Novel ALD Chemistry Enabled Low-Temperature Synthesis of Lithium

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Novel ALD Chemistry Enabled Low-Temperature Synthesis of Lithium Fluoride Coatings for Durable Lithium Anodes Lin Chen, Kan-Sheng Chen, Xinjie Chen, Giovanni Ramierez, Zhennan Huang, Natalie Geise, Hans-Georg Steinrück, Brandon L. Fisher, Reza Shahbazian-Yassar, Michael F. Toney, Mark C Hersam, and Jeffrey W. Elam ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b04573 • Publication Date (Web): 09 Jul 2018 Downloaded from http://pubs.acs.org on July 11, 2018

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ACS Applied Materials & Interfaces

Novel ALD Chemistry Enabled Low-Temperature Synthesis of Lithium Fluoride Coatings for Durable Lithium Anodes Lin Chen1,2, Kan-Sheng Chen3, Xinjie Chen3, Giovanni Ramirez1, Zhennan Huang4, Natalie R. Geise6,7, Hans-Georg Steinrück6, Brandon L. Fisher5, Reza Shahbazian-Yassar4, Michael F. Toney6, Mark C. Hersam3, and Jeffrey W. Elam*1,2 1

Energy System Division, Argonne National Laboratory, Lemont, Illinois 60439, USA

2

Joint Center for Energy Storage Research, Argonne National Laboratory, Lemont, Illinois 60439,

USA 3

Department of Materials Science and Engineering, Northwestern University, Evanston, Illinois

60208, USA 4

Department of Mechanical and Industrial Engineering, University of Illinois at Chicago, Chicago,

Illinois 60607, USA 5

Nanoscience & Technology Division, Argonne National Laboratory, Lemont, Illinois 60439,

USA 6

Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Center, Menlo Park,

CA 94025, USA 7

Department of Chemistry, Stanford University, Stanford, California 94305, USA

*

Corresponding Author, [email protected]

Keywords: Lithium Metal Anode, Atomic Layer Deposition, New Chemistry, Lithium Fluoride, High Shear Modulus

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Abstract: Lithium metal anodes can largely enhance the energy density of rechargeable batteries due to the high theoretical capacity and the high negative potential. However, the problem of lithium dendrite formation and low Coulombic efficiency (CE) during electrochemical cycling must be solved before lithium anodes can be widely deployed. Herein, a new atomic layer deposition (ALD) chemistry to realize the low-temperature synthesis of homogenous and stoichiometric lithium fluoride (LiF) is reported, which then for the first time, as far as we know, is deposited directly on lithium metal. The LiF preparation is performed at 150°C yielding 0.8 Å/cycle. The LiF films are found to be crystalline, highly conformal, and stoichiometric with purity levels >99%. Nanoindentation measurements demonstrate the LiF achieving a shear modulus of 58 GPa, seven times higher than the sufficient value to resist lithium dendrites. When used as the protective coating on lithium, it enables a stable Coulombic efficiency as high as 99.5% for over 170 cycles, about 4 times longer than that of bare lithium anodes. The remarkable battery performance is attributed to the nanosized LiF that serves two critical functions simultaneously: 1) high dielectric value to create a uniform current distribution for excellent lithium stripping/plating and ultrahigh mechanical strength to suppress lithium dendrites; 2) superior stability and electrolyte isolation by the pure LiF on lithium prevents parasitic reactions for much improved CE. This new ALD chemistry for conformal LiF not only offers a promising avenue to implement lithium metal anodes for high capacity batteries, but paves the way for future studies to investigate failure and evolution mechanisms of solid electrolyte interphase (SEI) using our LiF on anodes such as graphite, silicon and lithium. 1. Introduction Societal demands for energy storage in portable electronics, electric vehicles, and renewable energy are driving the development of advanced battery technologies1-3. For instance, the modest charge storage capacity of the graphite anodes used in commercial lithium ion batteries (372 mAh/g) has motivated research to discover a higher capacity alternative anode material4-7. Lithium metal is an ideal anode for rechargeable lithium-based batteries due to its high theoretical capacity (3,860 mAh/g) and low redox potential (-3.04 V vs. standard hydrogen electrode)5, 8-9. Lithium metal anodes also facilitate high-energy cathodes (e.g., sulfur or air) for next-generation systems10. However, the large-scale deployment of lithium metal anodes in rechargeable batteries has been hindered by severe technical hurdles, including large volumetric changes and lithium 2


