Influence of N-doped TiO2 on lithium ion conductivity of porous polymeric electrolyte membrane containing LiClO4

Solid State Ionics 212 (2012) 18–25

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Influence of N-doped TiO2 on lithium ion conductivity of porous polymeric electrolyte membrane containing LiClO4 Ki-Seok Kim a, Soo-Jin Park a, b,⁎ a b

Department of Chemistry, Inha University, Incheon 402-751, Republic of Korea Korea CCS R&D Center, Korea Institute of Energy Research, 152 Gajeongro, Yuseoung-gu, Daejeon 305-343, Republic of Korea

a r t i c l e

i n f o

Article history: Received 28 September 2011 Received in revised form 13 February 2012 Accepted 14 February 2012 Available online 9 March 2012 Keywords: PVDF-HFP TiO2 N-doping Polymer electrolyte Ion conductivity

a b s t r a c t Polymer electrolyte composites (PECs) based on poly(vinylidenefluoride-co-hexafluoropropylene) (PVDFHFP) that contained lithium perchlorate (LiClO4) were prepared. N-doped TiO2 (N-Ti) was used as a filler to enhance the Li+ ion conductivity. Structural modification and electrochemical properties of the PECs were investigated in order to understand the effect of N–Ti in the polymer matrix on ionic conductivity. The PECs prepared with different N–Ti contents were characterized by SEM, mapping, DSC, and a.c. impedance. From the results, the prepared N–Ti-filled PECs showed a good ability to absorb and retain the Li + ion. The presence of N–Ti in PVDF-HFP was resulted in increasing the ion conductivity significantly. The highest ion conductivity (6.7 × 10− 4 S/cm) was observed at a relative low N–Ti content (2.5 wt.%). It is believed that the addition of N–Ti could provide the foundation for the ionic transportation and ion-mobility of the PECs and the N-doping on TiO2 led to a more enhanced ionic conductivity by higher interaction between polar groups, such as O and N of N–Ti and Li+ ions. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Recently, portable electric devices, such as mobile telephones, movable computers, and hybrid electric vehicles have been remarkably developed with many advantages, leading to a strong need for safe and high energy lithium ion batteries. Lithium-ion polymer batteries have attracted a great deal of attention due to its higher energy density, improved safety hazards, and good processability. For polymer electrolytes, several polymers have been investigated and developed, including poly(ethylene oxide) (PEO), poly(acrylonitrile) (PAN), poly(methyl methacrylate) (PMMA), poly(vinyl chloride) (PVC), poly(vinylidene fluoride) (PVDF), and poly(vinylidene fluoride-hexafluoropropylene) (PVDF-HFP) [1–3]. Among them, PVDF-HFP has been considerably researched, in which the amorphous HFP phase helps in capturing a large amount of liquid electrolytes, while the crystalline PVDF acts as a mechanical support for the polymer matrix due to their semi-crystalline nature and high dielectric constant [4–6]. Organic/inorganic composites constitute a remarkable family of isotropic, flexible, and amorphous nanocomposites, which have been investigated extensively for structural materials and biomedical materials. In addition, numerous organic/inorganic composite membranes have fascinated researchers in the fuel cell field [7,8]. High ionic conductivity, adequate chemical and mechanical strength, ⁎ Corresponding author at: Department of Chemistry, Inha University, Incheon 402751, Republic of Korea. Tel./fax: + 82 32 860 8438. E-mail address: [email protected] (S.-J. Park). 0167-2738/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2012.02.024

extended thermal stability, and low price are the favorable characteristics of polymer electrolyte membranes. These characteristics can be obtained by synergistically combining both organic and inorganic materials together. The prepared hybrid materials possess enhanced properties compared to the single organic or inorganic materials. The organic/inorganic composite membranes have exhibited a specific interaction between the components, which has influenced the electrochemical stability [9,10]. The deficiencies of PVDF-HFP based polymer electrolyte membranes can be overcome through the addition of fillers, such as ceramic, silica, and titanium oxide. The fillers reduce the crystalline character, as well as promote the thermal behaviors of the membrane. Among the several fillers in polymer electrolyte membranes, nano-sized titanium oxide (TiO2) supports the ionic mobility due to its substrate characteristics, such as shape and surface nature, which effectively disturbs the order packing tendency of the host polymer chains [11,12]. Recently, the particle characteristics of the fillers are found to have a tremendous influence on the electrochemical properties of polymer electrolyte membranes. Indeed, in lithium battery studies, nano-sized fillers have exhibited higher ionic conductivity than the micro-sized particles [13]. In addition, the doping of metal, nitrogen, carbon, sulfur or the use of both nitrogen and sulfur as a dopant has been investigated in order to improve the electrical properties of the substrate. According to the theoretical studies, the doping of anions, such as nitrogen, carbon, or sulfur by substituting them for the lattice oxygen of titanium dioxide has been found to be effective in narrowing the band gap of titanium dioxide and not making any changes in the conducting band level [14–16].

