ∥ Department of Physics, Applied Physics, & Astronomy, Rensselaer Polytechnic Institute, Troy, New York 12180, United States ‡
S Supporting Information *
ABSTRACT: The physical and chemical behaviors of materials used in energy storage devices, such as lithium-ion batteries (LIBs), are mainly controlled by an electrochemical process, which normally involves insertion/extraction of ions into/from a host lattice with a concurrent flow of electrons to compensate charge balance. The fundamental physics and chemistry governing the behavior of materials in response to the ions insertion/extraction is not known. Herein, a combination of in situ lithiation experiments and large-scale ab initio molecular dynamics simulations are performed to explore the mechanisms of the electrochemically driven solid-state amorphization in Li−Si systems. We find that local electron-rich condition governs the electrochemically driven solid-state amorphization of Li−Si alloys. This discovery provides the fundamental explanation of why lithium insertion in semiconductor and insulators leads to amorphization, whereas in metals, it leads to a crystalline alloy. The present work correlates electrochemically driven reactions with ion insertion, electron transfer, lattice stability, and phase equilibrium. KEYWORDS: Li−Si alloys, electrochemical solid-state amorphization, ab initio molecular dynamics simulations, in situ TEM, electron rich
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et al.2 demonstrated that an amorphous Li−Si alloy forms when Si is electrochemically lithiated, which is considered as an electrochemically driven solid-state amorphization process. Although the amorphization process is assumed to be closely analogous to the diffusive solid-state amorphization observed nearly 20 years ago,15 the detailed mechanisms of the amorphization are still not established. Numerous studies have examined different LixSi phases, and several density functional theory (DFT) calculations from various groups have been performed to understand the mechanism of the structural changes and the lithiation behavior of Si-based materials.13,16−19 Experimental results from various groups have clearly shown that crystalline intermetallics have a much lower Gibbs energy as compared to the corresponding amorphous alloys. DFT calculations have shown that the formation energies of amorphous alloys are much higher than those of the corresponding crystalline alloys,11,17 which we show in Figure S1 of the Supporting Information. The longstanding question is why and how the amorphous LixSi phases, except the Li3.75Si phase, occur rather than forming the stable
tructural evolution of materials driven by electrochemical processes is not necessarily represented by a classic thermodynamically driven process. It has been well-documented that the electrochemically driven insertion of lithium ions into a semiconductor or insulator will lead to amorphization (typically Li into Si, ZnO2, SnO2, and Ge),1−4 a process that is not attainable with other known solid-state or wet-chemistry synthesis methods. Some materials, instead of being subjected to such a solid-state amorphization, directly alloy with the inserted ions to form crystalline new phase (typically Li into Al and Sn).5,6 The Li−Si phase diagram, based on thermal treatment of a mixture of Li and Si, presents a rich variety of solid crystalline compounds (LiSi, Li12Si7, Li13Si4, Li15Si4, and Li22Si5denoted as LiSi, Li1.71Si, Li3.25Si, Li3.75Si, and Li4.4Si in the following study).7 However, these compounds do not form during the electrochemical alloying of crystalline Si with Li.2 The phase transitions of LixSi alloys during Li insertion into Si have been investigated by scanning electron microscopy,8 in situ X-ray diffraction (XRD),9 ex situ XRD,10 in situ transmission electron microscopy (TEM),11−13 and real time nuclear magnetic resonance.14 It is widely accepted that, during the first lithiation cycle, crystalline Si undergoes a phase transition to form an amorphous LixSi, but highly lithiated amorphous LixSi can only recrystallize to Li3.75Si. Limthongkul © 2013 American Chemical Society
Received: July 2, 2013 Revised: August 2, 2013 Published: August 14, 2013 4511
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Figure 1. (a and b) STEM image showing the formation of a core−shell structure at the very early stage and after partial lithiation of the whole NW, respectively; (c) HRTEM image of the interface between amorphous LixSi shell and crystalline Si core during in situ lithiation as labeled by the white arrow; (d) schematic atomic model of the interface between amorphous LixSi and crystalline silicon core.