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dendrite formation during cycling11-12, resulting in low Coulombic efficiency and short cycling life13-14. Many of the problems plaguing lithium metal anodes result from deficiencies in the solid electrolyte interphase (SEI) that can form spontaneously when lithium metal contacts the organic electrolyte. For instance, this spontaneous SEI is spatially nonuniform, and exhibits variations in the local rate of lithium ion transport, which results in differential volume changes during cycling, cracking of the relatively fragile SEI, and exposure of fresh lithium. Consequently, new SEI forms and the cell eventually fails when the electrolyte is consumed. The nonuniform lithium ion flux through the SEI can also nucleate lithium dendrites causing an electrical short circuit failure15. In contrast to the natural SEI that forms on lithium, an artificial SEI can be fabricated with great uniformity and desired material properties9, 16-19. An ideal interface would possess the following attributes: 1) chemically stable in a highly reducing environment11; 2) uniform, self-limiting thickness and composition9,20; 3) excellent adhesion on lithium to resist delaminating; 4) electronically insulating9 to ensure lithium plating on the underlying metal; 5) mechanically robust to resist cracking and prevent breaching by lithium dendrites21-22; 6) high lithium ionic conductivity to facilitate uniform lithium stripping and plating. In previous work, numerous strategies have been explored to improve the lithiumelectrolyte interface and achieve dendrite-free lithium anodes, such as highly concentrated electrolytes23-24, engineered polymers14, 25-26, electrolyte additives4, 15, 27 and physical barriers28-29 on lithium metal. Among these approaches, physical barriers have the advantage that they add minimal weight and volume to the cell if they can be made sufficiently thin while maintaining the properties listed above. Recently, atomic layer deposition (ALD) has been evaluated as a technique for depositing ultrathin, conformal coatings on lithium metal.9, 30-32 Atomic layer deposition uses alternating, self-limiting chemical reactions between precursor vapors and a solid surface to grow materials in an atomic layer-by-layer fashion. Saturation of the individual ALD surface reactions provides sub-nm thickness control and ensures uniform composition and thickness across the substrate33. A broad variety of materials can be prepared by ALD so that the properties can be tailored to meet the requirements of a robust “artificial” SEI, and many ALD processes can be performed at temperatures below the lithium melting point (181°C). Previous studies of ALD coatings on lithium metal have examined Al2O3, which is soluble in acids and bases, and lithium aluminum sulfide, which is highly reactive with oxygen and water vapor. Degradation of these 3



coating materials during cycling can compromise their ability to protect the Li surface leading to reduced cycle life. In contrast, lithium fluoride is a common SEI component34 and is chemically inert on lithium metal. LiF is electrochemically stable from 0-6.4 volt35, is electronically insulating, and has a relatively high shear modulus of 55 GPa36. It has been reported that materials having a shear modulus at least two times higher than metallic Li (4 GPa) should suppress dendrite growth37. Previously, lithium fluoride has been prepared by ALD using lithium 2,2,6,6-tetramethyl3,5-heptanedionate (Li-thd) and titanium tetrafluoride (TiF4)38, magnesium bis(2,2,6,6tetramethyl-3,5-heptanedionate) (Mg(thd)2), Li-thd and TiF439, and lithium tert-butoxide (LiOtBu) and TiF440. However, all of these approaches require temperatures above the Li melting point, and can contaminate the film with Mg and Ti39. Lin et al. reacted lithium foil with Freon to form a LiF coating, but the thickness was difficult to control since the growth requires bare Li to react with the Freon, and the films contained carbon41. In this study, we present an alternative method for LiF ALD using LiOtBu and HF-pyridine solution (Figure 1a). The HF-pyridine solution comprises 70% HF and 30% pyridine, and is a liquid at room temperature with a high vapor pressure of ~100 Torr that is essentially pure HF42. We first study the LiF ALD using in situ quartz crystal microbalance (QCM) measurements to gain insight into the ALD chemistry and to verify self-limiting growth. Next, we deposit the LiF films on silicon coupons and characterize the films using spectroscopic ellipsometry (SE), X-ray photoelectron spectroscopy (XPS), atomic force microscopy (AFM), grazing incidence wide-angle X-ray scattering (GIWAXS), Rutherford backscattering spectrometry (RBS), hydrogen forward scattering (HFS), particle-induced gamma ray emission (PIGE), and nanoindentation. The ionic conductivity of the LiF was determined using impedance spectroscopy of films deposited on interdigitated platinum electrode substrates. Finally, we deposit LiF films on lithium metal anodes and fabricate symmetric (Li-Li) and asymmetric (Li-Cu) coin cells in order to assess electrochemical properties and cycling performance. 2. Experimental section 2.1 Lithium Metal Protection Commercial lithium was punched into disks with a diameter of 7/16 inches in an Ar-filled glovebox (O2 and moisture level lower than 0.5 ppm) and pressed to a flat surface. ALD was conducted in a custom, viscous flow reactor with the reaction zone consisting of a heated, 5 cm diameter stainless steel tube. The pressure within the reactor was maintained at ∼1 Torr using a 4


