Improving ionic conductivity with bimodal-sized Li7La3Zr2O12 fillers

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Improving ionic conductivity with bimodal-sized Li7La3Zr2O12 fillers for composite polymer electrolytes Yan Sun, Xiaowen Zhan, Jiazhi Hu, Yikai Wang, Shuang Gao, Yuhua Shen, and Yang-Tse Cheng ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b21770 • Publication Date (Web): 11 Mar 2019 Downloaded from http://pubs.acs.org on March 12, 2019

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ACS Applied Materials & Interfaces

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Improving ionic conductivity with bimodal-sized Li7La3Zr2O12 fillers for composite polymer

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electrolytes

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Yan Suna, b, †, Xiaowen Zhanb, †, Jiazhi Hub, Yikai Wangb, Shuang Gaob, Yuhua Shena, *, Yang-Tse Chengb, *

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aCollege

of Chemistry and Chemical Engineering, Lab for Clean Energy & Green Catalysis, Anhui University, Hefei 230601, China

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bDepartment

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†Both

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Email: [email protected]; [email protected]

of Chemical and Materials Engineering, University of Kentucky, Lexington 40506, USA

authors contributed equally to this work.

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Abstract: Ceramic-polymer composite electrolytes (CPEs) are being explored to achieve both

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high ionic conductivity and mechanical flexibility. Here, we show that, by incorporating 10 wt%

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(3 vol%) mixed-sized fillers of Li7La3Zr2O12 (LLZO) doped with Nb/Al, the room temperature

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ionic conductivity of a polyvinylidene fluoride (PVDF)-LiClO4-based composite can be as high as

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2.6 × 10 ―4 S/cm which is one order of magnitude higher than that with nano- or micro-meter

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sized LLZO particles as fillers. The CPE also shows a high lithium-ion transference number of

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0.682, a stable and low Li/CPE interfacial resistance, and good mechanical properties favorable

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for all-solid-state lithium-ion battery applications. XPS and Raman analysis demonstrate that the

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LLZO fillers of all sizes interact with PVDF and LiClO4. High packing density (i.e., lower porosity)

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and long conducting pathways are believed responsible for the excellent performance of the

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composite electrolyte filled with mixed-sized ionically conducting ceramic particles.

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Keywords: Lithium batteries, Li7La3Zr2O12, composite polymer electrolytes, ionic conductivity,

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ceramic-polymer interactions

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1. Introduction

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Lithium-ion batteries (LIBs) are ubiquitous in applications ranging from portable electronic

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devices to electric vehicles (EVs).1, 2 Most rechargeable LIBs contain organic liquid electrolytes,

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which may catch fire and cause explosion. Replacing organic liquid electrolyte with a solid

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electrolyte in all-solid-state LIBs is a promising approach to increasing volumetric energy density

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and cycle life as well as improving the safety of batteries. Inorganic solid electrolytes, such as

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oxides and sulfides, and solid polymer electrolytes (SPEs), such as polyethylene oxide (PEO)-

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based electrolyte, are two families of solid-electrolytes. The former has the advantage of high ionic

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conductivity, while the latter has better mechanical flexibility. Ceramic-polymer composite

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electrolytes (CPEs), consisting of inorganic fillers in polymer matrixes, have been explored to

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achieve both high ionic conductivity and mechanical flexibility.

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Several CPEs, including LLZO-PEO,3-5 LLZO-polyvinylidene fluoride (PVDF),6,

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LLZO-

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polypropylene carbonate (PPC)8, and LLZO-polyacrylonitrile (PAN),9 have shown enhanced ionic

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conductivity, chemical/electrochemical stability, and mechanical flexibility. The highest ionic

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conductivity was always observed at low LLZO mass contents that are below the percolation

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threshold of ionic conduction through LLZO particles alone. For example, Chen et al.3

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demonstrated that 10 wt% Li6.4La3Zr1.4Ta0.6O12-PEO-LiClO4-based CPE presented a maximum

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ionic conductivity of 1.15 × 10 ―4 S cm-1 at 25 C. Zhang et al.8 reported that free-standing

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poly(propylene carbonate)-5 wt% nano-sized Li6.75La3Zr1.75Ta0.25O12 CPE exhibited high ionic