model the computational supercell with a finite net charge, a compensating background is used to achieve charge neutrality.23 This is identical to that obtained from a similar system, which consists of the original charged system immersed in a jellium background which fills the supercell and neutralizes the charge, so that the net charge is zero.24 The amorphization and crystallization of the Li15Si4 phase is characterized in detail by in situ TEM, scanning TEM (STEM), and electron energy-loss spectroscopy (EELS) using intrinsic and P-doped Si nanowires (Si-NWs). As shown by the STEM Z-contrast image in Figure 1a and b, the lithiation of a singlecrystalline Si-NW proceeds as a core−shell structure with a crystalline Si core and an amorphous LixSi shell. For Z-contrast imaging in STEM, the contrast is integral to the electrons elastically scattered in the high angle annular dark field detector (HAADF), reflecting the average atomic number of the region.25 Therefore, the crystalline Si core shows higher intensity than the amorphous LixSi region because ZSi > Z(LixSi). In addition, Li+ is transported from the Li metal (bottom of Figure 1a) to Si-NW and lithiation propagates from Li/Li2O end and gradually moves upward. In Figure 1a, the bottom of the NW is partially lithiated and forms a core−shell structure, while the top part of the Si-NW remains intact. As lithiation proceeds, the top part of the Si-NW also begins to lithiate and forms the core−shell structure as well (Figure 1b). Concurrently, the diameter of the top part of the Si-NW increases from 88 nm (Figure 1a) to ∼136 nm after partial lithiation (Figure 1b). The lithiation process is recorded in situ and shown in Movie S1 in the Supporting Information. The high-resolution TEM (HRTEM) lattice image in Figure 1c indicates that the crystalline silicon is projected along the [111] zone axis, and the interface between the crystalline Si and amorphous LixSi aligns to the Si {1−12} planes. In addition, the HRTEM image in Figure 1c and corresponding atomic model clearly demonstrate that the interface between the amorphous LixSi shell and crystalline silicon core is atomically sharp. The distribution of Li within the LixSi is qualitatively viewed through in situ EELS mapping using an aberration-corrected STEM. During an in situ lithiation experiment, the Si-NWs at different lithiation states were analyzed by Z-contrast imaging and EELS mapping as shown in Figure 2a−c. In the early lithiation state in Figure 2a, the silicon core showed higher Zcontrast compared to the outer LixSi shell. The Li, Si, and overlaid Li and Si composite maps in Figure 2a identified the
crystalline phases (x = 1, 2, 3...) during lithiation. On the other hand, Li3.75Si phases have not been synthesized by other experimental methods, other than electrochemical lithiation, as far as we are aware. Overall, there is a lack of fundamental understanding of the factors controlling the solid-state amorphization process during electrochemical reactions. For the present study, both nominally intrinsic and phosphorus doped Si NWs were grown using chemical vapor deposition, and their lithiation speeds were compared. The in situ S/TEM measurements were performed on aberrationcorrected FEI Titan microscopes using a nanofactory STM holder. The electron energy loss spectroscopy (EELS) data was acquired by using Gatan Image Filter (Quantum 965) with a 2k × 2k CCD. The lithiation of the Si NWs is observed in situ by assembling a nanobattery inside the TEM with Li metal being used as the Li source in order to obtain lattice-resolution images during in situ lithiation process. The Li2O layer formed on the lithium metal surface was used as a solid electrolyte. The Si NWs were grown on a silicon substrate and doped with phosphorus to achieve high electronic conductivity. After forming a close circuit by connecting the Si NW to the Li/Li2O end, lithiation usually begins spontaneously. The lithiation process was illustrated by the STEM Movie S1. All of the calculations were performed using density functional theory, within local density approximation.20 Ceperly-Alder parametrization as implemented in the SIESTA code21 was used, which adopts a linear combination of numerical localized atomic orbital basis sets for the description of valence electrons and norm-conserving nonlocal pseudopotentials for the atomic core.22 The valence electron wave functions were expanded by using double-ζ basis set. The charge density was projected on a real space grid with a cutoff of 150 Ry to calculate the self-consistent Hamiltonian matrix elements, and the Γ point was used in the Brillouin zone sampling. The crystalline alloys were constructed from the crystalline structure database and experimental results. Large systems were used to model the structure and evolution. The numbers of atoms used for the Li−Si alloys are: LiSi (128/ 128), Li1.71Si (192/112), Li3.25Si (156/48), Li3.75Si (240/64), and Li4.4Si (352/80). The evolution of the system is involved using the MD method with verlet algorithm for all the atoms after adding extra electrons, with a time step of 1.0 fs. Fermi distribution is used for the added electronic charges to be placed into the empty states. If the crystal remains its perfect configurations, a long simulation time up to 10 ps is used. To 4512