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216 sccm flow of ultrahigh purity (UHP, 99.999%) argon carrier gas. The reactor was attached to a glovebag to avoid exposing the Li foils to air. LiF ALD was conducted using LiOtBu (Aldrich, 97%) and hydrogen fluoride (HF)-pyridine (Aldrich, 70% HF 30% pyridine). The LiOtBu was contained in a stainless steel bubbler heated to 135°C and delivered using 36-sccm UHP argon, while the HF-pyridine was contained in a stainless steel cylinder maintained at room temperature and delivered by its own vapor pressure. The Li foils were transported from the glovebox to the ALD reactor glovebag using Ar-filled mason jars to prevent air exposure. The LiF ALD timing sequences can be expressed as: t1−t2−t3−t4, where t1 is the LiOtBu exposure time, t3 is the HF exposure time, and t2 and t4 are the respective purge times for LiOtBu and HF, with all times in seconds. The timing sequence used here for LiF on Li is 5-5-1-5. After ALD coating, the samples were transferred back into a sealed mason jar and immediately transported to the glovebox. 2.2 Characterization To investigate the growth of ALD LiF, a quartz crystal microbalance (QCM) sensor housed in a holder was loaded into the reactor. The crystal was maintained at the 150°C ALD reactor temperature for several hours to achieve thermal equilibrium. The mass changes per unit area were calculated by measuring the frequency shifts and applying the Sauerbrey equation. To prepare TEM samples, 100 nm SiO2 nanoparticles were dispersed onto holy carbon grids and loaded into the reactor for LiF coating. The structure was analyzed using a spherical aberration corrected JEOL JEM-ARM200CF scanning transmission electron microscope (STEM) equipped with a 200 kV cold-field emission gun, annular bright field (ABF) and high angle annular dark field (HAADF) detectors, and an energy dispersive spectrometer (EDS) using windowless EDS 100mm Oxford detector. AFM characterization was performed using ~300 kHz Si cantilevers in normal tapping mode on an Asylum Cypher AFM. Scanning electron microscopy (SEM) was performed using a JEOL 7000 unit. For nanoindentation, the hardness (H) and elastic moduli (E) of silicon and LiF thin film were measured. These measurements were performed using a Hysitron TI 900 nanoindenter, equipped with a three-sided diamond Berkovich tip with a radius of curvature of ∼150 nm and 65.35° half-angle. Grazing incidence wide-angle X-ray scattering (GIWAXS) measurements were obtained at the Stanford Synchrotron Radiation Lightsource at beamline 11-3 using an incident beam energy 5



of 12.735 keV. A Raxyonics 225 detector was used with a sample to detector distance of 153 mm. An incident angle of 0.16° was selected, above the critical angle for the silicon substrate (0.14°). The sample was enclosed in a helium chamber during measurements. GIWAXS spectra were processed using the Nika software package for Wavemetrics Igor, in combination with WAXStools, an add-on script for Igor. For the ionic conductivity measurement, 50 nm LiF was deposited onto the active side of a comb chip (DropSens, Interdigitated platinum Electrode, IDE). These interdigitated Pt electrodes provide a 6.7x105 enhancement in ionic current compared to a simple square geometry. In-plane ionic conductivity measurements were conducted on the 50 nm LiF coated Pt electrode in an Arfilled glovebox using electrochemical impedance spectroscopy (EIS). Microprobes were used to contact the electrode, which was placed on a hot plate to enable the measurements in the temperature range of 100-425 ºC. EIS for the ionic conductivity was performed with a Solartron 1260 impedance/gain phase analyzer combined with a Solartron 1287 electrochemical interface. The frequency is from 1 MHz to 0.1 Hz with an amplitude of 5 mV. 2.3 Electrochemical Measurements Bare Li foils and LiF-coated Li foils with a diameter of 7/16 inch were used. 2032-type coin cells were fabricated in an Ar-filled glovebox for symmetric and asymmetric (Li-Cu) cell measurements. For the symmetric cells, Li was employed on both sides as the working and counter electrodes, and carbonate electrolyte with 1 M LiPF6 in EC/EMC (3:7 in volume) from BASF was used. Different current density and capacity were used for testing the Li symmetric cells, including 0.4 mA/cm2 to a capacity of 1 mAh/cm2 and 1 mA/cm2 to a capacity of 1 mAh/cm2. For the Li-Cu tests, Cu was applied as the working electrode and Li functioned as the counter electrode, where 1 M lithium bis(trifluoromethylsulphonyl)imide (LiTFSI) with 0.18 M Li2S8 and 2 wt.% LiNO3 in DOE/DOL solvent (1:1, volume ratio) was used. Li-Cu cells were initially discharged at a current density of 1 mA/cm2 to 1 mAh/cm2, followed by stripping from Cu to 1 V at 1 mA/cm2. Separators used in this work were Celgard 2325, and the cycling measurements were performed using an Arbin 2043 and a LAND instrument. Electrochemical impedance spectroscopy (EIS) of Li/Li symmetric cells before cycling was measured by a Solartron 1260 impedance/gain phase analyzer combined with a Solartron 1287 electrochemical interface. The frequency was swept from 1 MHz to 0.1 Hz with an amplitude of 5 mV. 6


