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Ionic conductivity properties of amorphous Li–La–Zr–O solid electrolyte for thin film batteries
Solid State Ionics 229 (2012) 14–19
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Ionic conductivity properties of amorphous Li–La–Zr–O solid electrolyte for thin film batteries D.J. Kalita a, S.H. Lee a, K.S. Lee a, D.H. Ko b, Y.S. Yoon a,⁎ a b
Department of Materials Science and Engineering, Yonsei University, Shinchon Dong 134, Seoul 120-749, Republic of Korea Korea Aerospace Research Institute/Payload Optics Team, 115 Gwahangno, Yuseong-gu, Daejeon 305-333, Republic of Korea
a r t i c l e
i n f o
Article history: Received 8 June 2012 Received in revised form 14 September 2012 Accepted 14 September 2012 Available online 17 November 2012 Keywords: Thin film battery Solid electrolyte RF-power Li-modifier
a b s t r a c t Lithium–lanthanum–zirconium–oxide (Li–La–Zr–O) thin films were deposited using radio-frequency (RF) magnetron sputtering by varying the RF power from 40 to 80 W. The characteristics are investigated with X-ray diffraction analysis (XRD) and Field Emission Scanning Electron Microscope (FE-SEM). Results showed that the thin films had amorphous structure accompanied with a smooth, dense, and homogeneous surface. Impedance analysis was conducted to investigate the ionic conductivity. The thin film deposited at 40 W gives the highest ionic conductivity of 4×10−7 S/cm when compared to both the 60 W and 80 W samples which show the ionic conductivity of 2×10−7 S/cm and 0.8×10−7 S/cm respectively. The temperature dependence of the ionic conductivities fit the Arrhenius relation and the thin film deposited at 40 W possessed the lowest activation energy. These properties make Li–La–Zr–O thin films as a promising candidate material for use in all-solid-state thin film lithium ion batteries. © 2012 Elsevier B.V. All rights reserved.
1. Introduction The need of microenergy sources with high power density is inevitable for the miniaturization of electronic devices like implantable medical devices, electronic paper, RFID-Tag, microelectromechanical systems (MEMS), etc. [1–4]. Although most of the modern battery technologies are able to partially satisfy this requirement, thin film lithium ion batteries (LIBs) are considered as the most promising candidate for their ability to provide the maximum voltage and energy storage densities [5,6] in their microstate. Electrolyte which is one of the major components in LIBs may be found either in solid-state or in liquid-state. In comparison with the liquid-state electrolytes, solid-state electrolytes have not provided satisfactory performance to meet the requirements of practical field applications that required high ionic conductivity, and chemical and electrochemical stability [7]. However, a wider range of potential field applications are provided by the solid state electrolytes than the liquid-state electrolytes. Liquid-state electrolytes are accompanied by the inherent disadvantages such as short lifetime, leakage of electrolyte and nonfunctioning of device below the freezing point and above the boiling point [8]. The fabrication of thin film forms of energy storage devices is possible with the solid state electrolytes. Recently, a number of solid electrolytes having perovskite structure such as Li0.35La0.55TiO3 (LLT) [9–11] and with NASICON structure such as LiTi2(PO4)3 (LTP) [12,13] have been reported with a high bulk ionic conductivity in the range of ~10−3 to 10−4 S/cm, respectively. Although these materials were accepted for LIBs, recent investigation reveals that anode materials ⁎ Corresponding author. Tel.: +82 2 2123 2847; fax: +82 2 2123 8145. E-mail address: [email protected] (Y.S. Yoon). 0167-2738/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.ssi.2012.09.011
having red-ox potential lower than 1.8 V vs. Li/Li+ cannot be employed because of the red-ox reaction of titanium with lithium that causes reduction of Ti4+ to Ti3+ in bulk [14], which restricted the energy and power densities of LIBs to a far lower value. During the last several years, a series of compounds with a nominal composition of Li5La3M2O12 (M=Nb, Ta) and garnet like structure has been investigated by Thangadurai et al. [15]. However, a better chemical stability against lithium anode and a relatively high ionic conductivity have been obtained for Li7La3Zr2O12 (LLZ) by Murugan et al. [16]. Both the phase property and fabrication processes of solid electrolytes have crucial roles in their electrochemical properties. Earlier investigations make it obvious that solid electrolyte with a garnet like or crystalline structure gives an ionic conductivity up to 10−4 S cm−1, preparation of which usually involves a high temperature step like conventional high temperature annealing, rapid thermal annealing (RTA) and deposition with a high temperature substrate heating, etc. All these high temperature processes have adverse effects when the fabrication involves thin films approximately 1 μm owing to the current demand for thin film LIBs. A crucial matter that remains as blur and muddling to the scientist is to optimize certain properties of the coating substrate composite such as hardness, evolution of residual stress, fracture roughness and other mechanical properties. The treatment of thin films at elevated temperature around 700 °C results in being prone to failure [17]. In most of the PVD coatings like sputtering, laser deposition, etc., this failure results from the development of stresses and interlayer inter-diffusion resulting to severe damages such as film cracking, de-cohesion by buckling and inter-phase [18–20]. A deposition annealing process always induces thermal stress in the thin films which may result to the formation of micro-cracks even at moderate temperature like 400 °C [19]. A number of studies
