Nanoscale characterization of ion mobility by temperature controlled Li

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Nanoscale characterization of ion mobility by temperature controlled Li-nanoparticle growth Valon Lushta, Dirk Dietzel, Bernhard Roling, and Andre Schirmeisen ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b16281 • Publication Date (Web): 08 Jan 2019 Downloaded from http://pubs.acs.org on January 11, 2019

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ACS Applied Materials & Interfaces

Nanoscale characterization of ion mobility by temperature controlled Li-nanoparticle growth Valon Lushta,†,‡ Dirk Dietzel,† Bernhard Roling,‡ and André Schirmeisen∗,† †Institute for Applied Physics, Justus-Liebig-Universität, 35392 Gießen, Germany ‡Department of Chemistry, University of Marburg, Hans-Meerwein-Strasse 4, 35032 Marburg, Germany E-mail: *[email protected]

Abstract

the AlPO4 -phase, while any influence of grain boundaries is related to subsurface constrictions of the current paths.

Detailed understanding of electrochemical transport processes on the nanoscale is considered not only as a topic of fundamental scienctific interest but also as a key to optimize material systems for application in electrochemical energy storage. A prominent example are solid state electrolytes, where transport properties are strongly influenced by the microscopic structure of grain boundaries or interface regimes. However, direct characterization of ionic transport processes on the nanoscale remains a challenge. For a heterogeneous Li+ -conducting glass ceramic we demonstrate quantitative nanoscopic probing of electrochemical properties on the basis of temperature controlled growth of nanoscopic Li-particles with conductive tip atomic force microscopy. The characteristic energy barriers can be derived from of the particle growth dynamics and are consistent with simultaneously recorded nano-voltammetry, that can be interpreted as an interplay between overpotentials, ion-conductivity and nanoscale spreading resistence. In the low temperature limit at around 170 K, where the particle growth speed is slowed down by several orders of magnitude with respect to room temperature, we demonstrate ion-conductivity mapping with lateral resolutions only limited by the effective tip-surface contact radius. Our mapping measurements reveal the insulating character of

Keywords: solid state electrolyte, AFM, lithium ion conductor, lithium ion conducting glass ceramic, batteries, ion migration, Li reduction

Introduction Performance optimization of modern electrochemical systems require a profound microscopic understanding of the underlying thermodynamics of the electrochemical processes. One particular example are solid state batteries, which are considered as a promising route to safe and reliable electrochemical energy storage based on solid state electrolytes. 1–6 However, widespread application of solid state batteries is currently hindered by insufficient ion conductivity of the solid electrolytes. One strategy to overcome this barrier is optimization of their electrochemical properties by nanostructuring, since grain boundaries or interface regimes in ion conducting materials have shown strongly enhanced transport properties. 7–14 This approach requires the exact knowledge of ion transport properties at the nanoscale. Inspired by established techniques like conductivity spectroscopy using microelectrode contacts, 15 current developments employ conductive atomic force microscopy (AFM) where

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a voltage between the nanoscale tip and the solid ion conductor sample is applied to induce localized ion diffusion. Different strategies have been followed to detect the subsurface ion movement, like e.g. detecting surface expansion in electrochemical strain microscopy (ESM) 9,16–24 or measuring electrostatic forces related to ion diffusion underneath the tip in time domain electrostatic force spectroscopy (TDEFS). 11,25–27

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lithium ion conductive glass ceramic (LI-CGC), it was found that these processes can lead to the formation of stable metallic Li-particles on the sample surface. 28,29 Analyzing the growth process of such particles in combination with monitoring the Faradaic current is one of the most direct methods to explore nanoscale electrochemistry. Scanning probe microscopy assisted nanoparticle growth was applied to different materials in recent years with special focus on the onset of particle growth, 30–32 the geometrical growth characteristics, 29,33 the reversibility of the growth process, 33,34 and the influence of back electrodes and atmospheric conditions. 35 Ultimately, these studies lead to a concept about the reaction kinetics of the nanoparticle growth process, where depending on the applied overpotentials between tip and substrate different regimes of nucleation, reaction limited growth or transport limited growth are anticipated. 35 In this context, scanning tunneling microscopy (STM) techniques have proven to be particularly suitable to induce nanoionic effects, as was e.g. demonstrated by analyzing the particle nucleation process on an atomic level 30 or more recently by active control of dopant concentration in solid electrolyte AgSnanodots, 36,37 with processes similar to the one relevant for quantized conductance atomic switching. 38 AFM techniques on the other hand are usually more suitable to analyze transport limited processes. In this work, e.g., nanoparticles of varying size were produced by electrochemical Li-reduction, which benefits from contact mode AFM operation during fast particle growth Electrochemical reaction kinetics are strongly dependent on temperature, however, experiments of particle growth as an explicit function of temperature are rare. 30 This imposes two significant limitations: First, it is not possible to extract quantitative energy barriers from the particle growth process at one temperature alone. Instead, determining quantitative values involves temperature dependent measurements, where characteristic activation energies can be derived from Arrhenius representations of the data. 11,27,39 Secondly, experiments at room temperature predominantly lead to the

