untitled ISSN 215 One‐pot PVdF‐H Agnieszka Mariusz W 1 Institute of No 2 Center of Adva * Corresponding Tel.: +48.061.27 ARTICLE IN DOI: 10.5155/e Received: 10 No Received in rev Accepted: 19 De Published onlin Printed: 31 Mar KEYWORDS Li‐ion battery Inorganic fillers Synthesis desig Sol‐gel processe Supercritical CO Composite gel e 1. Introduct Lithium energy stora mobile devic based on liq of gel‐like el operational polymer ma meet a numb stability rela good compa mechanical s good therma In rece propylene) ( electrolyte in the addition the polyme mechanical s The sub carbon dioxi on PVdF‐HF 53‐2249 (Print) t method FP compo a Martyla 1, R Walkowiak 1 n‐Ferrous Metals, D anced Technologies g author at: Institu 797815. Fax: +48.06 FORMATION eurjchem.7.1.121‐1 ovember 2015 vised form: 17 Dece ecember 2015 ne: 31 March 2016 rch 2016 S s gn es O2 drying electrolytes tion ion cells are o age technologie ces and vehicle uid electrolytes lectrolytes, whi safety and mi trix which is t ber of requirem ated with redox atibility with li strength and to al resistance [5] nt years, po (PVdF‐HFP) cop n lithium ion b of small partic r matrix imp stability and lith bject of this w ide as a drying FP. In a furth E / ISSN 2153‐225 htt Europ of synthe osite mem Robert Prze and Maciej Division in Poznan, s Adam Mickiewicz te of Non‐Ferrous M 61.2797897. E‐mai 127.1361 ember 2015 one of the faste es. This applie es [1‐3]. These s [4]. New tech ich are suppose inimizing the r o act as the so ments: have a g x processes at t iquid electrolyt o be easy to ma ]. ly(vinylideneflu polymer was w batteries. It is w cles of metal ox proves conduct hium cationic p work was the agent for comp her step com uropean Journal Europ 57 (Online)  20 tp://dx.doi.org/ pean Jo Journal web esis and su mbranes ekop 2, Monik Kopczyk 1 Central Laboratory University, 89c Um Metals, Division in P l address: agnieszk ABSTRACT This paper de gel electrolyt the precursor butoxide (Ti( supercritical electrochemic electrochemic Cite this: Eur. est growing are es both to pow e cells are prim hnologies involv ed to ensure gr risk of spillage olid electrolyte good electroche he lithium elec tes, low price, anufacture and uoride‐co‐hexa widely tested as well known fac ides (TiO2, SiO2 tivity, thermal properties [6‐9] use of superc posite systems posite polyme l of Chemistry 7 pean Journal of C 016 Atlanta Pub 10.5155/eurjche ournal bpage: www. upercritic ka Osinska‐ ry of Batteries and C multowska St., 61‐61 Poznan, Central La ka.martyla@claio.p escribes one‐pot tes based on po r of the inorga OC4H9)4) were u CO2. The memb cally, showing cal characteristic . J. Chem. 2016, eas of wering marily ve use reater e. The must emical ctrode, good d have fluoro s solid ct that 2) into l and . critical based er gel elec rou TiO SiO sol‐ syn con app dry sup rize IR. abs the elec test 2. E 2.1. ray (1) (2016) 121‐ Chemistry lishing House LL em.7.1.121‐127. of Che .eurjchem.co cal carbon Broniarz 1, Cells, 12 Forteczna 14 Poznan, Poland aboratory of Batter poznan.pl (A. Marty t method of the oly(vinylideneflu nic filler tetrae used. The drying ranes and gel e favorable pro cs in Li‐ion cells. 7(1), 121‐127 ctrolyte for Li‐ us polymer matr O2 systems drie 2 and SiO2‐TiO ‐gel method fr nthesis was ca ndensation of plication of dib ying of membra percritical CO2. ed using XRD an Specific surfa orption isother dry composit ctrolyte, were s ts. Experimental . Instrumentat Phase identific diffraction (X ‐127 LC ‐ All rights re .1361 emistry m n dioxide a St., 61‐362, Pozna ries and Cells, 12 Fo yla). e synthesis and uoride‐co‐hexafl thyl orthosilica g of membranes lectrolytes have perties in ter . ion batteries b rix and a family ed by this meth 2‐PVdF‐HFP sy rom TEOS and arried out wit reactants in a buthyl phthalat anes with filler The structure o nd SEM techniq ace areas we rms. Gel electro e membranes subjected to a w