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have been done to tackle these kinds of serious problems by using adhesive layers such as Cr, TiN. However these multilayer interconnects are also influenced by the development of thermal stress during heat treatment resulting in void formation and crack propagation [18]. Differences or mismatches in the thermal and mechanical properties causing large thermal stresses give rise to these menacing problems during annealing of thin films [21]. Employment of high temperature (>600 °C) during crystallization of thin films leads to a severe loss of Li content and may cause some unwanted phase separation. Investigation by Ahn et al. inferred that crystallizations of thin LTO thin films cause segregation of TiO2 phase incorporating holes at grain boundary for which the microstructures experience electric short cuts [22]. This phase separation has become one of the concerning factors preventing uses of dopants like Al to a certain extent, at high temperature. Chiang et al. [23] and Ohjuku et al. [24] had reported the use of Al substitution to accomplish high temperature stability to the Mn and Ni based oxides. However, later investigation by S. Buta et al. [25] showed that Al possesses a lower miscibility with a number of transition metal oxides. This results in the formation of LiAlO2 phase as it has positive mixing enthalpies with lithium transition-metal oxides. C. Wolverton's [26] study revealed that Al also has a high tendency to form precipitate phase in multicomponent systems containing transition metal. These kinds of drawbacks always make high temperature thin film processes accompanied with failure of experiments. Because of such constraints, amorphous electrolytes are broadly investigated for having resultant denser microstructure with negligible grain boundary effect, and improved chemical stability for applying in all-solid-state LIBs [27], although they result in moderate ionic conductivity compared to a crystalline one. Amorphous electrolytes which permit only the migration of lithium ion can be categorized either as sulfide [4] or an oxide series by element [28,31]. In comparison to oxide series sulfide series exhibit high ionic conductivity and thermal stability but in practical purposes it is hard to use them because of their high instability when they come into contact with moisture in the air. On the other hand, oxide series electrolytes are chemically and electrochemically more stable than sulfide series solid-state electrolytes at normal atmospheric conditions [32,33]. Therefore, solid-state oxide electrolytes are more potent for TFBs and TFSCs. In this study, we focused on the room temperature fabrication and characterization of lithium– lanthanum–zirconium–oxide (Li–La–Zr–O) thin films having amorphous structure by RF sputtering. Our aptness towards this material's amorphous synthesis lies on both the pros and cons associated with it. The prime interest about this material is its possession of high stability against lithium anode and also its high bulk ionic conductivity resulting it to show up against several solid electrolytes like LATP and LATO [34]. Both these features are the prime requirement for a solid electrolyte in thin film Li ion battery. Investigation by Geiger et al. [35] uncovered that Li–La–Zr–O results in a stable cubic phase with bulk ionic conductivity of 10−4 S/cm in the presence of Al as a phase stabilizer. However, the temperature requirement for this process is above 900 °C which results in severe damages to the thin film such as cracking and the total failure of the system. Again, Xie et al. [36] reported that processing of this material at around 850 °C gives rise to a tetragonal phase and ionic conductivity of 10−6 S/cm making the process unprofitable. A room temperature synthetic process of this material is more justifiable if it leads to a sufficient ionic conductivity (10−6 S/cm) expelling the cons related to high temperature processes. As compared with other fabrication methods radio frequency (RF) magnetron sputtering is considered to be an easy, established and environmentally benign process of deposition which can be easily controlled. The plasma generated during RF sputtering consists of virtues like high energy and extreme purity with a mean free path up to 10 cm long [37]. The abovementioned intrinsic worth makes the sputtered atoms diffuse quickly with high extra energy around the growing film leaving a homogeneous film avoiding the aggregation of atoms [27]. Herein, we investigate the effect of deposition power on the structural and ionic conductivity properties of as deposited films applied as solid electrolyte of TFBs.
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2. Experimental The preparation of Li–La–Zr–O target which involves a solid-state method is reported elsewhere [38]. Li2CO3, La2O3, and ZrO2 were used in powder form as the starting materials. The starting materials were predried and then ball milled with zirconia balls (⌀ =0.2 mm) for 6 h in 2-propanol. In order to compensate the loss of lithium by volatilization during heat-treatments, 10 wt.% excess of Li2CO3 (10 wt.%) was added in the starting mixture resulting in the molar ratio of Li:La:Zr=7:3:2. The milling process was repeated after each heat-treatment (first, at 900 °C for 6 h and second, at 1125 °C for 6 h) in air. As obtained mixture was finally pressed in the form of pellet under 50 MPa for 5 min and subsequently annealed at 1230 °C for 36 h at a heating rate of 1 °C/min in air. The crystal structure of Li–La–Zr–O was confirmed after each heat-treatment by X-ray diffractometer (XRD, Rigaku RINT) using Cu Kα radiation. SiO2–Si wafers used in the experiment were cleaned in isopropyl alcohol and then in acetone each for 15 min by sonication, and then rinsed thoroughly with deionized water. After cleaning the substrates, the bottom electrode was deposited maintaining a working pressure of 5× 10−3 Torr by giving a constant Ar flow of 40 sccm. Li– La–Zr–O thin films were deposited on the bottom SUS electrode using RF power of 40, 60, and 80 W at a working pressure of 5 × 10−3 Torr under Ar atmosphere. After the deposition of Li–La–Zr–O thin films an upper SUS electrode was grown on each sample by the same procedure used for the bottom electrode. These fabrication steps are schematically shown in Fig. 1. The final cell was a multilayered structure of SUS/Li–La– Zr–O/SUS/SiO2–Si. After deposition of the thin films the crystal structure characterization was performed by XRD (Rigaku RINT) measurement with CuKα1 radiation at a scan rate of 2°/min. X-ray photoelectron spectroscopy (XPS, ESCALAB MK II, V.G. Scientific Ltd.) and energy dispersive X-ray (EDX) spectroscopy (JEOL-JSM-700 IF) analyses were carried out to confirm the chemical state and atomic composition of the Li–La–Zr– O thin films. Surface morphology observation and thickness calculation of the samples were carried out by field emission scanning electron microscopy (FE-SEM, JEOL-JSM-700 IF). Additionally, impedance and activation energies of the multilayered SUS/Li–La–Zr–O/SUS/Si cell structures were measured by impedance analyzer (IM6ex, Zahner). A frequency range of 1 Hz to 3 MHz was used at room temperature. 3. Results and discussion The surface morphology and thickness of the films are shown in Fig. 2. From Fig. 2 (a, b and c) it can be concluded that the surfaces of the thin
Fig. 1. Schematic illustration of cell fabrication processes.