Figure 1: Schematic depiction of the measurement setup. The tip of an electronicallyconductive AFM cantilever is in contact with the sample surface, a lithium ion conductive glass ceramic (LICGC). Applying a sufficiently high bias voltage between the tip and the sample induces an electric field in the sample and forces mobil ions to move towards the tip, where they are reduced and form metallic Li particles on the surface. The sample is mounted on a variable temperature stage, which allows to control the sample temperature in a range between room temperature and 40K One particularly intriguing aspect about AFM based analysis of solid state electrolytes is that it essentially mimics a conventional electrochemical set-up with a nano-electrode in contact with an electrolyte. This allows transferring classical electrochemical approaches to nanoscale systems. Fig. 1 shows the conductive AFM tip in contact with sample. Redox reactions can be induced by applying a cathodic overpotential between tip and sample, which leads to a reduction of mobile ions at the solid electrolyte / tip interface and the formation of metallic particles. In previous experiments on a


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growth of larger particles with contact areas up to a few µm2 as observed previously. 28,29 This limits the ability to achieve high spatial resolution for mapping purposes, unless more complex strategies are applied to control the particle sizes by e.g. using either current I or height ∆z as a bias interrupting trigger. 31 Nonetheless, nanoscale particle mapping with grid sizes of a only a few 10nm with still separate particles has not yet been demonstrated. Here we present experiments of AFM assisted growth of nanoscale Li-particles on LICGC in a wide temperature range. Our method allows simultaneous in-situ monitoring of particle size, Faradaic current and nucleation threshold. Based on these three parameters we identify different growth regimes and find that particle growth is exponentially dependent on the sample temperature, while the voltage for nucleation stays constant. The subsurface migration energy barrier is independently determined by Faradaic current and particle growth analysis with matching results. However, the detection limit of the Faradaic current analysis is two orders of magnitude inferior to the particle size analysis, where the smallest identifiable particle consists of less than 15000 atoms. These nanoscopic particle sizes can only be produced at cryogenic temperatures, which slow down the migration processes. This low temperature approach is exploited for systematic ion conduction mapping on the LICGC surface with a lateral resolution below 25nm, essentially only limited by the AFM tip radius.

Figure 2: Average threshold voltage for the formation of metall Li particles calculated from the formation of Li-nanoparticles at 6 different temperatures between 180K and 280K. No temperature dependence is evident from the threshold values. In each case, the starting point of the growth process has been determined from the normal force feedback signals recorded during the linear voltage sweep. An example curve measured at T = 280 K is shown as inset. Here, the normal force feedback signal remain zero until the bias voltage U reaches the threshold value of Uth = 3.8 V and a sudden increase of the feedback signal occurs, which allows to pinpoint the onset of particle growth potential from the maximum of the derivative of the normal force curve. After that the feedback signal remains at a non-zero value required to compensate the topography’s z-shift during the continuous particle growth.