tion cation of sampl XRD) powder served ‐ Printed y drying of n, Poland orteczna St., 61‐362 properties of co luoropropylene) te (TEOS) and with fillers was e been examined ms of electrol based on PVdF y of SiO2, TiO2 an hod were chara stems were sy d Ti(OC4H9)4 p thin the polym a one‐step me te as porosity rs was also ca of membranes ques, thermal an re determined olytes, prepared in lithium con wide range of e es was perform diffractometer d in the USA f 2, Poznan, Poland. omposite polym ) (PVdF‐HFP). A titanium tetra‐ s carried out wi d structurally an lyte uptake an ‐HFP micropo‐ nd binary SiO2‐ acterized. TiO2, ynthesized by a recursors. The mer matrix by thod with the promoter. The arried out with was character‐ nalysis and FT‐ d by nitrogen d by soaking of nducting liquid electrochemical med using an X‐ r (Philips, PW er As n‐ th nd nd ‐ ‐ , a e y e e h ‐ ‐ n f d l ‐ W 122 Martyla et al. / European Journal of Chemistry 7 (1) (2016) 121‐127 1050) using CuKα radiation and a Ni filter. The diffraction pattern were recorded in the range of 5 to 90 [2Θ], with a step size of 0.04/[2Θ] and 2s/step speed. The percent of crystallinity was determined from XRD data and it was calculated by dividing the total area of crystalline peaks by the total area under the diffraction curve (crystalline plus amorphous peaks) [10] using Automatic Powder Diffraction Philips software. The real area of the crystalline peaks and the amorphous peak can be determined from/by a computer software package performing a mathematical deconvolution of the peaks. The percentage of the crystalline polymer can be determined from Equation 1. % Crystallinity 100 (1) For the cross‐section observation, the microporous membrane was freeze‐fractured in liquid nitrogen and then sputter coated. Imaging the surface was performed with SEM electron microscopy (Phillips, 515 SEM). Thermal stability of membranes was monitored by thermogravimetric analysis (TGA) and thermogravimetric analysis/differential thermal analysis (TGA/DTG) (Netzsch STA 409C 3F). TGA measurements were carried out under a nitrogen atmosphere at heating rate of 15 °C/min from 25 to 800 °C. The samples’ weight was in the range of 5‐10 mg. Nitrogen was used as a carrier gas with flow rate of 25 mL/min. Surface studies were carried out using a Fourier Trans‐ form Infrared Spectroscopy/Attenuated Total Reflectance (FT‐ IR/ATR) with FT‐IR spectrometer (Bruker, TENSOR 27) from ATR accessory (SPECAC). Measurement resolution was 4 cm‐1. The porous structure was determined by low temperature (‐196 °C) nitrogen adsorption measurements carried out on the Accelerated Surface Area and Porosimetry system model 2010 made by Micromeritics, using 200‐300 mg of sample with the grain size fractions between 0.1 and 0.2 mm. Prior to nitrogen adsorption, all samples were outgassed at 623 K, at 0.4 Pa until a constant mass was reached. Both adsorptive and desorptive branch of the isotherm was taken in the range of p/p0 0‐1. For the study of liquid phase uptake, small pieces were cut off from dry membranes. Once carefully dried and weighted, they were immersed in a container with pure propylene carbonate (anhydrous, 99.7%, Aldrich). At predefined moments in time, the membrane pieces were removed from the container, blotted lightly from the excess of liquid, weighted and immediately placed again in the container. Liquid phase uptake was followed as a result of swelling membrane weight increase: Membrane weight increase = (mt‐m0)/m0 × 100 [%] (2) where m0 is the weight of dry membrane, mt is the weight of membrane after a given time of swelling. The ionic conductivity of composite gel polymer electroly‐ tes is calculated from: σ = L/(Rb×A) (3) where L and A represent thickness and area of the electrolyte specimen, respectively. Rb is the bulk resistance of the gel electrolyte obtained from complex impedance measurements. For the conductivity measurements