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conductivity (5.2 × 10 ―4 S cm-1) at 20 C. This indicates that Li ions mainly pass through Li salts

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in the polymer matrix while LLZO fillers enhance the conductivity of the matrix by (1) suppressing

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the polymer crystallization thereby increasing the free volume available for motion of polymer

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chain segments and (2) facilitating the disassociation of Li salts to increase mobile Li ion

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concentration for conduction.6 Yang et al.9 investigated the morphology influence of ceramic 2 ACS Paragon Plus Environment

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fillers on ionic conductivity, and showed that the PAN-LiClO4-based CPE with LLZO nanowires

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exhibited significantly higher ionic conductivity than the one with LLZO nano particles. They

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suggested, based on NMR measurements, that the nanowires could partially modify the PAN

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matrix to create faster pathways for Li-ion conduction.9 Thus, a connected pathway is critical to

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fast Li-ion transport in CPEs, which may be accomplished by optimizing the morphology and

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volume fraction of the inorganic fillers. However, the size effects of the ceramic particles on the

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ionic transport and mechanical properties of CPEs are still unknown.

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Thus, this study focuses on the size influence of ceramic fillers on the electrochemical

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performance and mechanical behavior of CPEs. A modified sol-gel approach with a single heating

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step was employed to prepare nano-, micro-, and mixed-sized Nb/Al-doped LLZO

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(Li6.03La3Zr1.75Nb0.25Al0.24O12) (thereafter referred to as NLLZO, MLLZO, and XLLZO,

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respectively). CPE membranes consisting of PVDF, LiClO4 salt, and LLZO fillers of various sizes

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(thereafter referred to as N-CPE, M-CPE, and X-CPE, accordingly) were then prepared by a simple

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and scalable slurry casting method. The mechanisms responsible for significant conductivity

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enhancement in the CPE with mixed-sized LLZO fillers have been revealed. In addition,

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mechanical properties of CPEs were obtained by nanoindentation measurements.

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2. Experimental section

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Material Synthesis: Nb/Al-doped LLZO (Li6.03La3Zr1.75Nb0.25Al0.24O12) was synthesized with a

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modified sol-gel method similar to that described elsewhere.5 In a typical experiment, 0.0156 mol

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citric acid monohydrate (C6H8O7•H2O, Alfa Aesar, 99%), 0.0156 mol ethylene glycol (C2H6O2,

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Alfa Aesar, 99%), 0.014 mol lithium nitrate (LiNO3, Sigma Aldrich), 0.006 mol lanthanum (III)

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nitrate hexahydrate (La(NO3)3•6H2O, Alfa Aesar, 99.9%), 0.0035 mol zirconium (IV) 3 ACS Paragon Plus Environment


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acetylacetonate (C20H28O8Zr, Alfa Aesar), 0.00048 mol aluminum nitrate nonahydrate

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(Al(NO3)3•9H2O, Alfa Aesar, 98.0-102%), and 0.0005 mol niobium (V) chloride (NbCl5, Alfa

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Aesar, 99.9%) were dissolved in a 40 ml mixture of deionized water and ethanol under vigorous

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stirring at 80 C to obtain a homogeneous solution5, 10, 11. Subsequently, the solution was stirred

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overnight at 90 C to evaporate the solvent, and the resultant dry xerogel10, 11 was then heated at

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250 C for 2 h to form a highly porous brown gel. The gel was calcined at 850 C for 2 h to obtain

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nano-sized LLZO (NLLZO) or at 1100 C for 12 h to form MLLZO. XLLZO powders were

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obtained by simply mixing NLLZO and MLLZO at the weight ratio of 1:1. For preparation of

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sintered pellets, NLLZO powders were first ground in acetone using a planter mill (300 rpm/3 h).

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Subsequently, the powders were pelletized using cold isostatic press at 200 MPa and then sintered

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at 1100 C for 12 h. During sintering, the pellets were placed over a powder bed of identical

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composition at the bottom of the crucible to prevent reactions with alumina and covered by the

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same powder to reduce lithium loss. An extra 5 wt% lithium was used in the preparation of

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MLLZO or sintered pellets to compensate lithium evaporation at 1100 C.