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Fig. 2. SEM images of the surface (a, b, c) and cross section (d, e, f) of Li–La–Zr–O thin films deposited at 40 W (a, d), 60 W (b, e) and 80 W (c, f).
films are obviously smooth and dense without significant formation of pinholes and cracks. However, it is noteworthy that a lower deposition power leads to a better surface uniformity of the thin film as the high power deposition leads to a slightly vigorous growth which gives some non-uniformity in the films (Fig. 2b, c). Smooth morphology of the surfaces enables the thin film to avoid safety problems by decreasing the contact resistance with the electrodes. A gradual increase in the thickness of the thin films has been observed [39] (Fig. 2d, e and f) with increasing the RF power and can be controlled by tuning the sputtering rate and time. The bulk film is dense and homogeneous, as shown by the cross section of the films. The surface morphology of the films without significant grain boundaries indicates that RF sputtering is convenient in the formation of the amorphous film, and is augmented by X-ray diffraction analysis. Fig. 3 shows the XRD patterns of the Li–La–Zr–O thin films, the XRD patterns of the films deposited at different RF power are devoid of any dominant peaks and have almost flat diffraction pattern except for a broad XRD peak between 40° and 50° which starts to appear with increase in the deposition power. This result indicates that a very small amount of atomic arrangement increases when deposition is carried with relatively high RF power. This phenomenon can be explained as a consequence of acquiring relatively moderate kinetic energy by the sputter ejected species which increases the possibility of atomic alignment and crystallinity [38,40]. This atomic arrangement however is not sufficient to bring ionic transport through hopping mechanism like in crystalline phase and dominated by some hierarchical structure [41]. The compositions of thin films were investigated by EDX (Table 1) spectroscopy and further confirmed with the XPS data (Table 2) as EDX does not provide data about the Li content. The reason for the concentration variation of elements with deposition power is due to the differential sputtering yield for each under the same glow-discharge plasma conditions which is a typical phenomenon in the sputtering process [29,42].
The element content of thin films was almost reliable as they were protected in argon filled glove box. Li (1s) spectra for films deposited at different RF power (40, 60 and 80 W) are shown in the inset of Fig. 4. The peak positions at approximately 55 eV indicate the presence of Li in Li2O form. A gradual decrease in the peak intensity is observed with increasing RF-deposition power. This is in agreement with the O 1s peak observed at the 528.5 eV which is attributed to lithium oxide [43] shown in Fig. 4. This decrease in the peak intensity and the area under the curves shows a lower concentration of Li for the films depositing at higher RF power (40 W>60 W>80 W). The peak at 531 eV is generally assigned to the bonding metallic oxygen. The peak intensity shows a gradual increase with increasing deposition power highest for the 80 W sample. This indicates a greater concentration of La and Zr in
Fig. 3. XRD patterns of the Li–La–Zr–O thin films deposited at 40 W (a), 60 W (b) and at 80 W(c) deposited at room temperature.