gel electrolytes were prepared by immersing round pieces of dry composite membranes (10 mm diameter) for 1 hour in a liquid electrolyte typical for Li‐ion batteries, consisting of 1 M solution of LiPF6 (Aldrich 99.99%) in 1:1 (w:w) mixture of ethylene carbonate (EC, anhydrous, Sigma‐Aldrich 99%) and dimethyl carbonate (anhydrous, Sigma‐Aldrich 99%). In the next step, the round pieces of membranes swelled in the liquid electrolyte were placed in a two‐electrode Swagelok‐type cell with stainless steel electrodes. All these operations were conducted in a glove box, in dry argon atmosphere. The conductivities were determined at several temperatures (10‐ 60 °C) on the basis of impedance spectra obtained by means of PARSTAT 2263 (Princeton Applied Research) impedance analyzer in the frequency range of 100 kHz ‐ 1 Hz with 10 mV AC amplitude. Typically, each measurement was repeated several times to ensure good reproducibility of results. The cells were thermostated during measurements in a climatic chamber (Vötsch). Activation energy for composite polymer gel electrolytes have been calculated from the equation: σ = σ0 exp (‐Ea /kT) (4) where σ0 is the pre‐exponential factor, Ea is the activation energy, T is the absolute temperature in Kelvin scale and k is the Boltzmann constant. 2.2. Synthesis The poly(vinylidenefluoride‐co‐hexafluoropropylene) was purchased from Kynar Flex, Atofina. The dibutyl phthalate from Merck, tetraethyl orthosilicate, tetrabutyl titanate and acetic acid were from Alfa Aesar. The PVdF‐HFP gels were prepared according to a method similar to the so‐called Bellcore. PVdF‐HFP copolymer was added to acetone together with dibutyl phthalate and precursor of inorganic fillers. Inorganic filler was created in one step method with the formation of the membrane by the simultaneous hydrolysis and condensation of the corresponding alkoxide precursor [11]. CH3COOH + Si(OEt)4 → Si(OEt)3(CH3COO) + EtOH (5) EtOH + CH3COOH → CH3COOEt + H2O (6) H2O + Si(OEt)3(CH3COO) → HOSi(OEt)3 + CH3COOH (7) 2 HOSi(OEt)3 → (EtO)3Si‐O‐Si(OEt)3 + H2O (8) CH3COOH + Ti(OEt)4 → Ti(Ot‐Bu)3(CH3COO) + t‐BuOH (9) t‐BuOH + CH3COOH → CH3COOt‐Bu + H2O (10) H2O + Ti(Ot‐Bu)3(CH3COO) → HOTi(Ot‐Bu)3 + CH3COOH (11) 2 HOTi(Ot‐Bu)3 → (t‐BuO)3Ti‐O‐Ti(Ot‐Bu)3 + H2O (12) Such a synthesis strategy results from the experiments described in the work for the preparation of the gel systems of the oxides and hydroxides by hydrolysis and condensation in anhydrous solvents [12,13]. Tetraethyl orthosilicate (TEOS) was used as a source of silica. Tetrabutyl titanate (Ti(OC4H9)4) was a titanium oxide source. The weight ratio of the filler precursor (as final oxide) to the copolymer was 1:10. Acetic acid (CH3COOH) was added to methanol (CH3OH) as a hydrolysis and condensation agent. Mixtures were stirred and heated at 45 °C for several hours. Each solution was cast on a glass plate, covered with a Petri dish and left for slow evaporation (5 days). The resulting membranes were immersed in an apparatus for drying under supercritical conditions of CO2. The system was maintained at constant pressure of 1250psi and temperature of 35 °C for 6h. Next, the system was slowly depressurized for about 2 h at the same temperature. As soon as the depressurization process was completed, the resulting membranes were collected and kept in a desiccator for further use. The resulting samples were named respectively: PVdF‐HFP, PVdF‐HFP‐SiO2, PVdF‐ HFP‐TiO2, PVdF‐HFP‐SiO2‐TiO2. 