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3. Results and Discussions 3.1 Growth and Characterization of ALD LiF Films We performed in situ quartz crystal microbalance (QCM) measurements to establish the LiF ALD and to examine the growth mechanism. Figure 1a shows the in situ QCM measurements during alternating lithium tert-butoxide (LiOtBu or LTB) and HF-pyridine exposures at 150°C using the timing sequence: 5-5-0.2-5. Figure 1a reveals linear growth at ~23 ng/cm2. Assuming a bulk density for LiF of 2.64 g/cm3, this corresponds to a growth per cycle (GPC) of 0.88±0.03 Å/cycle where the error range is taken as the standard deviation of the growth per cycle values for the individual steps in Figure 1b. The inset to Figure 1b shows a detailed view of the QCM measurements during three successive LiF ALD cycles where the black trace shows the QCM data and the blue and red traces show the status of the LTB and HF-pyridine dosing valves, respectively. The mass increases by ~59 ng/cm2 during each LTB exposure and then decreases by ~35 ng/cm2 during each HF-pyridine exposure so that the net mass change is ~23 ng/cm2. We propose the following half-reactions for the LiF ALD: -(H)x* + LiOtBu → Li(OtBu)1-x* + x HOtBu

(1)

Li(OtBu)1-x* + HF → LiF-(H)x* + (1-x) HOtBu

(2)

where the asterisks designate surface species. In Eq. 1, LTB reacts with x surface hydrogen species to release x tert-butanol molecules (HOtBu) and leaving 1-x tert-butoxide ligands on the surface. In Eq. 2, the tert-butoxide terminated surface reacts with HF releasing the remaining 1-x ligands as HOtBu to form LiF and repopulate the surface with hydrogen. These equations assume that pure, stoichiometric LiF is formed, surface H (in the form of adsorbed HF or Li-H species) are the reactive functional groups for the LTB molecule, and HOtBu is the only gas-phase reaction product. The second two assumptions would require additional in situ measurements to verify, but the assumption of pure, stoichiometric LiF will be addressed below. The structure of the LiF QCM signals can be used to extract the value for x, the number of surface H reacting with each LTB molecule. We can define R=A/B where A is the mass change generated by the LTB exposure and B is the total mass change of one ALD cycle. From Figure 1b, R=2.5±0.1 where the error bars are the standard deviation from the individual QCM step values. Using Eqs. 1 and 2, and the atomic weights of the species, R=(79-73x)/26 so that x=0.18±0.03 implying that, on average, only 18% of the OtBu ligands are removed upon LTB exposure and the remaining 82% are released during 7