D.J. Kalita et al. / Solid State Ionics 229 (2012) 14–19
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Table 1 EDX analysis on the surface of Li–La–Zr–O thin films deposited with RF-powers (40, 60, and 80 W). Deposition power
Compositional analysis (wt.% and at.%) La
40 W 60 W 80 W
Zr
O
wt.%
at.%
wt.%
at.%
wt.%
at.%
34.60 42.81 46.42
8.13 11.54 14
24.67 23.54 25.1
8.82 9.67 11.55
40.73 33.65 28.43
83.05 78.79 74.45
La2O3 and ZrO2 form (confirmed from the La3d5/2 peak at 834 eV and Zr3d5/2 peak at 181.5 eV) in the films deposited at higher RF power which is in agreement with the data obtained from EDX analysis. Finally, a peak at 532.7 eV corresponding to oxygen in water molecules was observed which does not follow any such increasing or decreasing trend with deposition power. It may be due to exposure of the thin films to the atmosphere during experimental work [44]. From the above observations it can be inferred that deposition at a lower RF-power increases the deposition of Li which would lead to a higher ionic conductivity as it is not involved in the formation of network or acts as network modifier [41]. Investigation already accomplished by a number of research groups unveiled that network modifiers, basically alkali oxide such as Na2O, when added, result in partial intervention of the network resulting in higher non-bridging oxygen (NBO) [45]. These introduced Na+ ions are bounded to the non-bridging oxygen weakly, permitting to contribute to the ionic conductivity only at high temperature not at the room temperature like Li does in our study. It is undisputable to refer Li2O as the soft part that is primarily responsible for the ionic conduction as Li+. Fig. 5 shows the impedance spectrum of the 40 W, 60 W and 80 W samples with a SUS/Li–La–Zr–O/SUS/SiO2–Si cell structure at room temperature has been carried out with Zahner IM6 Impedance Analyzer. The appearance of the semicircle with a spike in the high frequency region clearly shows a bulk conduction mechanism as there are no grain boundaries present in the amorphous Li–La–Zr–O thin films. However, a significant increase in the semicircle radius had been observed with a higher deposition power which indicates a higher ionic resistance. The ionic conductivity (σ) of the Li–La–Zr–O films was calculated according to equation [30–35] σ¼
d RA
ð1Þ
where, d, R and A represent the thickness of solid-state electrolyte, electrolyte resistance, and electrode area respectively. The thin film deposited at 40 W gives the highest ionic conductivity of 4 ×10−7 S/cm when compared to both the 60 W and 80 W samples. This conductivity measurement reveals that in comparison to earlier investigated amorphous electrolytes such as Li–B–W–O (2.15× 10−7 S/cm) [23], LiPON– LLT–LiPON (1.29× 10−7 S/cm) [17] amorphous Li–La–Zr–O has higher ionic conductivity subsequently rising its possibility as solid electrolyte. This increase in the ionic conductivity with a lower deposition power can't be a structural effect like grain boundary imposed resistance to ionic conductivity for the films deposited at higher RF power. As from the SEM analysis, it is seen that the surfaces of the thin films are almost
Fig. 4. XPS spectrum of 0 1s and Li 1s (in the inset) peaks for 40, 60, and 80 W samples.
smooth, uniform, and without substantial grain formation with increasing deposition power. However, an explanation for this trend in increasing the ionic conductivity with lower deposition power can be drawn from the XPS analysis which shows a higher concentration of Li for the film deposited at 40 W RF. The mobility of Li ions has a direct effect on the ionic conductivity of the Li–La–Zr–O thin films. Unlike aqueous electrolyte, the mobility of ion retarded the mobility governed by the equation λi = μi (zi·F), in which solid electrolyte follows a different transport mechanism and induces an increase in ionic conductivity with increasing Li concentration [46]. Again, in amorphous oxy electrolytes Li2O acts as a Li soft short range order part and thin film composition with a higher concentration of Li soft part (Li2O in our electrolyte) gives a higher conductivity and lower activation energy for Li+ ion transport [47,48] which is similar to our experimental results. Activation energy (Ea) was calculated by the following equation [30–49] using the Arrhenius equation and measured conductivity at various temperatures. The plots are shown in Fig. 6 σT ¼ ðσt Þ exp
−Ea kT
ð2Þ
where, σ is the ionic conductivity, k is the Boltzmann constant, and T is the temperature. The activation energy of Li–La–Zr–O solid-state electrolyte deposited at 40 W shows lower activation energy of 0.70 eV than that of 60 W and 80 W samples with an activation of 0.81 eV and 0.87 eV, respectively. A minimum electrical conductivity with a high
Table 2 The molar ratio of each component of Li–La–Zr–O thin films deposited with RF-powers (40, 60, and 80 W) from XPS analysis. Deposition power
40 W 60 W 80 W
Elements (molar ratio) Li
La
Zr
O
2.9 2.6 2.2
0.68 0.73 0.81
1 1 1
8 11 15
Fig. 5. Electrochemical impedance spectroscopy of the Li–La–Zr–O thin films deposited at 40 W, 60 W and 80 W RF-power.
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4. Conclusions
Fig. 6. Arrhenius plots for ionic conductivity of Li–La–Zr–O thin films deposited at 40 W, 60 W and 80 W as a function of temperature.
ionic conductivity has been considered as a basic requirement in order to prohibit the self discharge. A high self discharge may lead to a loss of capacity along with internal heat up which may cause severe damage. To ensure the electrical conductivity the I–V curve was measured which depicts the amount of current flow at a particular voltage. The measurement was carried out up to 5 V at 10 mV/s scan rate. The results are shown in Fig. 7. Although an electric conductivity of about 9 μA has been permitted by the Li–La–Zr–O thin films with increasing the voltage, no short circuit pattern of the curves has been obtained. Again, this value is far lower than that shown by LLT sandwich by LiPON (deposited up to 30 min) [50]. This means that the micro-battery with Li–La–Zr–O solid electrolyte will show the low self-discharge during the un-operating period. The amorphous structure and non-equilibrium nature usually present obscureness in conceiving the ionic transport mechanism. An explanation of the observed phenomena can be accounted for by the presence of both the static component like zirconium and lanthanum oxides and the mobile soft component of lithium, whose, presence is confirmed from the XPS data. Ionic transport is most strongly affected by the presence of this soft part. With a decrease in the deposition power a higher content of the Li soft part has been achieved and it leads to a higher ionic conductivity of the films deposited at lower RF-power. These results suggest that Li–La–Zr–O thin films deposited with a lower RF power can serve with better response in solid-state thin film power devices.