3. Results an It is we three struct crystallizatio observed rec [2Θ] 18.4, 2 110, and 021 modified m inorganic co to the phas confirmed b additive of th at 38.9 °, as The smalle membranes, reflections. The inte indicating th composite m crystallinity PVdF‐HFP m crystallinity SiO2‐TiO2 as crystallinity Figu Figure 2. SEM i nd discussion ll known that tures: α, β, an on conditions. cently (Figure 0.0, 26.5 and 1 planes of crys membranes are omponent mixe se changes in y the appearan he SiO2 increas in the case of est changes w which caused nsities of the c he more amo membranes. W of the compos membranes wa of PVdF‐HFP m s additives was of doped mem Martyla et al ure 1. XRD pattern image of composite PVDF can cry d γ type [14] α type cryst 1A). Peaks of 38.9 ° corresp stalline PVdF. D e shown on F ed with the cry n the hybrid nce of a non‐cr ses the loss of i f binary oxide c were observe d only a decre crystalline peak orphous type We observe a site. The calcul as approximate membranes form s 26, 5 and 5% mbranes is low l. / European Jou n of A) PVdF‐HFP, B e membranes: A) P ystallize mainly , depending o tal structures interest here a onding to 100 Diffractograms Figures 1B‐1D ystalline polym matrix, which rystalline phase intensity of the component add ed for TiO2‐d ease in intens ks further decre morphology o also the chan lated crystallin ely 28%, whil med using TiO2 %, respectively wer than pure urnal of Chemistr B) PVdF‐HFP‐SiO2, PVdF‐HFP, B) PVdF y into on the were are at 0, 020, of the . The er led h was e. The e peak dition. doped ity of eased, of the ge of nity of le the 2, SiO2, y. The PVdF‐ HFP inor [15 rela and furt incr the mor drie cha mem is c sph Dur poly 2D) Add sur also ry 7 (1) (2016) 1 C) PVdF‐HFP‐TiO2 F‐HFP‐SiO2, C) PVd P. The results rganic filler an ]. The decrease ated to the form d thus increase ther conseque rease the ionic SEM observati membranes f rphology of pu ed composites m Results show anged the stru mbrane pores. learly observed herical particles ring synthesis, ymer network, ). This is confir dition of the in face area in com o reflected in th 21‐127 2, D) PVdF‐HFP‐SiO dF‐HFP‐TiO2, D) PV can confirm t d its uniform d of the crystalli mation of filler the role of the ence of these conductivity of ion was perfor formed in liqu ure polymer (F membranes. that the additi cture, morpho The nature of t d in Figure 2A. s build in the p other fillers al , causing a inc med by the su norganic filler le mparison to no he pore sizes. O2‐TiO2. VdF‐HFP‐SiO2‐TiO2 the formation distribution in nity degree of particles of a v polymer chains phenomena f the material [1 rmed on the cr id nitrogen. Fi Figure 2A) and ion of an inorg ology and size the structure of In situ generat polymer networ lso are incorpo rease in poros urface area stu ed to a large in on‐modified PVd 123 2. of particles of the membrane the material is very small size s interaction. A should be an 16,17]. ross‐sections of igure 2 shows modified CO2‐ ganic filler has of PVdF–HFP f the PVdF‐HFP ed SiO2 formed rk (Figure 2B). orated into the ity (Figure 2C, udy (Table 1). ncrease in BET dF‐HFP. This is f e s e A n f s ‐ s P P d . e , . T s 124 Biggest i membrane h smallest por Table 1. Compa Sample PVdF‐HFP PVdF‐HFP‐SiO2 PVdF‐HFP‐TiO2 PVdF‐HFP‐SiO2 The TG m thermal stab atmosphere) curve shape range of 17 attributed to absorbed by in agreemen curve at 171 rest of the sa It is obse TiO2 the deco °C. DTG curv This is well observed in polymer side sample can presence of observed tha place at 38 surface of Ti molecules. T the decompo [18]. Additio thermogram with peaks complete de 450‐500 °C. PVdF‐HFP ha noted from polymer ma electrolyte m thermal sta impact is obser has six times la es in compariso arison of surface a BET [m2/g] 6.14 2 38.95 2 15.65 2‐TiO2 11.75 measurements w bilities of the ) curve of the m shows prelimi 70‐190 °C. The o the removal y the sample du nt/compatible w °C for membra amples. erved at the the omposition sta ve of the samp l correlated w TG curve and i e chains. The r be due to the CF‐based PVd at the complete 6 °C with the iO2 nanoparticle The sharp exoth osition of chem on of SiO2 and m which indica on DTG curv ecomposition of There is no app as a decompos the above ana atrix slightly in medium. Only ability of PVd Ma Fi rved for SiO2, i arger specific s on to the unmod rea