the subsequent HF exposure. The large fraction of OtBu ligands released during the HF exposure accounts for the large mass decrease observed in the QCM measurements. Next, we prepared a series of LiF films on silicon coupons using 300-1500 ALD cycles at 150°C with the timing sequence 5-5-1-5, and measured the thickness using spectroscopic ellipsometry (SE, Figure 1c). The growth is linear with a GPC of 0.82±0.02 Å/cycle. This value is within the measurement error of the value 0.88±0.03 Å/cycle derived from QCM, indicating that the true density is close to the assumed bulk density of 2.64 g/cm3, consistent with stoichiometric LiF (vide infra). In support of this finding, the SE data were fit with a Cauchy model and yielded a refractive index of 1.392 at 633 nm, in agreement with the literature value of 1.391. To evaluate the conformality and morphology of the LiF, we prepared LiF films with a thickness of ~15 nm LiF on ~100 nm silica nanoparticles dispersed on a carbon transmission electron microscopy (TEM) grid. The TEM image in Figure 1d reveals that the silica nanoparticles are coated with a thin shell of ~15 nm having a granular morphology (~10-50 nm grains). This same morphology is evident on the underlying carbon TEM grid indicating that the LiF has also coated the carbon. Energy dispersive X-ray analysis (EDX) mapping performed in the TEM showed that the particles consist of Si and O, and that the coating contains F (Figure S1). Selected area electron diffraction (SAED) was performed on this specimen and showed four rings consistent with the (111), (200), (220), and (311) planes of lithium fluoride (Figure 1e), thus indicating that the ALD LiF is crystalline. Finally, we prepared a 60 nm LiF film on Si, and performed atomic force microscopy (AFM). Figure 1f presents a 1-µm × 1 µm AFM image and shows that the film consists of 20-100 nm crystals, some of which resemble cubes with vertices aligned normal to the surface. The root mean squared (RMS) roughness is ~6 nm. The morphology of our films is similar to previous ALD LiF films prepared using Mg(thd)2, Li-thd and TiF4 , although our roughness is 2-3 times lower, probably due to our lower growth temperature of 150°C compared to the previous study (275-350°C)39. Additional AFM images are provided in Figure S2.

8


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(a)

(b)

(c)

(d)

(e)

(f)

Figure 1. (a) Synthesis schematic for ALD LiF using LiOtBu and HF-pyridine. (b) In situ QCM measurements during LiF ALD. (c) Ellipsometric thickness measurements of LiF films deposited on silicon substrates. (d) TEM image of SiO2 nanospheres dispersed on carbon TEM grid and subsequently coated using 15 nm LiF. (e) Selected area electron diffraction of 15 nm LiF coating. (f) AFM of 60 nm LiF on silicon.

In order to confirm the crystallinity of the LiF, 60 nm films were prepared on silicon and analyzed using X-ray diffraction (XRD). The XRD data shown in Figure S3 reveal two peaks at 38.8º and 45.1º, corresponding to (111) and (200) LiF reflections, respectively43, whereas the unmarked peaks are from the silicon substrate. Due to the low signals in the XRD data (Figure S3), we performed grazing incidence wide-angle X-ray diffraction (GIWAXS) measurements on a 118 nm LiF film on silicon. Figure 2a shows the GIWAXS intensity versus q(Å-1), and reveals two broad peaks at 2.78 and 3.21 Å-1 corresponding to (111) and (200) LiF reflections, correlating with the brightest rings in the SAED image (Figure 1e). The lattice constant, calculated from positions 9



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of the (111), (200), and (220) reflections, is 3.95±0.03 Å indicating that the LiF is compressively strained compared to the 4.03 Å lattice parameter reported in the literature.44 This strain may affect Li ion transport through the LiF film, although origin of the strain is unclear, (possibly from point defects or a film confinement effect). The diffraction pattern also shows that the films have a preferential (111) and (220) crystallographic texture, (Figure S8) consistent with the vertically aligned cubes observed by AFM (Figure 1f). Rutherford backscattering spectrometry (RBS) coupled with hydrogen forward scattering spectrometry (HFS) and nuclear reaction analysis (NRA) was performed on the LiF film on silicon to detect the ratios of any elements including lithium, carbon, and hydrogen. Figure 2b shows the elemental concentration versus depth from these measurements. We found hydrogen at a concentration of ~6% in the top 20 nm of the lithium fluoride film accompanied by an equivalent deficiency in F in this region. Below 20 nm, the film is essentially pure, stoichiometric LiF. We next performed X-ray photoelectron spectroscopy (XPS) measurements of a LiF film on silicon. Analysis of the XPS survey spectrum recorded after performing Ar sputtering in Figure 2c yields a Li:F ratio of 1:1 and only trace oxygen (0.7%) and C (2.5%). These contaminants were decreased by a factor of 2 by brief Ar sputtering, so we attribute these peaks to atmospheric contamination. Figure 2d and 2e show high-resolution scans of the Li 1s region with a peak centered at 55.6 eV, and the F 1s region with a peak centered at 685.1 eV, respectively. These peak positions match well with the literature spectrum for bulk LiF45. (b)

(a)

(c)

(d)

(e)

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Figure 2. (a) GIWAXS scan of a 118 nm LiF film on silicon showing (111) and (200) peaks of cubic LiF phase. A sharp Si(111) peak has been masked for clarity (see Figure S8). (b) Elemental concentration profiles for 116 nm LiF film on silicon modeled from RBS/HFS/NRA measurements. (c) X-ray photoelectron spectroscopy (XPS) survey scan from a 103 nm LiF film on silicon, and high resolution scans of peaks (d) Li 1s peak indicating LiF and (e) F 1s peak indicating LiF.