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gradual increase of thickness from the surface to the center of the nanowire can directly affect the EELS signal. The increase of the thickness will lead to plural scattering, affecting the edge jump and fine structure, and increase background intensity. However, our sample thickness is in the range of ∼0.4 times of the electron mean free path (Figure S5), for which the thickness effect on the signal is not significant. Further, we noticed that, after the full lithiation, the Li distribution becomes uniform as shown in Figure 2c, indicating that the Li-rich characteristic at the reaction interface before reaching full lithiation is not a thickness effect. This result is also statistically supported by the maps collected on nanowires with different thicknesses as illustrated in the Supporting Information (Figures S6, S7, and S8). With further enrichment of Li in LixSi, the Li3.75Si phase forms, as evidenced by the electron diffraction pattern shown in the inset of the STEM image in Figure 2c. To continue the charge balance of the system, the enrichment of Li + concentration at the interface between LixSi and Si (Figure 2a and b) will lead to the outer portion of the silicon core (adjacent to the reaction front) to have a higher electron density. Therefore, our analyses strongly suggest that the interface or reaction front region is under an electron-rich condition during lithiation. Furthermore, we find that the lithiation speed of highly phosphorus-doped Si-NWs exceeds the intrinsic Si-NWs by ∼200 times. Similar results have been previously reported.13 With a better electron conductivity of phosphorus-doped Si-NWs, the electrons accumulate at the interface more quickly than those in intrinsic Si-NWs, leading to high-speed lithiation. We have performed climbing image nudged elastic band (CI-NEB) calculations,26 as implemented in the Vienna ab initio package (VASP),27 to investigate the effect of P-doping on the diffusion of Li in a Si lattice. We found that the diffusion barriers of Li in intrinsic and P-doped Si are 0.55 and 0.53 eV, respectively, as shown in Figure S2. It is clear that the P-doping does not affect the diffusion of Li but
Figure 2. (a, b, and c) STEM Z-contrast image, Li maps (52.5−84.5 eV), Si maps (96.25−210.25 eV), and Li/Si composite maps showing the distribution of Li and Si in different lithiation stages. Note the Li distribution profile as the inset of Li maps in a−c shows the Li distribution profile across the Si-NW, indicating a Li-rich interface. The electron diffraction pattern in c indicates the formation of Li15Si4 phase after full lithiation of the Si-NW. The EELS maps are acquired at different nanowire regions to avoid possible beam damage (EELS acquisition parameters: 0.25 eV/channel, 0.005 s/pixel, 300 kV).
lithium-ion distribution within the NWs. Most noticeably, the Li map in Figure 2a showed an enhanced Li+ concentration at the interface between LixSi and Si. The EELS maps of Li distribution at the middle state of lithiation further revealed that this interfacial Li+-rich condition is maintained until the Si-NW is fully lithiated (Figure 2b). It should be pointed out that the
Figure 3. Structural snapshots of Li1.71Si at 300 K with Ne = 2 at different simulation times and time evolutions of the pair correlation function upon amorphization of Li1.71Si at 300 K with Ne = 2, where the red spheres indicate Li ions and the green spheres Si atoms. 4513