Fig. 7. I–V curve of Li–La–Zr–O thin films deposited at 40 W, 60 W and 80 W as a function of temperature.
Li–La–Zr–O thin films were successfully deposited by RF magnetron sputtering at different power and their ionic conductivity study has been investigated. From the XRD patterns, the films deposited at room temperature present amorphous structure. SEM results of thin films show a smooth surface without cracks and pits which is very important for film electrolytes to avoid short-circuiting in the battery application. The ionic conductivity of the films increases with a lower deposition power, up to a value of 4 × 10−7 S/cm with the lowest activation energy of 0.70 eV. This enhancement of ionic conductivity with a lower deposition power can be explained from the XPS analysis. XPS analysis shows a higher concentration of Li2O in the films deposited at 40 W RF power than that of 60 and 80 W samples which act as the Li modifier in oxy electrolytes. The higher content of Li2O should lower the activation energy for Li+ ion transport which is confluent to our experimental results. We propose that Li–La–Zr–O thin films may be a promising solid electrolyte for thin film Li ion batteries (LIBs) of all-solid-state. Acknowledgments We are grateful to the KARI-University Partnership Programfor their financial support. This work was supported by the Priority Research Centers Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (2009-0093823). This research was supported by the Pioneer Research Center Program through the National Research Foundation of Korea funded by the Ministry of Education, Science and Technology (2010-0019313). We are also thankful to the Hyundae Motor Group for supporting this research. References [1] S.J. Lee, H.K. Baik, S.M. Lee, Electrochem. Commun. 5 (2003) 32–35. [2] F. Albano, M.D. Chung, D. Blaauw, D.M. Sylvester, K.D. Wise, A.M. Sastry, J. Power. Sources 170 (2007) 216–224. [3] P.H.L. Notten, F. Roozeboom, R.A.H. Niessen, L. Baggetto, Adv. Mater. 19 (2007) 4564–4567. [4] F.C. Liu, W.M. Liu, M.H. Zhan, Z.W. Fu, H. Li, Energy Environ. Sci. 4 (2011) 1261–1264. [5] J.M. Tarascon, M. Armand, Nature 414 (2001) 359–367. [6] P. Birke, W. Weppner, Electrochim. Acta 42 (1997) 3375–3384. [7] Y.S. Yoon, S.B. Cho, S.C. Nam, Electrochim. Acta 71 (2012) 86–91. [8] P. Vinatier, P. Vinatier, B. Pecquenard, A. Levasseur, H.J. Sohn, Solid State Ionics 160 (2003) 51–59. [9] X. Yu, J.B. Bates, G.E. Jellison Jr., F.X. Hart, J. Electrochem. Soc. 144 (1997) 524–532. [10] Y. Inaguma, C. Liquan, M. Itoh, T. Nakamura, T. Uchida, H. Ikuta, M. Wakihara, Solid State Commun. 86 (1993) 689–693. [11] T. Abe, M. Ohtsuka, F. Sagane, Y. Iriyama, Z. Ogumi, J. Electrochem. Soc. 151 (2004) A1950–A1953. [12] H. Aono, E. Sugimoto, Y. Sadaoka, N. Imanaka, G. Adachi, J. Electrochem. Soc. 137 (1990) 1023–1027. [13] J. Fu, Solid State Ionics 96 (1997) 195–200. [14] P. Knauth, Solid State Ionics 180 (2009) 911–916. [15] V. Thangadurai, H. Kaack, W. Weppner, J. Am. Ceram. Soc. 86 (2003) 437–440. [16] R. Murugan, V. Thangadurai, W. Weppner, Angew. Chem. Int. Ed. 46 (2007) 7778–7781. [17] F. Vaz, L. Rebouta, P. Goudeau, J.P. Riviere, E. Schaffer, G. Kleer, M. Bodmann, Thin Solid Films 402 (2002) 195–202. [18] M. Bromark, M. Larsson, P. Hedenqvist, M. Olsson, S. Hogmark, E. Bergmann, Surf. Eng. 10 (1994) 205–214. [19] G. Kleer, R. Kassner, E.M. Meyer, M.G. Schinker, W. Do¨ll, Surf. Coat.Technol. 54 (1992) 167–170. [20] R.W. Hoffman, Surf. Interface Anal. 3 (1981) 62–66. [21] M.A. Moske, P.S. Ho, D.J. Mikalsen, J.J. Cuomo, R. Rosenberg, J. Appl. Phys. 74 (1993) 1716–1724. [22] J.K. Ahn, S.G. Yoon, Electrochem.Solid State lett. 8 (2) (2005) A75–A78. [23] Y.-M. Chiang, D.R. Sadoway, Y.I. Jang, B. Huang, H. Wang, Electrochem. Solid-State Lett. 2 (1999) 107–110. [24] T. Ohzuku, T. Yanagawa, M. Kouguchi, A. Ueda, J. Power. Sources 68 (1997) 131–134. [25] S. Buta, D. Morgan, A. Van der Ven, M.K. Aydinol, G. Ceder, J. Electrochem. Soc. 146 (1999) 4335–4338. [26] C. Wolverton, Acta Mater. 49 (2001) 3129–3142. [27] G. Tan, F. Wu, L. Li, Y. Liu, R. Chen, J. Phys. Chem. C 116 (2012) 3817–3821. [28] H. Kitaura, A. Hayashi, T. Ohtomo, S. Hama, M. Tatsumisago, J. Mater. Chem. 21 (2011) 118–124. [29] E. Quartarone, P. Mustarelli, Chem. Soc. Rev. 40 (2011) 2525–2540.