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Preparation of CPEs. CPEs were prepared using a slurry casting method. Polyvinylidene fluoride

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(PVDF, Arkema, Kynar 741) and lithium perchlorate (LiClO4, Alfa Aesar, 99.99%) were dried at

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100 C under vacuum for 12 h to remove water. PVDF and LiClO4 were dissolved in N, N-

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dimethylformamide (DMF) under stirring at 50 C for 4 h to obtain a homogeneous solution. The

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PVDF concentration is 10 wt.% and the ratio (by weight) of PVDF and LiClO4 is 3:1. The LLZO

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powders were added to the solution with a weight percentage of 10% in the total amount of PVDF

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and LLZO. The homogeneous solution was cast onto a clean glass plate using a doctor blade and

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then dried in vacuum at 60 C for 16 h to obtain the CPEs. Subsequently, the CPEs were punched

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into round-shaped membranes with a diameter of ½ inch and pressed at 200 MPa using a uniaxial 4 ACS Paragon Plus Environment

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press inside the glovebox before further measurements. For comparison, PVDF-SPE without

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LLZO fillers was prepared. Additionally, CPEs with 20, 30, and 40 wt% XLLZO were fabricated

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to explore the influence of ceramic loading.

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Material Characterization: Phase analysis was conducted using X-ray diffraction (XRD; Siemens

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D5000, Cu Kα, 40 kV, 30 mA). The thermal stability of CPEs was evaluated by thermogravimetry

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analysis (TGA; NETZSCH STA 449 F3 Jupiter) by heating the specimens to 600 C in N2

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atmosphere with a ramp rate of 10 C min-1. Microstructure and composition were characterized

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using scanning electron microscopy-energy dispersive spectroscopy (SEM-EDX; Quanta 250) and

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the grain size distribution was determined using ImageJ. The surface chemistry of CPEs was

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studied using X-ray photoelectron spectroscopy (XPS, Thermo Scientific KAlpha). To avoid

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impurities such as LiOH and Li2CO3 formed on the surface, the samples were transferred from an

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Ar-filled glovebox into the XPS chamber with a vacuum sample holder. Raman spectra were

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obtained with a DXR micro-Raman instrument (Thermo Scientific). A diode pumped Nd:YVO4

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laser was used as excitation source (532 nm wavelength) for Raman characterization.

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The n-butanol adsorption method was used to determine the porosities of CPEs. The electrolyte

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membrane was immersed in n-butanol for 2 h and was then dried using filter papers. The porosity

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values were calculated according to Eq. (1):12 (𝑚𝑤 ― 𝑚𝑑)

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𝑝𝑜𝑟𝑜𝑠𝑖𝑡𝑦 =

(𝑚𝑤 ― 𝑚𝑑)

𝑛

𝑚 𝑛 + 𝑑 𝑝

(1)

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where mw and md are the weight of the wet and dry membranes, respectively, and n and p the

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densities of n-butanol and PVDF, respectively.



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Mechanical Test: Mechanical properties of PVDF-SPE and CPEs were measured by an

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instrumented nanoindentation (with a Berkovich indenter) inside an argon-filled glovebox. To

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study the instantaneous modulus of these membranes, nanoindentation relaxation measurements

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were conducted with a constant loading/unloading rate of 20 nm/s. The indenter was held at the

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maximum depth (3000 nm) for 100 s to monitor the relaxation of the load. The instantaneous

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modulus was determined by Eq. (2),13

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𝐸(0) =

|

𝑑𝐹

1 ― 𝑣2 𝑑𝐹 2𝑎 [𝑑ℎ ℎ = ℎ 𝑚

|

+

𝑑𝐹/𝑑𝑡|𝑡 ― 𝑚

𝑣ℎ

]

(2)

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where

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before unloading, and 𝑣ℎ is the unloading rate. The elastic modulus and hardness were derived

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from load-displacement curves based on Oliver-Pharr’s method.14 Reversible work, Wu (the total

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area under the unloading curve), to total work, Wt (the total area under the loading curve) were

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calculated based on load-displacement curves of depth-controlled nanoindentation measurements.

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Since ℎ𝑚𝑎𝑥< 1/10 thickness of membranes in both test modes, the substrate effect is negligible.