of membranes Average pore diameter [nm 3.44 4.62 4.77 6.45 were performe composites. T membrane is sh inary mass loss e mass loss of of moisture or uring loading [1 with the peak o ane with SiO2 an ermogram that rts at 257 °C an le shows peaks with the mass is related to th remaining 17 w e presence of dF‐HFP polyme e decomposition e correspondin es is in fact mo hermic peak at c mical absorptio SiO2‐TiO2 binar ates two weigh ves. It is also f the samples t preciable weigh ition temperatu lysis that addi nfluences therm y the addition dF‐HFP. This artyla et al. / Eur gure 3. TG/DTG a in case of whic surface area an dified PVdF‐HF s. e m] Pore volum [cm3/g] 0.07 0.04 0.02 0.01 ed to characteri G/DTG (at nit own in Figure 3 s in the temper f about 4% ca r water and so 18]. The mass l observed in the nd at 186.6 °C f t for the sample nd also occurs a s at 265 and 3 loss of the sa he degradation wt.% of the ana TiO2 or due t er matrix. It is n of the sample ng weight loss odified by the al ca. 300 °C belon n of organic sp ry system lead ht losses corre observed tha takes place bet ht loss above 5 ure of 457 °C. I tion of fillers t mal stability o n of SiO2 incr sample star ropean Journal of nalysis of pure and ch the nd the FP. me ze the trogen 3. The rature an be olvent loss is e DTG for the e with at 308 08 °C. ample of the alyzed to the s also e takes s. The lcohol ngs to pecies ds to a elated at the tween 20 °C. It was to the of the reases ts to dec mai Thi aro due bas at 7 HFP stre resp at 8 cm‐ vibr any assi red fille ben sam tion poly pre whi grai inte agg 146 stre cha der 1 is pre wat sati follo PVd upo afte of Chemistry 7 (1) d modified membr compose at 42 inly due to the s is confirmed und 480 °C. The remaining e to the presen ed PVdF‐HFP p For pure PVdF 760, 870 and 9 P phase [19]. Ba etching vibrati pectively. The 871 cm‐1 [19]. B ‐1 are related t rations of the C y band disappea In the Figure 4 igned to Si‐O‐ uces its intensi er showed peak nding vibration For all drying me shape as a p n between ino ymer leading to Using the sy cursors, the ter ich leads to th ins [20]. This eractions of gra gregation. This 64 and 1040‐11 etching and C aracters were ob rived particles [ derived from t sence of water ter confirms th ion processes owing reaction dF‐HFP did not Figure 5 illust on immersion i er 1 hour of swe ) (2016) 121‐127 ranes. 0 °C and the w e degradation o by the peak o g 20 wt. % of nce of oxides o polymer matrix F‐HFP spectra in 73 cm‐1 can be ands at 1060 an ions of CF2 a amorphous ph Bands detected a to symmetrical CF2 group. There ared when the f 4B we observe ‐Si stretching ity. A membran ks at 1658 and of ‐OH groups samples, the F pure PVdF‐HFP organic fillers o the formation ynthesis metho rmination of th he reduction of s phenomenon ains and thus d effect is visibl 130 cm− 1 (Figu C–O vibration bserved for alk [23]. Broad ban the OH group st in the investiga he occurrence o of alkoxide p ns (Equation 3‐ reveal presenc trates how the in propylene ca elling is also pre 7 weight loss of of side chains o observed in the the modified s or due to the p [18]. n Figure 4A vib assigned to cry nd 1180 cm‐1 ar and CF of the ase of PVdF‐HF at 1148, 1203, and asymmetr e was no pheno fillers were add a peak at 1100 vibration. Add ne with SiO2‐Ti d 1572 cm‐1 att [20]. FT‐IR spectra h P which confirm and oxygen n of polymer com ods of TiO2 he alkoxide grou f the contact a n leads to the decreases the d e in the absor ure 4C), corresp n, respectively koxy groups in m nd observed at a tretching and r ated samples. T of the hydrolys precursors, des 10). Analysis o e of this band ( membrane we arbonate. The esented in Figu about 50% is of the polymer. e DTG curve at samples can be presence of CF‐ brational bands ystals of PVdF‐ re attributed to e vinyl group FP is observed 1179 and 1382 rical stretching omena in which ded. 0 cm‐1, which is dition of