K.-S. Kim, S.-J. Park / Solid State Ionics 212 (2012) 18–25

As mentioned above, many researches on the polymer electrolyte containing various fillers have been reported for a long time, however, there has been limited work on the effect of heteroatom-doped TiO2 on the electrochemical properties of polymer electrolytes. Therefore, in this work, nitrogen-doped TiO2 (N–Ti) was prepared by using the urea mixing and calcinations at 500 °C. The N–Ti/poly(vinylidene fluoride-hexafluoropropylene) (PVDF-HFP) membrane prepared by film casting was investigated as the polymer electrolyte. The effect of N–Ti on the electrochemical and physicochemical properties of the N–Ti/PVDF-HFP is discussed.

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Table 1 Composition of polymer electrolytes. Sample

PVDFHFP

0 N–Ti/PVDF-HFP 2.5 N–Ti/PVDF-HFP 5.0 N–Ti/PVDF-HFP 7.5 N–Ti/PVDF-HFP 10 N–Ti/PVDF-HFP

Plasticizer (1:1, w/w)

30 30 30 30 30

EC

DEC

32.5 31.25 30.0 28.75 27.50

32.5 31.25 30.0 28.75 27.50

LiClO4

TiO2

5 5 5 5 5

0 2.5 5.0 7.5 10

2. Experimental 2.1. Materials Poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP) (MW: ~ 400,000, Tm: 140–145 °C) and titanium tetrachloride (TiCl4) were obtained from Aldrich. Acetone, lithium perchlorate (LiClO4), and diethyl carbonate (DEC) were procured from Aldrich and used without further purification. Ethylene carbonate (EC) and urea were supplied from TCI and Aldrich, respectively. 2.2. Preparation of N–Ti An amount of TiCl4 was added into distilled water and vigorously stirred for 1 h. Then, the hydrolysis of TiCl4 in the water was followed as [17]: þ

−

TiCl4 þ H2 O→TiO2 þ 4H þ 4Cl :

ð1Þ

The hydrolysis reaction created a colloid solution with several nano-sized TiO2 particles. Finally, ammonia solution was added to the colloidal solution in order to precipitate TiO2. The resultant was then washed with distilled water and dried at 100 °C. For the N doping, the obtained TiO2 and urea were mixed with a molar ratio of 1:3 for a few minutes using mortar. For the solid phase reaction, the mixture was calcined at 500 °C for 2 h under N2 gas. The temperature rose to 500 °C by 5 °C/min with an N2 flow rate of 200 ml/min. The prepared nitrogen doped TiO2 was named as N–Ti.

equipped with an Mg Kα (1253.6 eV) X-ray source, and a highperformance multichannel detector that was operated at 200 W. The textural characteristics of TiO2 and N–Ti were analyzed at 77 K by gas adsorption analysis (BELSORP, BEL Japan). The samples were degassed at 423 K for 20 h to obtain a residual pressure of b10 − 6 mm Hg. The specific surface areas and pore volume of the samples were determined using the Brunauer–Emmett–Teller (BET) equation and Dubinine–Radushkevitch (DeR) equation, respectively. The amount of N2 adsorbed at a relative pressure (P/P0) of 0.98 was used to examine the total pore volumes, which are a combination of the micropore and mesopore volumes. The morphologies of pure PVDF-HFP and N–Ti/PVDF-HFP films were observed by a scanning electron microscopy (SEM, S-4200, Hitachi). The thermal properties, such as melting temperature, ΔHm, and crystallinity of the N–Ti/PVDF-HFP films with different N–Ti contents were determined in the range of 90–220 °C by using Differential Scanning Calorimetry (DSC, DSC200F3/NETZSCH). The thermal stability of the N–Ti/PVDF-HFP films with different N–Ti contents was measured using thermogravimetric analyses (TGA, Du-Pont TGA-2950 analyzer) from 30 to 700 °C at a heating rate of 10 °C/min in a nitrogen atmosphere. Impedance spectroscopy was used to determine the ionic conductivity of the PECs. The measurements were carried out in the frequency range from 100 kHz to 10 Hz. The PECs (1 × 1 cm) were sandwiched between two polished stainless-steel electrodes. The conductivity values (σ) were calculated from the bulk resistance (Rb), which was determined by equivalent circuit analysis software