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Electrochemical Testing: Ionic conductivity was measured by the AC impedance method with two

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stainless-steel (SS) foils as ion-blocking electrodes in the frequency range of 6 MHz to 1 Hz using

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a Solatron 1260 analyzer. In the case of pellet samples, gold paste was used and cured at 700 °C

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for 1 h. The ionic conductivity () was calculated from the thickness of electrolyte film (L),

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electrode area (S), and electrolyte resistance (R), according to Eq. (3):

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is the initial unloading slope,𝑑𝐹/𝑑𝑡|𝑡𝑚― is the load relaxation rate immediately

𝑑ℎ ℎ = ℎ 𝑚

𝐿

 = 𝑅𝑆

(3)

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The linear sweep voltammetry (LSV) was measured in the potential range from 2.5 to 6.0 V at a

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scanning rate of 1 mV s-1 in a SS/CPE/Li asymmetric cell. Lithium-ion transference number was 6 ACS Paragon Plus Environment

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determined using a combination of DC polarization and AC impedance measurements in a

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Li/CPE/Li symmetrical cell configuration based on Eq. (4):15 𝐼𝑠𝑠(𝑉 ― 𝐼0𝑅0)

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𝑡𝐿𝑖 + = 𝐼 ( 0

𝑉 ― 𝐼𝑆𝑆𝑅𝑆𝑆)

(4)

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where I0 and Iss denote respectively the initial and steady-state current in response to a DC

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polarization potential of 10 mV. R0 and Rss are respectively initial and steady-state interfacial

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resistances and can be extracted from the impedance spectra obtained before and after polarization

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of the cell.

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3. Results and discussion

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Figure 1. (a) XRD patterns of calculated LLZO, NLLZO, MLLZO, PVDF, Kapton tape, and X-

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CPE sealed with Kapton tape. SEM images of (b) NLLZO, (c) MLLZO, (d) N-CPE, (e) M-CPE,

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and (f) X-CPE. (g-h) Photographs of flexible X-CPE films. (i) A cross-sectional SEM image and

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(j-m) EDS mappings of X-CPE.

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As seen in Figure 1a, the XRD diffraction peaks of NLLZO and MLLZO powders match well

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with those in the calculated pattern, indicating that both materials crystalize into a single cubic 8 ACS Paragon Plus Environment

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phase. The four peaks at 18.5, 20.1, 26.9, and 38.7 can be assigned to the (020), (110), (022), and

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(211) planes of the γ-phase of PVDF, respectively.16 Major peaks of LLZO remain in the pattern

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of X-CPE, implying that the crystal structure of LLZO is well retained. However, no obvious

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diffraction peaks for PVDF can be observed, suggesting that the degree of crystallinity of PVDF

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in X-CPE is reduced. This may be caused by the interactions between the dispersed LLZO particles

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and the PVDF matrix.6, 7, 17 The small peak at 21.1 in the XRD pattern of X-CPE is most likely

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from LiClO4, suggesting crystallization of undissolved excess salt in the CPE. SEM images of

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LLZO powders with different particle sizes are shown in Figure1b-c. The NLLZO shows crystal-

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like particles while the MLLZO powders are round-shaped with a considerable degree of

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agglomeration. The average cluster sizes of NLLZO and MLLZO are 0.44 and 4.31 m (Figure

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S1), respectively, confirming the formation of nano- and micro-sized LLZO samples by applying

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different heating protocols. The EDX mappings of NLLZO and MLLZO (Figure S2) reveal

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uniform distributions of Al, Nb, La, Zr, and O. The morphologies of CPEs made with NLLZO,

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XLLZO, and MLLZO particles are compared in Figure 1d-f. All CPE membranes have similar

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surface features, with branch-like strips spreading over the porous matrix. From the Cl mapping

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in Figure S3, the strips are LiClO4 salt precipitates. This confirms the recrystallization of excess

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LiClO4 as identified in the XRD pattern (Figure 1a).18 Based on the F and Al signals, PVDF

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constitutes the matrix in which LLZO particles are uniformly distributed. Enlarged views of the

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matrixes are presented in the inserts in Figure 1d-f. The pore size of X-CPE is much smaller than

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that of N-CPE and M-CPE. The SEM images of X-CPEs with different filler contents are shown