We further performed nanoindentation measurements on LiF thin films prepared by ALD in order to evaluate mechanical properties. Lithium metal has a shear modulus of 4 GPa46. Newman et al. predicted that a coating with a shear modulus two times that of lithium can effectively suppress dendrite formations21. Similarly, a recent simulation concluded that a shear modulus > 8.5 GPa is sufficient to suppress lithium dendrites22. To ensure the accuracy of our measurements, we first measured force-displacement curves for bare silicon as a function of penetration depth (Figure S4a). Using the Oliver-Pharr method (see SI Eqns. S1 and S2)47, we calculated a hardness of 11 GPa (Figure S4b) and an elastic modulus of 174 GPa (Figure S4c) similar to literature precedent48. Figure 3a shows the force-displacement curves measured for LiF on Si, from which the reduced modulus is 148 GPa (Figure 3b) and the shear modulus is 58 GPa (Table S1). This value agrees well with the previously reported value of 55.1 GPa36. Consequently, the ALD LiF has a shear modulus 6-7 times higher than the value required to suppress lithium dendrites. We note that the mechanical properties of the ALD LiF may be influence by the liquid electrolyte in a working cell.

This effect could be investigated using nanoindentation

measurements in the presence of the liquid electrolyte49 (b)

(a)

Figure 3. (a) Force-distance curves for a 122 nm LiF thin film on silicon measured using nanoindentation. (b) Reduced modulus values for lithium fluoride calculated from Figure 3a using Eqn. S1.

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`

3.2 Battery Performance of the Pure ALD LiF Coated Li We prepared LiF coatings of 8 nm thickness on lithium or 6 nm thickness on copper foils in order to evaluate the barrier properties of these coatings for battery applications. Symmetric cells having the same electrodes as both cathode and anode were fabricated using both bare Li and LiF-coated Li foils. A constant current was applied for a fixed time to achieve the desired charging and discharging capacity. Voltage profiles were recorded to determine overpotential during the cycling process. Figure 4a shows the results of a Li ∣ Li cell (black) and a LiF-coated ∣ LiF-coated Li cell (red), which were assembled with 20-μL electrolyte and tested at a current density of 1 mA/cm2 up to a capacity of 1 mAh/cm2. As seen in Figure 4a, the bare Li symmetric cell shows a higher overpotential compared to the LiF-coated symmetric cell over the entire measurement period. At ~170 hours, the bare Li cell exhibits a sudden voltage drop and subsequent voltage fluctuations consistent with a short circuit. In contrast, the voltage of the LiF coated cell remains very stable until 260 hours when the experiment was stopped. (b)

(a)

(d)

(c)

Figure 4. (a) Voltage versus time during electrochemical cycling at a current density of 1 mA/cm2 up to a capacity of 1 mAh/cm2 for symmetric cells using 20 μL carbonate electrolyte with pristine Li (black) and Li coated with 8 nm LiF (red). (b) Duplicate test to Figure 3a but using 5-μL carbonate electrolyte. (c) Nyquist plots of a 50 nm LiF film on Pt patterned silica versus substrate temperature. (d) Arrhenius plot for ionic conductivity of the ALD LiF.