2.3. Preparation of N–Ti/PVDF-HFP films

1 Rb A ¼ σ t

The PECs were prepared using a solution casting method. The PVDF-HFP was dried at 100 °C under a vacuum for 24 h in order to eliminate all of the attached water molecules. LiClO4 was used as the lithium ion contributor in PECs after being dried at 120 °C under a vacuum for 24 h. PVDF-HFP was dissolved in acetone containing 5 wt.% propyl alcohol for 1 h by stirring at 60 °C. Then, the different contents (0 wt.% to 10 wt.%) of N–Ti were added to the PVDF-HFP solution by vigorously stirring the solution for 3 h. LiClO4, EC, and DEC were added to the mixture solution and the resultant mixture was stirred for 24 h at 50 °C until a homogeneous solution was formed. The mixture solution was poured onto a glass plate and evaporated slowly at room temperature in a vacuum oven. The film thickness prepared was about 50 μm and the composition of the film preparation is presented in Table 1.

where t is the film's thickness and A is the film's area.

2.4. Characterization N–Ti was confirmed by the X-ray diffraction (XRD, Rigaku D/Max 2200 V) at 40 kV and 40 mA by using Cu Kα radiation. The XRD patterns were obtained in 2θ ranges between 5° and 70° at a scanning rate of 2°/min. The functional groups and content of the nitrogen group of N–Ti were examined by a X-ray photoelectron spectroscopy (XPS, K-Alpha) that used a VG Scientific ESCALAB MK-II spectrometer

ð2Þ

3. Results and discussion 3.1. Characterization of N–Ti For the N doping of TiO2, the mixture of TiO2 and urea was calcined at 500 °C for 2 h. Fig. 1 shows the XRD patterns of pristine TiO2 and N–Ti. From the XRD patterns, the synthesized TiO2 shows an anatase phase while the N doping of TiO2 seems to not have caused any changes in the crystalline structure of TiO2. However, in the case of N–Ti, the intensity of the XRD patterns slightly decreased Table 2 Textural properties of TiO2 and N–Ti. Sample

SBETa (m2/g)

VTb (cm3/g)

dMc (nm)

VMed (cm3/g)

TiO2 N–Ti

46.65 10.90

0.142 0.060

12.18 22.18

0.142 0.060

a b c d

Specific surface area. Total pore volume. Mean pore diameter. Mesopore volume.

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(a) TiO2

Intensity (a. u.)

(b) N-Ti

(b)

(a) 0

10

20

30

40

50

60

70

2θ Fig. 1. XRD patterns of pristine TiO2 and N–Ti.

compared to pristine TiO2, indicating the fact that the TiO2 was substituted by the N group after the calcinations of the mixture of urea/TiO2 by the solid reaction process. Also, it is suggested that the decomposition of urea in the mixture was able to restrain the formation and growth of the TiO2 crystal phase transformation of the amorphous sample [18]. The structural change of TiO2 after N-doping is further determined by N2 adsorption analysis as shown in Fig. 2 and Table 2. It is clear that all samples show type III curves with mesoporous features and the specific surface area of TiO2 decreased with decreasing total pore volume after N-doping, indicating the structural change by solid phase reaction between TiO2 and urea. In addition, average pore size of

Sample

Ti

O

N

TiO2 N–Ti

29.14 15.12

61.63 32.17

0.49 36.03

TiO2 increased and mesopore volume is opposite due to the blocking effect of TiO2's pores by the N-doping. It is expected that lower porous feature of the N–Ti leads to the more easy and direct Li ion transfer in TiO2 networks, as a result of decreasing the Li ion adsorption into the TiO2 pores. The N-doping of the pristine TiO2 was determined by using XPS spectroscopy. Fig. 3 shows the XPS spectra of pristine TiO2 and N–Ti. Pristine TiO2 consists only of oxygen (530 eV) and titanium (458 eV) and no peaks are observed in the pristine TiO2 from the detecting ranges, whereas N–Ti had an N1s peak at 398 eV, indicating the chemisorbed nitrogen and substitutional nitrogen. The N content of N–Ti was 36 wt.%, as shown in Table 3. 3.2. Morphologies of N–Ti/PVDF-HFP PECs The morphologies of the pure PVDF-HFP and N–Ti/PVDF-HFP films were observed using SEM, as shown in Fig. 4. The pure PVDF-HFP film (thickness: about 50 μm (inset of Fig. 4a)) showed a smooth surface and homogeneous phase, whereas the addition of N–Ti caused changes in the morphologies of the PVDF-HFP films, indicating structural changes in the composites. The PVDF-HFP films that contained N–Ti gained a rough surface while the porous structure and the