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Ionic conductivity properties of amorphous Li–La–Zr–O solid electrolyte for thin film batteries D.J. Kalita a, S.H. Lee a, K.S. Lee a, D.H. Ko b, Y.S. Yoon a,⁎ a b
Department of Materials Science and Engineering, Yonsei University, Shinchon Dong 134, Seoul 120-749, Republic of Korea Korea Aerospace Research Institute/Payload Optics Team, 115 Gwahangno, Yuseong-gu, Daejeon 305-333, Republic of Korea
a r t i c l e
i n f o
Article history: Received 8 June 2012 Received in revised form 14 September 2012 Accepted 14 September 2012 Available online 17 November 2012 Keywords: Thin film battery Solid electrolyte RF-power Li-modifier
a b s t r a c t Lithium–lanthanum–zirconium–oxide (Li–La–Zr–O) thin films were deposited using radio-frequency (RF) magnetron sputtering by varying the RF power from 40 to 80 W. The characteristics are investigated with X-ray diffraction analysis (XRD) and Field Emission Scanning Electron Microscope (FE-SEM). Results showed that the thin films had amorphous structure accompanied with a smooth, dense, and homogeneous surface. Impedance analysis was conducted to investigate the ionic conductivity. The thin film deposited at 40 W gives the highest ionic conductivity of 4×10−7 S/cm when compared to both the 60 W and 80 W samples which show the ionic conductivity of 2×10−7 S/cm and 0.8×10−7 S/cm respectively. The temperature dependence of the ionic conductivities fit the Arrhenius relation and the thin film deposited at 40 W possessed the lowest activation energy. These properties make Li–La–Zr–O thin films as a promising candidate material for use in all-solid-state thin film lithium ion batteries. © 2012 Elsevier B.V. All rights reserved.
1. Introduction The need of microenergy sources with high power density is inevitable for the miniaturization of electronic devices like implantable medical devices, electronic paper, RFID-Tag, microelectromechanical systems (MEMS), etc. [1–4]. Although most of the modern battery technologies are able to partially satisfy this requirement, thin film lithium ion batteries (LIBs) are considered as the most promising candidate for their ability to provide the maximum voltage and energy storage densities [5,6] in their microstate. Electrolyte which is one of the major components in LIBs may be found either in solid-state or in liquid-state. In comparison with the liquid-state electrolytes, solid-state electrolytes have not provided satisfactory performance to meet the requirements of practical field applications that required high ionic conductivity, and chemical and electrochemical stability [7]. However, a wider range of potential field applications are provided by the solid state electrolytes than the liquid-state electrolytes. Liquid-state electrolytes are accompanied by the inherent disadvantages such as short lifetime, leakage of electrolyte and nonfunctioning of device below the freezing point and above the boiling point [8]. The fabrication of thin film forms of energy storage devices is possible with the solid state electrolytes. Recently, a number of solid electrolytes having perovskite structure such as Li0.35La0.55TiO3 (LLT) [9–11] and with NASICON structure such as LiTi2(PO4)3 (LTP) [12,13] have been reported with a high bulk ionic conductivity in the range of ~10−3 to 10−4 S/cm, respectively. Although these materials were accepted for LIBs, recent investigation reveals that anode materials ⁎ Corresponding author. Tel.: +82 2 2123 2847; fax: +82 2 2123 8145. E-mail address: [email protected] (Y.S. Yoon). 0167-2738/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.ssi.2012.09.011
having red-ox potential lower than 1.8 V vs. Li/Li+ cannot be employed because of the red-ox reaction of titanium with lithium that causes reduction of Ti4+ to Ti3+ in bulk [14], which restricted the energy and power densities of LIBs to a far lower value. During the last several years, a series of compounds with a nominal composition of Li5La3M2O12 (M=Nb, Ta) and garnet like structure has been investigated by Thangadurai et al. [15]. However, a better chemical stability against lithium anode and a relatively high ionic conductivity have been obtained for Li7La3Zr2O12 (LLZ) by Murugan et al. [16]. Both the phase property and fabrication processes of solid electrolytes have crucial roles in their electrochemical properties. Earlier investigations make it obvious that solid electrolyte with a garnet like or crystalline structure gives an ionic conductivity up to 10−4 S cm−1, preparation of which usually involves a high temperature step like conventional high temperature annealing, rapid thermal annealing (RTA) and deposition with a high temperature substrate heating, etc. All these high temperature processes have adverse effects when the fabrication involves thin films approximately 1 μm owing to the current demand for thin film LIBs. A crucial matter that remains as blur and muddling to the scientist is to optimize certain properties of the coating substrate composite such as hardness, evolution of residual stress, fracture roughness and other mechanical properties. The treatment of thin films at elevated temperature around 700 °C results in being prone to failure [17]. In most of the PVD coatings like sputtering, laser deposition, etc., this failure results from the development of stresses and interlayer inter-diffusion resulting to severe damages such as film cracking, de-cohesion by buckling and inter-phase [18–20]. A deposition annealing process always induces thermal stress in the thin films which may result to the formation of micro-cracks even at moderate temperature like 400 °C [19]. A number of studies