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in Figure S4. As the loading of XLLZO particles increases from 10 to 30 wt%, the matrix becomes

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slightly more porous and the membrane becomes darker. When the XLLZO content reaches 40

23

wt%, it is difficult to cast a membrane from the slurry (Figure S5). This will be further discussed



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in combination with porosity and conductivity results. Figure 1g-h show a flexible X-CPE

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membrane of a thickness of 160 m. The cross-section SEM image together with the EDX

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mapping of Cl (Figure 1i-m) indicates that many LiClO4 strips are aligned approximately along

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the thickness direction. The well-overlapped Al/La (from LLZO) and F signals (from PVDF)

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suggest that LLZO particles are evenly dispersed in the PVDF matrix.

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Figure 2. (a) Impedance spectra, (b) room-temperature ionic conductivities of CPEs with 10 wt%

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LLZO fillers and PVDF-SPE. (c) Arrhenius plots for a sintered LLZO pellet and N-, X- and M-

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CPE. The equivalent circuit used to fit the impedance spectra is shown in the bottom of panel a.

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Figure 2a shows the Nyquist plots of PVDF-based polymer electrolyte (PVDF-SPE), N-, X-,

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and M-CPEs, respectively, measured at room temperature using stainless steel foils as ion-

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blocking electrodes. The resistance values of CPEs were obtained by fitting spectra with an

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equivalent circuit model (Figure 2a). As compared in Figure 2b, X-CPE exhibits an ionic

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conductivity of 2.6 × 10 ―4 S/cm at 25 C, which is higher than that of N-CPE (1.0 × 10 ―4 S cm-1)

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and M-CPE (8.1 × 10 ―5 S cm-1). The ionic conductivity of PVDF-SPE membrane without LLZO

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fillers was 1.4 × 10 ―5 S cm-1 at 25 C, which is 17 times lower than that of X-CPE. Table S1 10 ACS Paragon Plus Environment

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collects the ionic conductivity values of these membranes from different batches with the

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corresponding standard errors displayed in Figure 2b, which indicates excellent reproducibility of

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the results. The Arrhenius plots of the three different CPEs are shown in Figure 2c, with the data

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for a sintered LLZO pellet included for comparison. Both the room-temperature ionic conductivity

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(3.6 × 10 ―4 cm-1 at 22 °C) and activation energy (0.32 eV) of the as-sintered LLZO pellet are

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consistent with previous reports for c-LLZO.6, 7, 19-21 In addition to its high ionic conductivity, X-

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CPE also has a slightly lower activation barrier (0.33 eV) than that of N-CPE (0.39) and M-CPE

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(0.38 eV), indicating that Li-ion migration in X-CPE is energetically more favorable.6

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Figure 3. XPS spectra of (a) F 1s and (b) Cl 2p for different membranes. (c) Porosities of PVDF-

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SPE, N-CPE, M-CPE, and X-CPE. (d) Possible Li-ion pathways in CPEs with 10 wt% LLZO

3

fillers.

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To identify interactions among distinct components in CPEs, XPS spectra of PVDF-SPE and

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different CPEs were analyzed in Figure 3a-b. The major F1s peaks at 687.4 and 684.8 eV in PVDF-

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SPE are assigned to the F-C and F-Li bonds, respectively.22, 23 Incorporation of LLZO fillers,

8

regardless of their sizes, results in a higher peak intensity of F-Li in CPEs than in PVDF-SPE,

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which can be attributed to the formation of LiF resulting from interactions between LLZO particles

10

and PVDF chains. Raman spectra provide additional evidence for this structural modification of

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PVDF by LLZO. As show in Figure S6, the peak at 2980 cm-1 of PVDF-SPE corresponds to the

12

CH2 bending vibration mode.