TiO2 iO2 (Figure 4D) tributed to the have almost the ms the interact‐ atoms in the mplex [21]. from alkoxide ups may occur, among anatase e reduction of degree of their ption bands at ponding to C–H y [22]. These many alkoxide‐ about 3500 cm‐ esults from the The presence of is and conden‐ scribed by the f non‐modified (Figure 4A). eights increase solvent uptake ure 6. s . t e ‐ s ‐ o p d 2 g h s 2 ) e e ‐ e e , e f r t H e ‐ ‐ e f ‐ e d e e Fig Figure 5. PC Figure gure 4. FT‐IR spect C uptake of the com e 6. Value of absorp Martyla et al tra of pure and mo mposite polymer m ption solution afte l. / European Jou odified membranes membrane as a fun er 1 hour of swellin urnal of Chemistr s: A) PVdF‐HFP, B) ction of time. PVdF ng in PC. PVdF‐HFP ry 7 (1) (2016) 1 ) PVdF‐HFP‐SiO2, C F‐HFP, B) PVdF‐HF P, B) PVdF‐HFP‐SiO 21‐127 C) PVdF‐HFP‐TiO2, FP‐SiO2, C) PVdF‐H O2, C) PVdF‐HFP‐Ti D) PVdF‐HFP SiO2 HFP‐TiO2, D) PVdF‐ iO2, D) PVdF‐HFP‐S 125 2‐TiO2. HFP‐SiO2‐TiO2. SiO2‐TiO2. 126 Figure 7. Temp SiO2, C) PVdF‐H F Membran much better other cerami For all uptake is ac solvent upta for the hybr These value Membranes less than all o It is know electrolyte s rutile or an improvemen electrochem Figure 7 conductivity conductivity calculated fr The incr linked to the chain flexib increases, th increase, thu As temp produces em segments, m these free v promotes th transfer from perature dependen HFP‐TiO2, D) PVdF‐ igure 8. Specific co nes with SiO2 a r than the mem ic filler that wa composite po chieved quickly ake after one ho id filler up to 3 es are compa with hybrid SiO other test samp wn that additio uch as silica (S natase), magne nt of ionic co ical parameters 7 shows the te y plot of PVdF‐ y of composi om Equation 11 rease in condu e decrease in v bility of the he mobility and us improving th perature incre mpty areas tha mobile ions or volumes. The e transfer of io m one site to an Ma ncies of ionic cond ‐HFP‐SiO2‐TiO2. onductivities of me absorb solvent mbrane with n s used in exper lymer membra y (about one h our of swelling 300% for the m arable with li O2/TiO2 fillers a ples [25]. n of inorganic f SiO2), alumina ( esium oxide (M onductivity an s as well [26‐28 emperature de ‐HFP based me te gel polym 1. uctivity with t viscosity and, h electrolyte. A the dissociatio e conductivity o eases, the pol at cause specie solvated mole segmental mo onic motion by nother in the sa artyla et al. / Eur ductivities determi embranes at 25 °C t extremely we no filler or wit imental work. anes the max hour). The obs is from about membrane with terature data absorb solvent fillers to the po (Al2O3), titania MgO) results i nd mechanical 8]. ependencies of embranes. The mer electrolyt temperature ca hence, increase As the temper on rate of Li ion of the electroly ymer expands es such as po ecules to move ovement of po allowing the io me polymer ch ropean Journal of ined for the compo : A) PVdF‐HFP‐SiO ell and th any ximum served 140% h SiO2. [24]. much olymer (TiO2, in the l and f ionic ionic es is an be es the rature ns also yte. s and lymer e into lymer ons to hain or to n ioni con The syst Thi par poly com con of S the ads favo elec PVd con pro orig role dim can che crys In F bee con of Chemistry 7 (1) osite membranes p O2, B) PVdF‐HFP‐SiO neighboring po ic conductivity As commonl nductivity is not e highest condu tem has almos s is associate rticles of oxid ymer matrix mponents. This nsequently to th SiO2 to PVdF‐H ionic conduct orbed on the orable transpor The role of ctrolyte is well dF‐HFP polyme nductivity. Ion omoted by hi ginated from it e of TiO2, it i minish the ion‐a n be originated emical interactio In case of the stallinity and lo Figure 8 logarit en plotted as nductivities at 2 ) (2016) 121‐127 presented as Arhen O2‐TiO2.C) PVdF‐H olymer chain r [19]. y found in t a linear functi uctivity is for Si t the same con ed with struc e fillers are and can inte s