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have been done to tackle these kinds of serious problems by using adhesive layers such as Cr, TiN. However these multilayer interconnects are also influenced by the development of thermal stress during heat treatment resulting in void formation and crack propagation [18]. Differences or mismatches in the thermal and mechanical properties causing large thermal stresses give rise to these menacing problems during annealing of thin films [21]. Employment of high temperature (>600 °C) during crystallization of thin films leads to a severe loss of Li content and may cause some unwanted phase separation. Investigation by Ahn et al. inferred that crystallizations of thin LTO thin films cause segregation of TiO2 phase incorporating holes at grain boundary for which the microstructures experience electric short cuts [22]. This phase separation has become one of the concerning factors preventing uses of dopants like Al to a certain extent, at high temperature. Chiang et al. [23] and Ohjuku et al. [24] had reported the use of Al substitution to accomplish high temperature stability to the Mn and Ni based oxides. However, later investigation by S. Buta et al. [25] showed that Al possesses a lower miscibility with a number of transition metal oxides. This results in the formation of LiAlO2 phase as it has positive mixing enthalpies with lithium transition-metal oxides. C. Wolverton's [26] study revealed that Al also has a high tendency to form precipitate phase in multicomponent systems containing transition metal. These kinds of drawbacks always make high temperature thin film processes accompanied with failure of experiments. Because of such constraints, amorphous electrolytes are broadly investigated for having resultant denser microstructure with negligible grain boundary effect, and improved chemical stability for applying in all-solid-state LIBs [27], although they result in moderate ionic conductivity compared to a crystalline one. Amorphous electrolytes which permit only the migration of lithium ion can be categorized either as sulfide [4] or an oxide series by element [28,31]. In comparison to oxide series sulfide series exhibit high ionic conductivity and thermal stability but in practical purposes it is hard to use them because of their high instability when they come into contact with moisture in the air. On the other hand, oxide series electrolytes are chemically and electrochemically more stable than sulfide series solid-state electrolytes at normal atmospheric conditions [32,33]. Therefore, solid-state oxide electrolytes are more potent for TFBs and TFSCs. In this study, we focused on the room temperature fabrication and characterization of lithium– lanthanum–zirconium–oxide (Li–La–Zr–O) thin films having amorphous structure by RF sputtering. Our aptness towards this material's amorphous synthesis lies on both the pros and cons associated with it. The prime interest about this material is its possession of high stability against lithium anode and also its high bulk ionic conductivity resulting it to show up against several solid electrolytes like LATP and LATO [34]. Both these features are the prime requirement for a solid electrolyte in thin film Li ion battery. Investigation by Geiger et al. [35] uncovered that Li–La–Zr–O results in a stable cubic phase with bulk ionic conductivity of 10−4 S/cm in the presence of Al as a phase stabilizer. However, the temperature requirement for this process is above 900 °C which results in severe damages to the thin film such as cracking and the total failure of the system. Again, Xie et al. [36] reported that processing of this material at around 850 °C gives rise to a tetragonal phase and ionic conductivity of 10−6 S/cm making the process unprofitable. A room temperature synthetic process of this material is more justifiable if it leads to a sufficient ionic conductivity (10−6 S/cm) expelling the cons related to high temperature processes. As compared with other fabrication methods radio frequency (RF) magnetron sputtering is considered to be an easy, established and environmentally benign process of deposition which can be easily controlled. The plasma generated during RF sputtering consists of virtues like high energy and extreme purity with a mean free path up to 10 cm long [37]. The abovementioned intrinsic worth makes the sputtered atoms diffuse quickly with high extra energy around the growing film leaving a homogeneous film avoiding the aggregation of atoms [27]. Herein, we investigate the effect of deposition power on the structural and ionic conductivity properties of as deposited films applied as solid electrolyte of TFBs.