13

deprotonation of CH2.24 The rise of the two peaks at 1121 cm-1 and 1510 cm-1 for CPEs are

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characteristics of C=C stretching vibration mode of polyenes, revealing the dehydrofluorination

15

of PVDF chain,25 which is consistent with the suggestion that LLZO can create an alkaline-like

16

condition and cause the partial dehydrofluorination of PVDF.6, 24, 26 This structural change induced

17

by LLZO can account for the abovementioned color difference between PVDF-SPE and CPEs,

18

and also strongly supports the tendency to form Li-F bonds in CPEs as demonstrated by XPS. In

19

the Cl 2p scans (Figure 3b), the characteristic peak at 208.7 eV present in all can be ascribed to

20

the ClO4- group.27 Interestingly, an additional peak at 206.8 eV corresponding to Cl in ClO3-

21

enviroment28 appears in the CPEs, indicating interactions between Li salt and LLZO.29 The XPS

22

spectra of Li1s and O1s are also collected in Figure S7. For all CPEs, the major peaks can be

23

assigned to Li (56.9 eV) or O (533.3 eV) in LiClO4. The contribution from LLZO is not seen

The absence of this peak in the three CPEs indicates the


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probably due to its low surface concentration. Although the Nb3d and Cl2p peaks overlap in the

2

190–220 eV range, the XPS peaks shown in Figure S7c should be from Cl2p instead of Nb3d since

3

other elements in the Nb/Al doped LLZO were undetected.

4

Based on XPS and Raman results, we propose four possible mechanisms responsible for

5

conductivity enhancement due to LLZO: (1) increased Li+ concentration as a result of acid-base

6

interactions between LLZO and dehydrofluorinated PVDF,6, 30 (2) enhanced Li salt disassociation

7

caused by the interactions between LLZO and Li salts as indicated by the formation of ClO3-, (3)

8

enhanced segmental motion of PVDF chains as a result of reduced crystallinity for PVDF,6, 26 and

9

(4) provided sites for Li+ hopping and can therefore enhance conductivity.

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Moreover, as shown in Figure 3c, the porosity of X-CPE (12.5%) derived from the n-butanol

11

absorption method6, 31 is smaller than that of N-CPE (24.8%) and M-CPE (34.9%). The same trend

12

was evidenced by the porosities estimated based on the volume of individual components as

13

detailed in Table S2 (Supporting Information). This is unsurprising since, based on the Furnas

14

model,32 the theoretical maximum packing efficiency of a mixture of coarse and fine particles is

15

always higher than that of single-sized fine or coarse particles.32, 33 Because pores can block Li-

16

ion migration pathways that are critical to fast ion conduction, we believe that the higher

17

conductivity of X-CPE compared to N-CPE and M-CPE is caused by its low porosity. Several

18

factors may affect the porosity of the CPE membranes after compression. For example, the

19

pressure (300 MPa in present study) may be insufficient to cause flow of polymers to fill the pores,

20

especially when the pore diameter is small (e.g., Laplacian pressure effect). In addition, the

21

interface between polymer and inorganic particle may not be sufficiently strong to sustain the

22

elastic rebound of the polymer, causing delamination and, thus, formation of pores at the interfaces

23

upon the removal of the external pressure. Furthermore, closed pores may also re-open in the



1

absence of external pressure.34 All these issues can contribute to the residual porosity in the

2

membranes.

3

A schematic illustration of ionic conduction mechanisms in CPEs is shown in Figure 3d: fast

4

Li-ion migration channels exist because of Li salts embedded in the PVDF matrix. Additional Li-

5

ion pathways are created by LLZO particles. For a given mass (volume) fraction of LLZOs, the

6

Li-ion conductivity increases with decreasing porosity because LLZO particles can facilitate Li-

7

ion hopping, shortening jump distance, and lowering activation barrier. As a result, X-CPE with

8

the lowest porosity has the highest ionic conductivity.

9

The thermal stability of different membrane samples was characterized using thermogravimetric

10

analysis (TGA). The initial weight loss (below 150 C) for PVDF-SPE shown in Figure S10 is due

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to the evaporation of trapped water molecules. Most lithium salts such as LiClO4 are usually

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hygroscopic as their small, highly polarizing lithium ions interact strongly with water molecules

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(i.e., LiClO4·3H2O).35 Interestingly, water absorption of CPEs with LLZO fillers is negligible. One

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possible reason is that the interaction between LLZO and LiClO4 facilitates the disassociation of

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LiClO4, which weakens the polarity of lithium ions and reduces H2O absorption. The weight loss

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of N-, M-, and X-CPE due to melting and gradual degradation of PVDF start at much lower

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temperatures (

Improving ionic conductivity with bimodal-sized Li7La3Zr2O12 fillers

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