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the simulation time, the sharp peaks at the distances larger than 6 Å disappear at the simulation time of 90 fs, and the longrange order disappears, forming a disordered phase. We also note that, by further increasing the simulation time, the shortrange order remains consistent with an amorphous state, as shown by broadening the peaks below 3.0 Å. At higher temperatures (500 K), all of the LixSi alloys transform to an amorphous structure with N e = 2, demonstrating that extra electrons indeed play an important role in the electrochemical process. Figure S3 shows the atomistic structures of LiSi and Li3.75Si alloys annealed at 500 K for 2.5 ps with Ne = −1, 0, and 2. For the crystalline LiSi, 3-fold coordinated Si ions form the interconnected chains and pucker with eight-membered rings, while cavities in the network are occupied by Li ions. Crystalline Li3.75Si is formed by six Li atoms with a surrounding Si atom. Both LiSi and Li3.75Si alloys transform to an amorphous structure with Ne = 2 at 500 K. All of the Li−Si alloys investigated with DFT show the same amorphization process with an excess of electrons: Si−Si bond breakage and amorphization of the Li−Si alloys. The effect of temperature and Ne on the stability of Li−Si alloys was investigated, and the PCFs are shown in Figure S4. At 500 K, crystalline LiSi and Li1.71Si with Ne = 1 transform to amorphous structures after a simulation time of 2.5 ps, while Li3.25Si, Li3.75Si, and Li4.4Si retain their long-range order, at least for a simulation time of ∼10 ps. However, when Ne is increased to 2, all of the alloys transform to amorphous phases within 2.5 ps, and all of the crystalline Li−Si alloys are unstable with sufficient electron enrichment. These results clearly indicate that the electron-induced solid-state amorphization depends on the composition of LixSi alloys and electron density. As the temperature is decreased to 300 K, Li3.75Si can retain its longrange order, even with Ne = 2, as shown in Figure S4c. All of the other alloys become amorphous at much shorter simulation times. Figure S4d shows the structures of Li3.75Si at temperatures ranging from 500 to 900 K with Ne = 1. It is clearly seen that Li3.75Si can be stable up to 900 K. These results suggest crystalline Li3.75Si is the most stable phase among all of the Li− Si alloys, which agrees with the fact that Li3.75Si is known to be the fully lithiated phase at room temperature.32,33 This also explains why only crystalline Li3.75Si has been observed during lithation.30,34 Nevertheless, it should be pointed out that all crystalline Li−Si alloys are unstable at high temperatures and high electron densities (e.g., Ne = 2 at 500 K). Compositions of intermetallic phases and related solid-state systems are often successfully estimated using an electron counting rule ((8 − N) rules) developed by Zintl.35,36 The formal counting rule strategy corresponds to the assignment of the valence electrons of the electropositive atoms to the valence bond (VB) type states of the electronegative component, allowing for the formation of an octet configuration of the latter. The Li−Si alloys are likely to be correlated to the Zintllike phases.37 In Li−Si alloys, the Si atoms is more electronegative than Li atoms; thus the Si atom accepts electrons from the Li atoms. If all of the Li atoms donate their valence electrons to Si atoms, the charge transfer in LixSi can be represented by (Li+)xSix−. There are four valence electrons for an isolated Si atom, which means that the isolated Si atom can accept four electrons at most. The armorphous and crystalline atomic configurations are shown in Figure S1. The Si connectivity decreases with increasing Li content: 3-fold coordinated Si ions forming interconnected chains puckered with eight-membered rings in LiSi, 5-Si-rings, and 4Si-“Y” in
will enhance the transportation of electrons, creating better electron conductivity of phosphorus-doped Si-NWs. These results clearly suggest that the formation of an electron-rich environment plays a dominant role in the lithiation process. During lithiation, electrons transfer to the Si through the external circuit. As shown previously, the silicon lithiation process is ultimately limited by the interface reaction, involving the rate of breaking of Si−Si bonds and formation of a Si−Li bond.28−30 According to this reaction controlled theory, Li+ can diffuse easily through the LixSi shell layer and accumulate at the interface between LixSi and Si core. EELS mapping of Li distribution does not directly reveal any charge distribution. An electron-rich condition can be produced to counter the Li ion enrichment at the reaction front between amorphous LixSi and crystalline Si. The total circuit remains charge neutral, but the reaction interface between LixSi and Si core is electron-rich. In principle, such a positive−negative charge layer at the reaction front can be measured using electron holography. Further detailed work is needed along this direction. The Li3.75Si formation is observed only in the electrochemical lithiation of Si. No other methods, such as solid-state reaction or wet chemistry, can produce a crystalline Li3.75Si phase. As discussed previously, the uniqueness of the electrochemical method is due to the inducement of an electron-rich environment in the Si-NWs, which is not present in solidstate and wet-chemistry methods. To better