RESULTS AND DISCUSSION

tween 180K and 280K. To grow Li-particles on the sample surface, the AFM was operated in contact mode, while the voltage U was ramped up within 1s from 0V to -5V and switched off immediately, once the maximum voltage was reached. A precise determination of the onset of particle growth can be made based on the AFM normal force feedback signal (Fig. 2, inset) recorded during the growth process. Prior to the onset of growth, the feedback signal remains zero. Only when the particle starts to grow, the feedback signal shows a steplike increase and remains

We have performed temperature dependent measurements of the particle growth process on a standard LICGC sample (AG01) with a thickness of 150 µm, that was purchased from Ohara Corp. Our first objective was to identify the characteristic activation energy related to the Li-particle formation. Experiments were performed with an Omicron VT-AFM under ultrahigh vacuum (UHV) conditions, where the sample was analyzed at 6 different temperatures be-


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Figure 3: Temperature dependence of Li-particle growth. a) Examplary topography image of Li-particles grown by applying voltage ramps in a 5x5 grid at 240K. The missing particles in the upper left and lower right corner of the grid can be related to an ionically insulating AlPO4 impurity phase breaching the surface. Additionally the red circle in a) marks a particle, where further growth was partially impeded, when the insulating AlPO4-phase was reached at the lower left fringe of the particle. b) Topography profiles of nanoparticles grown at different temperatures between 180K and 280K. The red profile obtained for 240K corresponds to the profile line shown in the topography image of the 5x5 grid. c) Topography images of nanoparticles grown at different temperatures between 180K and 280K. at a elevated level, which drives the continuous compensation of normal forces during particle growth. The threshold voltages can be related to the reduction potential of the Li+ ions, which determine the minimum voltage, that must be applied to induce the reduction process. 28,35 Throughout our experiments, we found an average threshold voltage of about 3.8 ± 0.5 V, that matches well with the values found at room temperature 29,35 and shows no apparent temperature dependence (Fig. 2). For each temperature we have applied voltage ramps to the AFM-tip on 25 different positions in a 5 x 5 grid. After the particles had been grown, a control image without any bias voltage between tip and sample was taken to quantify the resulting particle sizes. A typical grid of Li-particles prepared at T = 240 K is shown in Fig. 3a. At first glance, it seems to be surprising, that we do not see 25 Li-particles, but instead only 21 particles are found. This discrepancy can be related to the particular struc-

ture of the LICGC, where the LATP-type ionconducting phase is interdispersed with small volumes of piezoelectric AlPO4 . 39 No particles can develop if the tip is in contact with the ionic-insulator AlPO4 . Consequently, such positions have been left out for the quantitative evaluation of ionic transport parameters. Additionally, particle growth related Li+ depletion effects in the sample as reported by Arruda et al. 35 have not been observed in this work, most likely due to the generally small sizes of our Li-particles. Therefore also the exact characteristics of our back-electrode do not influence the interpretation of the experimental results. We find that the size of our nanoparticles depends sensitively on the temperature, with lower temperatures leading to smaller particle sizes. This is illustrated in Fig. 3b and Fig. 3c, which show the topography of typical particles obtained at different temperatures and their cross-section, respectively. For a quantitative analysis the particle vol-


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particle growth process. This can best be done by analysing the linear sweep voltammograms of the Faradaic current recorded during particle growth as shown in Fig. 5. As discussed previously, it is assumed that the threshold voltage Uth for particle growth is determined by the standard reduction potential of the Li+ ions. Once this threshold is reached, the potential difference ∆U = U − Uth is divided between the charge transfer reaction at the sample/tip interface and the ion transport in the solid electrolyte. The reaction rate at the interface is typically described based on the Butler Volmer equation 40 with an exponential dependence of the current on the charge transfer overpotential. Therefore, only the initial stage of particle growth is reaction limited, i.e. if the overpotential is small. With increasing overpotential the growth process is quickly shifting from the reaction limited regime to the transport limited regime, 35 where ion migration towards the tip is the growth limiting factor. In this case, the migration related Faradaic current is determined by the applied overpotential 1 41,42 and the spreading resistance Rsp ∝ σion ·r with r the radius of the particle/substrate contact area and σion the ionic conductivity. In order to understand the resulting growth dynamics, let us first consider the case of a time-independent overpotential η, as it was previously analyzed in. 29 Under the assumption that the increase of the particle radius r with time t is described by a power law r ∝ tα , we expect the Faradaic current to increase as I ∝ η/Rsp ∝ tα . Consequently, the integral amount of reduced Li-ions and thus the particle volume V should scale as V ∝ t1+α . At the same time, the experiments in 29 and this work have revealed, that the Li-particles always grow in cone-like shape with the cone height h ∝ r. Based on this, the particle volume can be expressed as V ∝ r3 ∝ t3α . By comparing both presentations of the particle volume we can ultimately derive α = 0.5 (i.e. r(t) ∝ t0.5 ) as was indeed observed in the experiments of Ref. 29 As a next step, we analyze the particle growth for a linear increase of the overpotential η with time. Based on the previous assumption of r ∝ tα we can now derive V ∝ t2+α by integration of the