TiO2

100

TiO2

O1S

N-Ti

N1S

N-Ti

80

Ti2P

Counts/s

Volume adsorbed (cm3/g)

Table 3 Elemental composition of TiO2 and N–Ti (unit: at.%).

60 40 20 0 0.0

0.2

0.4

0.6

0.8

1000

1.0

800

P/P0 0.15

400

200

0

(a) TiO2

N1S

30000

TiO2

(b) N-Ti

N-Ti

25000 0.10

Counts/s

Differential pore volume (cm3/g)

600

Binding energy (eV)

0.05

20000 15000 10000 5000

(b) (a)

0.00 0 1

10

100

Pore width (nm) Fig. 2. (a) N2 adsorption/desorption isotherm and (b) mesopore distribution of TiO2 and N–Ti.

395

400

405

410

Binding energy (eV) Fig. 3. XPS wide scan of (a) pristine TiO2 and N–Ti and (b) N1S peak of pristine TiO2 and N–Ti.

K.-S. Kim, S.-J. Park / Solid State Ionics 212 (2012) 18–25

porosity of the films increased with an increase in the N–Ti content. At a low N–Ti content (2.5 wt.%), the surface of the PVDF-HFP film became rough with lower porosity, indicating a high dispersion of N–Ti.

21

At a high N–Ti content (over 2.5 wt.%), however, the porosity of the PVDF-HFP film significantly increased with an increase in the N–Ti, which was attributed to the growth of the aggregation of N–Ti and

Fig. 4. SEM images of the N–Ti/PVDF-HFP films with different N–Ti contents; (a, b) 0 wt.% (inset: film thickness is about 50 μm), (c, d) 2.5 wt.%, (e, f) 5.0 wt.%, and (g, h) 10 wt.%.

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high solvent retention ability of the PVDF-HFP, indicating that the pores in microstructure occur due to slow solvent removal [19]. It is believed that the aggregation of the N–Ti decreased the ionic conductivity of the N–Ti/PVDF-HFP films, resulting from the decrease in ion mobility [20]. In addition, the internal phase of the films is observed by the cross-section of the films, as shown in Fig. 5. As mentioned above, the internal phase of pure PVDF-HFP film is smooth, whereas N–Ti/PVDF-HFP films show the increase in the formation of the pore and aggregates between N–Ti with increasing the N–Ti. The dispersion of N–Ti in the polymer matrix is further examined using the mapping method. The Ti and nitrogen mapping (Fig. 6a–d) of the 2.5 N–Ti/PVDF-HFP films showed a homogeneous distribution of the respective elements, whereas the 5.0 and 7.5 N–Ti/PVDF-HFP films exhibit a heterogeneous dispersion of Ti due to localized aggregation of Ti.

3.3. Thermal properties of N–Ti/PVDF-HFP PECs To confirm the effect of the addition of fillers on the crystalline structure of the polymeric composite electrolytes, a crystallinity value was obtained by using the DSC method, as shown in Fig. 7. Compared to pure PVDF-HFP, PCE PVDF-HFP reveals the broad peak around melting point due to the addition of EC and DEC as plasticizer. Furthermore, the intensities of the characteristic peaks are decreased and broaden with the addition of N–Ti to the PVDF-HFP/LiClO4 system, indicating the crystalline-to-amorphous transformation of PVDF-HFP matrix. N–Ti fillers can decrease the crystallinity of PVDFHFP due to the following reasons: 1) Lewis acid–base interaction between the fluorine of PVDF-HFP and the Lewis acid sites on the surface of N–Ti, as in the cases of SiO2 and Al2O3, and between Li + and nitrogen groups and 2) the prevention of recrystallization of the PVDF-HFP chains by the N–Ti intercalated into polymer chains [21].