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A recent report described the importance of low electrolyte volumes in Li-S batteries incorporating lithium metal anodes to achieve a sufficiently high volumetric capacity for commercialization50. Electrolyte volume is a critical metric for batteries incorporating lithium metal anodes, which can consume electrolyte continuously during operation. Unfortunately, most lithium battery studies do not report the effects of electrolyte volume on performance. To study the effect of electrolyte volume on lithium anodes coated with LiF, we prepared symmetric cells using only 5-μL electrolyte while keeping all other parameters the same as in Figure 4a. As shown in Figure 4b, failure occurs after 80 hours for the bare Li cell with 5-μL electrolyte, compared to ~170 hours for the cell with 20-μL electrolyte. In comparison, the symmetric cell with LiF coated Li survives for ~180 hours, twice as long as the bare Li cell. Figure S5 shows photographs of separators harvested from the Li symmetric cells incorporating 5-μL electrolyte with and without the LiF protective coatings after prolonged cycling. Figure S5a shows one side of the separator from the bare Li cell after 100 hours of cycling, revealing a thick layer of “dead” lithium and SEI, as well as multiple black spots indicative of short circuits. The reverse side of this separator shows the same characteristic features (Figure S5b). In contrast, the separator extracted after 200 hours of cycling from the LiF-coated symmetric cell is much cleaner (Figs. S5c-d) with minimal discoloration or debris on this separator associated with dead lithium and SEI. We attribute the more stable voltage profile (Figs. 4a-b) and the reduced dead lithium and SEI (Figs. S5c-d) to chemical passivation of the Li surface provided by the LiF film. As discussed in the introduction, an ideal passivation coating on lithium metal should have a high ionic conductivity to ensure a uniform ion flux that promotes planar electrodeposition, while also avoiding lithium deposition on the top surface of the coating. Figs. 4c and 4d display the results from electrochemical impedance spectroscopy (EIS) performed on a LiF coating versus measurement temperature. The Nyquist plots (Figure 4c) show one semicircle followed by a tail. These data were fit with an equivalent circuit consisting a resistor and constant phase element (CPE) in parallel, which implies that the ionic resistance is given by the intersection of the linear tail on the x-axis. Figure 4d shows an Arrhenius plot for the calculated ionic conductivities, and the linear fit yields an activation energy of 0.93±0.02 eV and an extrapolated room temperature ionic conductivity of 10-14(±0.5) S/cm. Experimental values for the activation energy of Li ion diffusion in LiF range from 0.65-1.1 eV51-53, and the extrapolated room temperature ionic 13



conductivity has been shown to be 10-12 S/cm53. A recent density functional theory (DFT) study examined a broad range of possible conductivity mechanisms in LiF and calculated activation energies between 0.27 and 1.89 eV, with an average value of 0.89 eV, and an average room temperature ionic conductivity of 10-13 S/cm54. Our measured activation energy and room temperature ionic conductivity are comparable to previous measured and predicted values, all of which indicate that the ionic conductivity of LiF is low. In fact, our measured ionic conductivity predicts an impedance of ~108 Ω for our coin cells rather than the measured value of ~100 Ω (estimated from the voltage of 100 mV with a current of 1 mA from Figure 4a). This discrepancy probably results from the calculations and measurements being performed in a “dry” state where the LiF is either isolated or in contact with relatively inert electrodes such as platinum52-53. In the coin cells, the LiF contacts both lithium metal and liquid electrolyte, and these interactions could dramatically enhance the ionic conductivity. For instance, the low electrochemical potential of Li (-3.05 V vs. standard hydrogen electrode) can reduce the LiF generating Li interstitial defects or a new, Li-rich LiF compound. Similarly, fluorine from the LiF can diffuse into the Li creating F vacancies. These defects can increase the charge carrier concentration and/or new compounds formed in the Li-LiF reaction can increase the mobility, with both effects increasing the lithium ion conductivity. Similar arguments can be made regarding interactions between the LiF and the liquid electrolyte. Coulombic efficiency (CE) is another important parameter for evaluating lithium metal performance and cycle life. Lithium-copper cells were employed to measure the efficiency of lithium stripping and plating. Figure 5a presents CE results for a bare Li-Cu cell and a Li-Cu cell where the Li is coated with 8 nm LiF, both measured at 0.4 mA/cm2 with 20-μL electrolyte using 1-hour charge and discharge. For the bare cell, the CE is low during the initial ~5-10 cycles due to SEI formation that consumes lithium on both the lithium anode and the copper cathode, after which the CE stabilizes at 99.5%. In contrast, the LiF-coated cell shows much higher CE even in the initial cycles (before settling at 99.5%), suggesting reduced lithium depletion, and a thinner SEI compared to the bare Li case. Since SEI formation consumes liquid electrolyte, the LiF coated cells should better preserve the electrolyte during cycling. Indeed, the coated Li cell maintains a high CE for more than 170 charge-discharge cycles compared to only ~50 cycles for the bare cell, implying a greater than 3-fold enhancement in cycle life.

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(b)

(c)

(d)

Figure 5. Coulombic efficiency versus charge-discharge cycle number measured from Li-Cu coin cells with and without ALD LiF coatings (a) at 0.4 mA/cm2 with ALD LiF coating on Li (b) 1 mA/cm2 with ALD LiF coating on Li, and (c) at 1 mA/cm2 with ALD LiF on the Cu. (d) Nyquist plots of LiF-coated Li paired with copper cathode using different ether electrolyte volumes.