Figure 4: Arrhenius plot of the logarithm of the individual particle volumes vs. 1/kB T with kB the Boltzman constant and T the sample temperature (blue spheres). For each temperature the plot contains the volume of 18-23 particles and a linear fit to the experimental data yields an activation energy of 306 ± 6 meV. Additionally, the plot shows parameter A1 representing the degree of quadratic curvature d2 I/dU 2 for each temperature obtained by a quadratic fit to the voltammogram recorded during particle growth (see Fig. 5). From this curve an energy barrier of Ecurrent = 349 ± 50 meV is determined. (Please note, that the Faradaic current signals for T < 200 K were generally too low for quantitative analysis) umes were calculated. If the particles are sufficiently large, the sample can essentially be considered as flat. In this case, it is sufficient to use an automatic height threshold detection routine in the vicinity of each particle to assess its volume. By using this approach, our experiments revealed a decrease in volume by 3 orders of magnitude for a temperature reduction of 100K. The logarithmic values of all these volumes have then been plotted vs. the inverse temperature in the form of a semi-logarithmic Arrhenius plot (Fig. 4, blue spheres). A linear fit to the experimental data results in an effective activation energy barrier during particle growth of EP G = 306 ± 30 meV. For the interpretation of this activation energy, we have to consider the different electrochemical processes, which take place during the


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barrier for ion-conduction in the LATP-phase of the Ohara LICGC of Emacro = 330±3 meV. 39 This close match confirms our interpretation of the ion transport-limited current at higher overpotentials. In principle, both observables (i.e. particle size and quadratic curvature d2 I/dU 2 ) appear to be equivalent, meaning that within the margin of error closely matching energy barriers can be extracted. However, we found, that a reliable current measurement was already difficult or even impossible below 220K, where the already low Faradaic currents can be additionally obscured by electronic current contributions or background noise. At the same time the resulting particle sizes could still be easily quantified. This highlights one of the main advantages of Li-particle growth as a method to analyze nanoscopic ionic transport properties: By decoupling the particle growth and the subsequent evaluation of particle sizes, we can best benefit from the high spatial resolution of the atomic force microscope, which can be operated with optimum imaging parameters without the requirement of in-situ growth monitoring. This advantage is particularly important for imaging purposes. To demonstrate the imaging capabilities of low temperature Li particle growth, we have reduced the sample temperature even further down to 168 K, while still applying the previously described voltage ramps. In this case, typical particle volumes are as low as 300 nm3 . By considering the density of metallic Li we find that such particles only consist of approximately 15,000 Li-atoms, which is equivalent to an average reduction current of about 2 fA, that is well below our capabilities of current measurements. In principle, the growth of such small metallic particles may seem similar to previous works performed on superionic conductors 30 or small mixed-conductor nanodots. 37 E.g. within the latter work even atom-by-atom precipitation of metal nanoparticles originating from nonstoichiometric ions could be observed. In our work, we carry out measurements on a macroscopic sample of an ion conductor, so that atom-by-atom precipitation of Li-particles in not expected. Furthermore, instead of focusing

Figure 5: Exemplary linear sweep voltammograms recorded during Li-particle growth measured at three different temperatures and plotted versus the potential difference ∆U = U − Uth . At high cathodic overpotentials, all voltammograms can well be fitted by a function I(η) = I0 + A1 (η − η0 )2 , with three fit parameters I0 , A1 , and η0 . With all values of η0 in the range of η0 = 0.3 ± 0.1 V . Thus, η0 indicates the transition from the reaction limited regime (region I) to the transport limited regime (region II). (Please note, that the low currents do not allow for an in-depth analysis of region I) Faradaic current. Comparing this result again to V ∝ r3 ∝ t3α we can determine α = 1, which ultimately results in a current vs. voltage characteristic described by I ∝ η 2 . To check if this relation can indeed describe the experimentally observed current vs. voltage characteristics, we have analyzed the Faradaic current as a function of the overpotential η = U − Uth for different temperatures (see Fig. 5). In all cases, the voltammogram can be fitted well with a square dependence on the overpotential η, if a small voltage offset η0 of about 0.3 V related to the reaction limited regime is considered. We find that the prefactor A1 = d2 I/dη 2 of the quadratic increase of the current also shows a strong temperature dependence. From a corresponding Arrhenius-plot (see Fig. 4, red triangles) an activation energy of Ecurrent = 349 ± 50 meV can be extracted. The values of EP G and Ecurrent both agree well with macroscopically determined energy