The value of the crystallinity (xc) of the N–Ti/PVDF-HFP films has been defined as the enthalpy ratio of PVDF-HFP with different N–Ti contents. It can be calculated with the equation:

xc ¼

ΔH f

∘

ΔHf

ð3Þ

where ΔHf° is the melting enthalpy (150 J/g) of a completely crystalline PVDF-HFP and ΔHf is the experimental enthalphy [22]. The crystallinity value can describe any relative changes in the crystalline or amorphous phase of composite electrolytes. The melting temperature, melting enthalpy (ΔHm), and crystallinity (xc) values are summarized in Table 4. It is clear that the ΔHm of pure PVDF-HFP is significantly decreased with the addition of plasticizer (about 60 wt.%). In addition, the ΔHm and xc values became decreased with an increase of the N–Ti content. This indicates that the crystalline domain of PVDF-HFP decreased by adding plasticizer and N–Ti lead to more homogeneous PEC system, resulting in the enhanced ion conductivity by easy mobility of lithium ion. Fig. 8 shows the TGA thermogram of the PEC films prepared with different N–Ti contents. In all of the samples, a slight weight loss occurred from about 70 to 100 °C, which was attributed to the evaporation of the residual solvent or moisture in the samples. All of the samples were stable up to about 250 °C. It could be seen that a decomposition range from about 250 to 500 °C is due to the decomposition of PVDF-HFP and LiClO4. The decomposition temperature of each sample was decreased with increasing N–Ti content, relating the decrease of polymer ratio as the increase of N–Ti content. In the range from 500 to 700 °C, the final weight loss was proportional to the incorporated N–Ti content. These results clearly indicate that N–Ti loaded polymer electrolyte has a good thermal stability up to 250 °C. In addition, this PEC system has a higher thermal stability compared to

Fig. 5. Cross-section images of the N–Ti/PVDF-HFP films with different N–Ti contents; (a) 0 wt.%, (b) 2.5 wt.%, (c) 5.0 wt.%, and (d) 10 wt.%.

K.-S. Kim, S.-J. Park / Solid State Ionics 212 (2012) 18–25

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Fig. 6. (a, b) Ti mapping and (c, d) nitrogen mapping of 2.5 N–Ti/PVDF-HFP, (e) Ti mapping of 5.0 N–Ti/PVDF-HFP, and (f) Ti mapping of 7.5 N–Ti/PVDF-HFP.

general liquid electrolytes (about 80 °C) used in lithium ion battery [23].

Heat flow (mV)

3.4. Electrochemical properties of N–Ti/PVDF-HFP PECs Electrochemical impedance analysis is a powerful technique for determining the electrical properties of interfaces between Table 4 Parameters from the DSC curves for the PVDF-HFP membranes as a function of N–Ti content.

Pure PVDF-HFP PCE PVDF-HFP 2.5 N-Ti/PVDF-HFP 5.0 N-Ti/PVDF-HFP 7.5 N-Ti/PVDF-HFP 10 N-Ti/PVDF-HFP

100

120

140

160

180

200

220

Temperature (oC) Fig. 7. DSC curve of the N–Ti/PVDF-HFP films with different N–Ti contents.

Sample

Tm ( °C)

ΔHm

xc (%)

PVDF-HFP 0 N–Ti/PVDF-HFP 2.5 N–Ti/PVDF-HFP 5.0 N–Ti/PVDF-HFP 7.5 N–Ti/PVDF-HFP 10 N–Ti/PVDF-HFP

143.8 140.0 142.6 142.6 142.4 141.5

105 37.1 31.3 23.9 22.9 13.3

35.3 29.8 22.8 21.8 12.7

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80

60

40 N-Ti residue

20 100

200

300

400

500

N-Ti/PVDF-HFP

6.0x10-4 5.0x10-4 4.0x10-4 3.0x10-4 2.0x10-4

Decomposition temp.

0

TiO2/PVDF-HFP

7.0x10-4

Ion conductivity (S/cm)

Weight loss (%)

8.0x10-4

PVDF-HFP 2.5 N-Ti/PVDF-HFP 5.0 N-Ti/PVDF-HFP 7.5 N-Ti/PVDF-HFP 10 N-Ti/PVDF-HFP

100

600

1.0x10-4

700

0

Temperature (oC)

2

4

6

8

Fig. 8. TGA thermogram of the N–Ti/PVDF-HFP films with different N–Ti contents.

electronically and ionically conducting phases. It may be used to investigate the dynamics of bound or mobile charges in the bulk or interfacial regions of a variety of solid and liquid materials: ionic, semiconducting, mixed electronic, ionic, and even insulators (dielectrics). In order to enhance the Li ion conductivity, the effect of TiO2 and N–Ti as filler on ion conducting behaviors of PECs systems was investigated. The impedance plots of the PECs prepared with different TiO2 and N–Ti contents are presented in Fig. 9. The approximate bulkresistance value could be obtained by referring to the touch point at the x-axis. It can be seen that the ionic resistance of the PECs decreased with increasing the filler content. Compared to TiO2/PVDF-

Fig. 10. Ionic conductivity of the TiO2/PVDF-HFP and N-Ti/PVDF-HFP films with different filler contents.