leads to th he conductivity HFP based gel p tivity [27] whe filler grain sur rt environment TiO2 nanopar l known. Addit er matrix leads n migration w gher TiO2 di ts inherent dip is well agreed aggregation in t d from their d on between TiO e membrane co ow surface area thms of specific a function of 25 °C range fro 7 nius plots: PVdF‐H HFP‐TiO2, D) PVdF‐ resulting in the composite m ion of the filler iO2 and TiO2 fil nductivity as p cture of mem uniformly disp eract with oth he fast ion t y enhancement polymer electro en thin layers rfaces can ensu t for ions. rticles in the tion of TiO2 na to improveme within nanopo electric consta ole property. R d that the nan the electrolyte m dielectric natu O2 and PVdF‐HF ontaining SiO2‐ a lead to low con c conductivities f filler content om 1.0×10−3 S/ HFP, B) PVdF‐HFP‐ HFP. e increment of materials, the concentration. ller. The binary ure PVdF‐HFP. mbranes, when persed in the her electrolyte transport and t. The addition olyte increases of electrolyte ure sufficiently e gel polymer anoparticles to ent of the ionic ores may be ant, which is Relating to this noparticles can medium, which re and/or the FP [29,30]. ‐TiO2 filler low nductivity [21]. s at 25 °C have t. The specific cm to 3.5×10−3 ‐ f e . y . n e e d n s e y r o c e s s n h e w . e c 3 Martyla et al. / European Journal of Chemistry 7 (1) (2016) 121‐127 127 S/cm. The highest conductivity was obtained for SiO2, TiO2, a binary filler, lead to lower conductivity than pure PVdF‐HFP. It is the further evidence of the effect of the membrane structure on its conductivity and this is consistent confirmed by other observations [29,30]. Apparent activation energies have been calculated as a function of filler using the Equation 12. Activation energies and room temperature ionic conductivities of different compositions are summarized in Table 2. The minimum value of Ea is 0.141eV for the pure PVdF‐HFP. There is an increase in Ea for a binary filler content. For the rest of samples the values of Ea indicate no clear correlation with the filler content. Table 2. The value of activation energy as a function of filler. Composite polymer gel electrolyte Ea [eV] PVdF‐HFP 0.141 PVdF‐HFP‐ SiO2 0.142 PVdF‐HFP‐TiO2 0.145 PVdF‐HFP‐SiO2‐TiO2 0.175 4. Conclusions 1. Composite polymer‐ceramic membranes were manu‐ factured by a modified Bellcore procedure, in which dibuthyl phthalate extraction and in‐situ oxide formation were carried out simultaneously by the application of supercritical CO2. 2. XRD patterns reveal that crystallinity degree of the polymer matrices decreased in the presence of fillers. 3. Gel electrolytes obtained by activation of dry memb‐ ranes exhibit significant, very high conductivities, exceeding 6.06×10‐3 Scm‐1. 4. The obtained results are promising from the point of view of possible application in Li‐ion battery techno‐ logies. PVdF‐HFP‐SiO2, ‐TiO2 and ‐SiO2‐TiO2 composite polymer electrolyte membranes were studied to investigate the role CPD and inorganic filler on gel polymer electrolyte. The addition of filler to PVdF‐HFP based gel polymer electrolyte allowed to obtain microporous composite polymer electrolyte. Reduction in crystallinity after addition of inorganic fillers and their interaction with the polymer were established from the XRD and FTIR studies. Incorporation of fine, well dispersed oxide particles prevents reorganization of polymer chains, resulting in a decrease in polymer crystallinity which gives rise to an increase in ionic conductivity and stabile thermal properties [30]. Morphological study shows porous structure of the membranes after supercritical CO2 drying that can be more effective change of the structure and ionic conductivity. Based on these results it can be suggested that silica obtained by one‐pot reaction method can be good filler for polymer electrolyte membranes. 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