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Figure 6: Mapping of ion conductivity by analysing the particle height promptly before and after the particle formation. Bias was applied in a 80x80 spectroscopy grid at 168 K. (a) Topography image of the probed sample area before the spectroscopy scan. (b) Topography image of the probed sample area after the spectroscopy scan. (c) Inset of the topography image. (d) Descriptive image of particle height calculation routine. In the retrace movement of the AFM tip between scan line n − 1 and n the bias is applied every 25 nm. The mean height of the particle represented by red dots is subtracted from the mean height of the topography before particle formation (shown by green dots). (e) Difference value of the particle height without topography crosstalk. on nucleation processes at relatively low overpotentials (typically 50-100 mV), we focus on ion transport limitations at high overpotentials up to 1 V. Fig. 6a shows a topography image of 2µm × 2µm before particle growth. After this reference scan Li-particles have been deposited in an 80 × 80 grid within this area. The resulting topography after particle growth is shown in Fig. 6b. Within this image we can clearly resolve single Li-particles (see Fig. 6c), which means that our principle resolution is well below the 25 nm, that characterizes the grid peri-

odicity and is at the same time similar to the tip size. However, for the example shown in Fig. 6, separating the Li particle sizes from the topography is not straightforward. For images like shown in Fig. 3a, determining the Li-particle sizes can easily be done based on the previously described automatic threshold detection approach. This approach is not feasible for high resolution images like shown in Fig. 6, where high particle density and surface roughness prevent accurate threshold detection. At the same time also lateral drift effects can become relevant due to the small particle sizes and the long


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time (t . 2 h) required to grow all particles of the 80 × 80 grid. Therefore, also an alternative approach based on subtracting images recorded before and after particle growth is not accurate enough for imaging applications. Thus, we ultimately applied a strategy, which is outlined in Fig. 6d and relies on the fact, that particle growth is conducted during the retrace of every 5th scan-line during an otherwise conventional scan of 2µm × 2µm. If e.g. particles are grown during the retrace of line n − 1, then the trace lines n − 1 and n are the ones immediately before and after particle growth. While the trace of line n−1 still yields the undisturbed topography, line n and n + 1 are directed almost perfectly through the center of the newly grown Li-particles. Consequently, the height of the particle can be calculated by subtraction of the corresponding averaged z signals measured during lines n − 1 to n +1 (Please see Fig. 6c for details on the exact procedure). In this case, drift effects become negligible, since subsequent scan lines are used to extract the particle height. With this approach we can quantify the height of each particle of the 80 × 80 grid and use this data to derive a spatially resolved representation of the particle sizes (Fig. 6e). From Fig. 6d we can see that the resulting particle height is very small with maximum values around 2 nm. Additionally, we find that the particle growth is very inhomogenous. Most prominently, we can recognize large patches where no Li-particles have grown. These areas can be identified as ionically insulating AlPO4 phases, the occurrence of which is also evident from the topography images. However, also the particle sizes grown on the ion conducting phase are varying considerably, but in this case, a direct correlation with the topography images is difficult. Instead particle sizes might strongly be influenced by the subsurface characteristics of the sample in correlation with insulating AlPO4 impurity phase. An indication of subsurface effects might be derived from Fig. 6d, where especially the areas around the AlPO4 -phases are interesting. In many cases, we can observe only relatively small particle sizes close to the edges of the AlPO4 -phases surface areas, while only occasionally also larger

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particle sizes can be found. A possible explanation for this behaviour lies in the subsurface extension of the AlPO4 -phases. When the tip is close to the grain boundary of an ion conducting phase and the insulating AlPO4 phase, the ionic current can only flow across a fraction of the subsurface volume determining the spreading resistance, namely the fraction containing the ion conducting phase. This leads to a reduction of the ionic current and of the particle volume