understand the process of electrochemical solid-state amorphization, the experimental observation is further explored using theoretical calculations of the crystalline lithium silicide phase transformation (LiSi, Li1.71Si, Li3.25Si, Li3.75Si, and Li4.4Si) under electron-rich conditions using a large-scale DFT molecular dynamics (DFTMD) method. The number of extra electrons (Ne) added per LixSi is used to characterize the electron density. Thus, Ne = 0 represents the neutral state without extra electrons, while Ne > 0 and Ne < 0 represent the states with electrons added to or removed from the system, respectively. Simulation results show that the LixSi retain their initial configurations without structure changes with Ne = −1 and 0 at temperatures ranging from 300 to 600 K. However, all of the LixSi alloys, except Li3.75Si, transform to an amorphous structure with Ne = 2 at room temperature (300 K). To understand the structural changes at various stages of amorphization, in Figure 3 we show several snapshots of the structural evolution as a function of time for Li1.71Si at 300 K with Ne = 2. A Li1.71Si crystal has planar 5-Si-rings and planar 4Si “Y” (three-pointed planar Si4 stars), each of which forms distinct polyanions.31 From the direct observation of the atomic configurations, it is clear that the atoms vibrate around its crystal lattice before 40 fs, and the 5-Si-rings and planar 4-Si ’Y’ remain at their initial arrangements. However, some Li atoms leave their lattice positions at 90 fs. The Li1.71Si begins to lose its crystallinity at 140 fs, transforming to an amorphous phase, increasing the annealing time. It is of interest to note that the amorphization started from the disordering of Li atoms, followed by the Si−Si bond breakage taking place with increasing simulation time, which may be due to the fast diffusion of Li compared with Si. However, Si−Si dumbbells can be clearly observed in the amorphous Li−Si alloy. To provide confirmed evidence of the crystal-to-amorphous transition, the pair-correlation function (PCF) is calculated and shown in Figure 3. The PCF curve remains during the apparent sharp peaks for simulation time ∼40 fs, indicating that the Li−Si alloy is still in its crystalline structure. By increasing 4514
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amorphization in LixSi alloys, under electron-rich conditions, may be caused by the softening of the Si−Si bonds. On the other hand, the isolated Si can accept more electrons, as described previously, which allows for the crystalline Li3.75Si alloy to accept more electrons than other Li−Si alloys. Therefore, the crystalline Li3.75Si alloy is the most stable configuration among all of the Li−Si alloys considered during lithiation. Based on experimental observations and large-scale DFTMD simulations, we have proposed that the electron-induced solidstate crystalline−amorphous transition is a key mechanism controlling the phase transition in electrode materials. Phenomenologically, the amorphization appears to be caused by the weakening of Si−Si bonds in the presence of high concentration of Li at the crystalline−amorphous interface.14 In essence, the Si−Si bond weakening is directly related to electrons. With Li being introduced to Si, the charge from Li is transferred to Si to fill up the antibonding sp3 states of Si, which in turn weakens the corresponding Si−Si bonds. Therefore, the Si network is destabilized by lithiation, leading to the formation of various Li−Si amorphous phases.38 This approach has been applied to study phase transformation of Li−Si alloys during lithiation, as an example, because the electrochemically stable states of Li−Si alloys at room temperature may not correspond to thermodynamically stable states. Using in situ STEM, we directly observed the amorphization of Si nanowires and the formation of Li3.75Si domains during lithiation. The Li distribution profile at the interface suggests that the reaction front region of Li−Si alloys is under an electron-rich condition. To adapt to the electron-rich environment, Si−Si bond breakage occurs, forming isolated Si in these amorphous alloys. It has been observed that when x in the amorphous LixSi reaches a critical value of 3.75, the amorphous Li3.75Si transforms to crystalline Li3.75Si. This phase transformation is characterized by a congruent process.39 The Li3.75Si crystalline phase is very stable and can accept all the electrons from the inserted Li atoms without completely filling their 3p orbital. The structural stability under an electron-rich environment gives insights as to why the crystalline Li15Si4 phase is stable, and electron-induced bond-breakage provides the mechanism driving the electrochemically solid-state amorphization.