We repeated these experiments with fresh coin cells using a higher current density of 1 mA/cm2, again for 1-hour charge and discharge (Figure 5b). For both cells, the CE stabilized at a slightly lower value of ~98.5% compared to 99.5% at 0.4 mA/cm2, suggesting that Li stripping from Cu may be kinetically limited. As in the previous experiment, the bare Li-Cu cell failed after only ~40 cycles compared to >160 cycles for the LiF coated cell. It should be noted that ~5 microns lithium is stripped and plated during each charge-discharge cycle in these experiments. It is remarkable that the effects of the 8 nm LiF coating persists for >170 cycles when subjected to such large volume changes. To the best of our knowledge, these LiF films yield the greatest lifetime improvement for Li metal anodes of any coating at current density of 1 mA/cm2 to the capacity of 1 mAh/cm2 to date20, 29, 55-57. As discussed above, a low electrolyte volume is critical for practical lithium batteries incorporating Li anodes. We found that the LiF coatings provide a significant cycle lifetime enhancement even using an electrolyte volume of 5 µL (Figure S6). We performed EIS 15



measurements of LiF coated Li-Cu cells using different electrolyte volumes (Figure 5d). The depressed semicircles in the high and medium frequency regions represent the charge transfer resistance (Rct). The 20-μL electrolyte leads to the smallest Rct, suggesting that the electrolyte covers the electrode surface better compared to the 5 and 10 μL cases. We hypothesize that the lower electrode coverage for the 10-μL case increases the interfacial impedance, and may produce the CE fluctuations observed in Figure S6. Thus far, we have described Li-Cu cell testing where the LiF was deposited on the Li electrode. However, during cycling, the Li stripping and plating occurs on the Cu electrode as well, and this suggests that perhaps a LiF coating on the Cu electrode could improve cyclability. To test this hypothesis, we prepared a Li-Cu cell with bare Li and 6 nm LiF on the Cu. Figure 5c displays the CE data for this cell using a rate of 1 mA/cm2 and the results are nearly identical to the case where the Li is coated (Figure 5b). In fact, the CE is higher and the stability extends for more cycles in the LiF on Cu case (98.8%, 180 cycles) compared to the LiF on Li case (99.5%, 160 cycles). Figure 6 shows SEM images for a bare lithium anode (Figures 6a, 6b) and an ALD LiFcoated Li anode (Figures 6c, 6d) after 60 and 160 charge-discharge cycles, respectively, performed using Li-Cu coin cells under identical cycling conditions of 1 mA/cm2 and 1 mAh/cm2 using 20 μL electrolyte. The bare lithium anode (Figures 6a and 6b) is covered with lithium dendrites, and this is consistent the abrupt decrease in Coulombic efficiency at ~60 cycles observed in Figure 5b. In Figure 6a, small patches of relatively smooth surface are visible between the dendrites suggesting that the dendrites nucleate from discrete locations. In contrast, the ALD LiF-coated Li anode shows no evidence of dendrites even after160 cycles (Figure 6c and 6d). This suggests that the ALD LiF coating inhibits dendrite formation during charge-discharge cycling, and this produces a higher Coulombic efficiency and greater cycle life (Figure 5b).

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(a)

(c)

(b)

(d)

Figure 6. SEM images of lithium anodes following charge-discharge cycling in Li-Cu cells at 1 mA/cm2 and 1 mAh/cm2 using 20 μL electrolyte. (a) and (b) show bare Li anode after 60 charge-discharge cycles at 1000x (10 μm scale bar) and 5000x (1 μm scale bar) magnification, respectively; (c) and (d) show ALD Li-F coated Li anode after 160 charge-discharge cycles at 1000x (10 μm scale bar) and 5000x (1 μm scale bar) magnification, respectively.

Based on the remarkable battery performance and the interesting fact of lithium residues on separators in cycled cells, we believe that the electron-insulating and compositionhomogeneous LiF first creates a uniform current distribution to facilitate lithium stripping and depositing process. Thus, lithium is readily deposited onto the lithium, which is beneath of the LiF instead of staying on the anode surface. The ultrastable LiF also serves as a perfect isolating interlayer between lithium metal and the electrolyte to hinder any side reactions, which incredibly preserve electrolyte content and only requires a small volume for great battery operation. Furthermore, the high shear modulus of the LiF film suppresses lithium dendrites that would disrupt the coating. Evidently, the naturally formed SEI lacks the mechanical strength to endure the lithium dendrites while the very high shear modulus of 58 GPa (Table S1) for our LiF can resist the fractures.

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4. Conclusion In conclusion, we demonstrate for the first time a ALD chemistry to prepare LiF thin films that can be completed at low temperature using HF and LiOtBu at 0.82-0.88 Å/cycle. The assynthesized film is so pure with contamination level

Novel ALD Chemistry Enabled Low-Temperature Synthesis of Lithium

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