CONCLUSIONS To conclude our measurement show that temperature is an excellent control parameter for AFM-induced growth of nanometer sized Liparticles. Systematic variation of the sample temperature allows extracting the characteristic energy barrier EP G from the temperature dependent size variation of the Li-particles. These energy barriers are comparable to macroscopic values. 39 They have simulateously been derived from the slope dI/dU curves of the corresponding voltammogramms during growth, with similar results. Thus we can conclude that EP G reflects the predominant influence of ion migration on the particle growth process and characterizes the mobility of Li-ions within the LICGC. Secondly, we have demonstrated how low temperature particle growth can be applied for high resolution imaging of ionic transport parameters. By lowering the sample temperature down to 168 K, typical faradaic currents dropped to the fA range, which is well below our sensitivity for current measurement. But still, the resulting particles could well be detected by the AFM. This resulted in a lateral resolution of about 25 nm, which is already close to the resolution limit imposed by the finite diameters of the AFM-tip. We found that the AlPO4 phases play an important role for the resulting particle sizes. If the AlPO4 -phases are directly at the surface no particles are formed due to the ion insulating properties of AlPO4 . At the same time, these AlPO4 -phases also seem to affect particle sizes in their immediate vicinity.


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Often, reduced particle sizes are found close to the AlPO4 -phases, an effect that might be related to their potential influence on subsurface ion diffusion by subsurface constrictions of the current paths. In further experiments, the scope of this experimental technique can now be broadened, either by analysing further solid state electrolyte materials or e.g. by spatially resolved characterisation of ion conducting thin films on electrode materials (e.g. SEI-films or artifical layers like LiNbO3 43 ). Especially the simultaneous combination of nanoscale voltammetry and particle growth is intriguing. While the first allows for a detailed analysis of localized Faradaic currents, the latter offers ultimate sensitivity spatial resolution by ex-situ measurement of integral currents (i.e. particle sizes) at low temperatures.

erally around 10 nN. The comparatively low normal loads and the generally good durability of the diamond cantilevers helped to minimize any effects of tip changes throughout the course of our measurements. Cooling was facilitated using liquid nitrogen within a flow cryostat, while an active temperature controller employed additional Joule heating to maintain a constant sample temperature. After each change of temperature, we waited for at least 30 minutes for thermal equilibrium and a stationary sample temperature. The bias ramps applied to generate particle growth were generated by the internal waveform generator of the AFM control electronics. During particle growth processes the AFM feedback always remained active, thereby preventing normal load changes within the electrical contact between the cantilever-tip and the sample. Acknowledgement Financial support was provided by the German Research Foundation (Project DI917/7-1), and in part by the Laboratory for Material Science of the Justus Liebig University Giessen. We acknowledge financial support within the LOEWE program of excellence of the Federal State of Hessen (project initiative STORE-E).

MATERIALS AND METHODS All measurements presented in this work have been performed on a commercial Lithium-Ion Conducting Glass-Ceramic (LICGC - AG01), that was purchased from Ohara Inc., Japan. The Ohara LICGC has a high Li-ion conductivity of appr. 10−4 S/cm at room temperature, 44 while at the same time the electrical conductivity is extremely low. The Ohara LICGC is stable in air, water, and aqueous solutions 45,46 and could thus be used in model systems demonstrating Li-Air and Li-water batteries. 47 The sample had a thickness of 150 µm and was fixed to the sample holder using a conductive glue. Throughout the experiments, the sample holder and thus also the LICGC were permanently grounded, while varying voltages were applied to the conductive AFM tip (see Fig. 1). The AFM cantilevers (Adama Innovations, AD-E-0.5-AS) had a nominal force constant of k = 0.5 N/m and all measurements shown in this work were performed using an Omicron VT-AFM operated in contact mode at ultra high vacuum conditions with base pressures of about 3 × 10−10 mbar. Normal forces during AFM measurements were gen-

Notes: The authors declare no competing financial interest.

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Graphical TOC Entry

AFM assisted Li-particle growth on Li-ion conductive glass ceramics by applying sufficient voltages between tip and sample to reduce the Li-ions. The size of the resulting metallic Li- particles sensitively depends on the temperature and allows to quantify characteristic ion transport barriers.


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Nanoscale characterization of ion mobility by temperature controlled Li

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