HFP films, N–Ti/PVDF-HFP films have a lower ionic resistance and the highest ionic conductive property was observed at 2.5 wt.% N– Ti, indicating that the ion mobility of the PECs is increased by the strong interaction between polar groups, i.e., O and N, of N–Ti and lithium ions. Fig. 10 shows the ion conductivity versus N–Ti contents for the LiClO4-loaded PECs at room temperature. A rapid increase in the ion conductivity was observed by adding a small quantity of the N–Ti. The maximum ionic conductivity (6.7 × 10 − 4 S/cm) was achieved at 2.5 wt.% N–Ti content. This value was about four-orders-of-magnitude higher than that of the pure LiClO4/PVDF-HFP and this value 500

6000

PVDF-HFP 2.5 TiO2/PVDF-HFP

a

b

5.0 TiO2/PVDF-HFP 7.5 TiO2/PVDF-HFP 10 TiO2/PVDF-HFP

-Z'' (Ohm)

-Z'' (Ohm)

4000

PVDF-HFP 2.5 TiO2/PVDF-HFP

2000

5.0 TiO2/PVDF-HFP 7.5 TiO2/PVDF-HFP 10 TiO2/PVDF-HFP

0

0 0

2000

4000

6000

8000

0

200

Z' (Ohm)

400

600

Z' (Ohm)

6000

500 PVDF-HFP 2.5 wt.% N-Ti 5.0 wt.% N-Ti 7.5 wt.% N-Ti 10 wt.% N-Ti

c

d

-Z'' (Ohm)

-Z'' (Ohm)

4000

2000

PVDF-HFP 2.5 N-Ti/PVDF-HFP 5.0 N-Ti/PVDF-HFP 7.5 N-Ti/PVDF-HFP 10 N-Ti/PVDF-HFP

0

0 0

2000

4000

Z' (Ohm)

10

N-Ti content (wt.%)

6000

8000

200

400

600

Z' (Ohm)

Fig. 9. Electrochemical impedance spectra of (a, b) TiO2/PVDF-HFP and (c, d) N–Ti/PVDF-HFP films with different filler contents.

K.-S. Kim, S.-J. Park / Solid State Ionics 212 (2012) 18–25

was higher than the TiO2/PVDF-HFP and previous results [24]. In the previous study, Aravindan et al. [20] reported that the highest ionconductivity of LiDFOB-loaded PVDF-HFP electrolyte was observed at 5.0 wt.% TiO2. This result indicated that as mentioned in the Introduction, N-doping of TiO2 leads to the enhanced electrochemical properties compared to pure TiO2 due to the decreased band gap and homogeneous dispersion at low content. However, when the N–Ti content was increased over 2.5 wt.%, the ion conductivity decreased remarkably from the maximum value, which has been attributed to the blocking effect on the transporting of charge carriers by the aggregation of N–Ti [25]. It can be concluded that the addition of an optimum N–Ti content can provide the most suitable environment for ionic transportation and achieve the highest conductivity. The enhanced ionic conductivity of the PECs is mainly dependent on polar groups increased by the nitrogen doping onto TiO2.

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]

4. Conclusions [14]

In this work, nitrogen doped TiO2 (N–Ti)/poly(vinylidenefluorideco-hexafluoropropylene) (PVDF-HFP) that contained lithium perchlorate (LiClO4) was prepared for polymer electrolyte composites (PECs). The electrochemical properties of PECs were investigated as a function of N–Ti content. It was found that the morphology changed with the addition of N–Ti and the crystallinity of the PVDF-HFP decreased. In addition, the ion conductivity of pure PVDF-HFP significantly increased with the addition of N–Ti. The highest ion conductivity (6.7 × 10 − 4 S/cm) was observed at a relatively low N– Ti content (2.5 wt.%). It is believed that the enhanced ionic conductivity of PECs was dependent on the reduced crystallinity and enhanced ion mobility by nitrogen doping onto conventional TiO2. Acknowledgement We acknowledge the financial support by grants from Korea CCS R&D Center, funded by the Ministry of Education, Science and Technology of Korean government.

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