Li12Si7, Si2 dumbbells in Li3.25Si, and isolated Si in Li3.75Si, and Li4.4Si. In crystalline LiSi, each Si atom has three Si neighbors and receives one electron from the Li atom, allowing it to fill its octet. The Li12Si7, also denoted as Li24Si14, has two planar 5-Sirings and a planar 4-Si “Y”. The 5-Si-ring can receive six electrons from Li atoms, providing a total of 26 valence electrons. The “Y” can receive eight electrons from the Li atoms. Therefore, it has a total of 24 valence electrons. In Li13Si4, half of the Si atoms are in a dumbbell configuration, while the other half of Si atoms are surrounded by the Li atoms. The Si dumbbells can receive four electrons, giving it a total of 12 valence electrons. The isolated Si receives four electrons, resulting in a total of eight valence electrons. It is of interest to note that the Li15Si4 phase has only isolated Si atoms. If all of the Li atoms donate their electrons, the average charge on the Si atoms would be −3.75e, and the Li15Si4 can accept more electrons. However, the Li22Si5 phase also has only isolated Si atoms. Once the Si atoms have filled their valence orbitals, there are still two extra electrons, and the Li22Si5 cannot accept more electrons. For the armorphous LixSi alloys, the Si connectivity decreases compared with crystalline LixSi alloys, which can accept more electrons than crystalline ones. So LixSi alloys prefer the armorphous under electron-rich conditions. With the decreasing of the Si connectivity, the Si−Si interaction will decrease, which can be seen from the decreasing of the s−p hybridization. The density of states (DOS) projected on Si atoms in the crystalline and amorphous Li−Si alloys are shown in Figure 4.
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ASSOCIATED CONTENT
S Supporting Information *
Atomistic configuration, diffusion energy barriers, pair correlation functions, and STEM images (pdf), plus a movie showing a STEM-HAADF image during lithiation of silicon to a-LixSi. This material is available free of charge via the Internet at http://pubs.acs.org.
Figure 4. Density of states (DOS) projected on Si atoms in various crystalline and amorphous Li−Si alloys, and the evolution of DOS for the 3s and 3p states of the Li1.71Si alloy as a function of simulation time at 300 K with Ne = 2, where the Fermi energy level is set to zero.
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The Fermi energy level is set to zero, which is used as a reference. As the Li fraction changes from 1 (LiSi) to 4.4 (Li4.4Si) per Si in crystalline Li−Si alloys, the splitting between the 3s and 3p states grows larger, and the distribution of the 3s and 3p states becomes narrower, which is attributed to the decreasing Si−Si bond interactions. The changes of the 3s and 3p states in the amorphous Li−Si alloys also show the same trend. On the bottom of Figure 4, we demonstrate the changes of the 3s and 3p states of the Li1.71Si alloy as a function of simulation time at 300 K with Ne = 2. The multipeaks of the 3s and 3p states merge into one peak with increasing simulation time, indicating the Si−Si bond breakage. The solid-state
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected] (Z.W.); chongmin.wang@pnnl. gov (C.M.W.);
[email protected] (F.G.). Author Contributions
Z.W. and M.G. contributed equally. Z.W., F.G., M.G., and C.W. designed this study. Z.W. and F.G. carried out simulations, while M.G. and C.W. performed the in situ electrochemical experiments and TEM analysis. Z.W., F.G., M.G., and C.W. analyzed and discussed the results and wrote the manuscript. X.Z., J.X., J.G.C., D.P., L.J.L, J.B., and S.Z. contributed to the 4515
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scientific discussion. All authors read and commented the manuscript. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The work described in this paper is part of the Chemical Imaging Initiative at Pacific Northwest National Laboratory (PNNL). It was conducted under the Laboratory Directed Research and Development Program at PNNL, a multiprogram national laboratory operated by Battelle for the U.S. Department of Energy (DOE). The work was conducted in the William R. Wiley Environmental Molecular Sciences Laboratory, a national scientific user facility sponsored by DOE’s Office of Biological and Environmental Research and located at PNNL. PNNL is operated by Battelle for DOE under Contract DE-AC05-76RLO1830. Nanowires grown for this study were synthesized at the Center for Integrated Nanotechnologies, a DOE, Office of Basic Energy Sciences user facility at Los Alamos National Laboratory (Contract DE-AC5206NA25396), and Northwestern University (NSF DMR1006069). Work at Northwestern University was supported by NSF DMR-1006069. J.B. and S.Z. were supported by DOE under Grant No. DE-SC0002623.
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dx.doi.org/10.1021/nl402429a | Nano Lett. 2013, 13, 4511−4516