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2. Experimental The preparation of Li–La–Zr–O target which involves a solid-state method is reported elsewhere [38]. Li2CO3, La2O3, and ZrO2 were used in powder form as the starting materials. The starting materials were predried and then ball milled with zirconia balls (⌀ =0.2 mm) for 6 h in 2-propanol. In order to compensate the loss of lithium by volatilization during heat-treatments, 10 wt.% excess of Li2CO3 (10 wt.%) was added in the starting mixture resulting in the molar ratio of Li:La:Zr=7:3:2. The milling process was repeated after each heat-treatment (first, at 900 °C for 6 h and second, at 1125 °C for 6 h) in air. As obtained mixture was finally pressed in the form of pellet under 50 MPa for 5 min and subsequently annealed at 1230 °C for 36 h at a heating rate of 1 °C/min in air. The crystal structure of Li–La–Zr–O was confirmed after each heat-treatment by X-ray diffractometer (XRD, Rigaku RINT) using Cu Kα radiation. SiO2–Si wafers used in the experiment were cleaned in isopropyl alcohol and then in acetone each for 15 min by sonication, and then rinsed thoroughly with deionized water. After cleaning the substrates, the bottom electrode was deposited maintaining a working pressure of 5× 10−3 Torr by giving a constant Ar flow of 40 sccm. Li– La–Zr–O thin films were deposited on the bottom SUS electrode using RF power of 40, 60, and 80 W at a working pressure of 5 × 10−3 Torr under Ar atmosphere. After the deposition of Li–La–Zr–O thin films an upper SUS electrode was grown on each sample by the same procedure used for the bottom electrode. These fabrication steps are schematically shown in Fig. 1. The final cell was a multilayered structure of SUS/Li–La– Zr–O/SUS/SiO2–Si. After deposition of the thin films the crystal structure characterization was performed by XRD (Rigaku RINT) measurement with CuKα1 radiation at a scan rate of 2°/min. X-ray photoelectron spectroscopy (XPS, ESCALAB MK II, V.G. Scientific Ltd.) and energy dispersive X-ray (EDX) spectroscopy (JEOL-JSM-700 IF) analyses were carried out to confirm the chemical state and atomic composition of the Li–La–Zr– O thin films. Surface morphology observation and thickness calculation of the samples were carried out by field emission scanning electron microscopy (FE-SEM, JEOL-JSM-700 IF). Additionally, impedance and activation energies of the multilayered SUS/Li–La–Zr–O/SUS/Si cell structures were measured by impedance analyzer (IM6ex, Zahner). A frequency range of 1 Hz to 3 MHz was used at room temperature. 3. Results and discussion The surface morphology and thickness of the films are shown in Fig. 2. From Fig. 2 (a, b and c) it can be concluded that the surfaces of the thin
Fig. 1. Schematic illustration of cell fabrication processes.
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Fig. 2. SEM images of the surface (a, b, c) and cross section (d, e, f) of Li–La–Zr–O thin films deposited at 40 W (a, d), 60 W (b, e) and 80 W (c, f).
films are obviously smooth and dense without significant formation of pinholes and cracks. However, it is noteworthy that a lower deposition power leads to a better surface uniformity of the thin film as the high power deposition leads to a slightly vigorous growth which gives some non-uniformity in the films (Fig. 2b, c). Smooth morphology of the surfaces enables the thin film to avoid safety problems by decreasing the contact resistance with the electrodes. A gradual increase in the thickness of the thin films has been observed [39] (Fig. 2d, e and f) with increasing the RF power and can be controlled by tuning the sputtering rate and time. The bulk film is dense and homogeneous, as shown by the cross section of the films. The surface morphology of the films without significant grain boundaries indicates that RF sputtering is convenient in the formation of the amorphous film, and is augmented by X-ray diffraction analysis. Fig. 3 shows the XRD patterns of the Li–La–Zr–O thin films, the XRD patterns of the films deposited at different RF power are devoid of any dominant peaks and have almost flat diffraction pattern except for a broad XRD peak between 40° and 50° which starts to appear with increase in the deposition power. This result indicates that a very small amount of atomic arrangement increases when deposition is carried with relatively high RF power. This phenomenon can be explained as a consequence of acquiring relatively moderate kinetic energy by the sputter ejected species which increases the possibility of atomic alignment and crystallinity [38,40]. This atomic arrangement however is not sufficient to bring ionic transport through hopping mechanism like in crystalline phase and dominated by some hierarchical structure [41]. The compositions of thin films were investigated by EDX (Table 1) spectroscopy and further confirmed with the XPS data (Table 2) as EDX does not provide data about the Li content. The reason for the concentration variation of elements with deposition power is due to the differential sputtering yield for each under the same glow-discharge plasma conditions which is a typical phenomenon in the sputtering process [29,42].
The element content of thin films was almost reliable as they were protected in argon filled glove box. Li (1s) spectra for films deposited at different RF power (40, 60 and 80 W) are shown in the inset of Fig. 4. The peak positions at approximately 55 eV indicate the presence of Li in Li2O form. A gradual decrease in the peak intensity is observed with increasing RF-deposition power. This is in agreement with the O 1s peak observed at the 528.5 eV which is attributed to lithium oxide [43] shown in Fig. 4. This decrease in the peak intensity and the area under the curves shows a lower concentration of Li for the films depositing at higher RF power (40 W>60 W>80 W). The peak at 531 eV is generally assigned to the bonding metallic oxygen. The peak intensity shows a gradual increase with increasing deposition power highest for the 80 W sample. This indicates a greater concentration of La and Zr in
Fig. 3. XRD patterns of the Li–La–Zr–O thin films deposited at 40 W (a), 60 W (b) and at 80 W(c) deposited at room temperature.
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Table 1 EDX analysis on the surface of Li–La–Zr–O thin films deposited with RF-powers (40, 60, and 80 W). Deposition power
Compositional analysis (wt.% and at.%) La
40 W 60 W 80 W
Zr
O
wt.%
at.%
wt.%
at.%
wt.%
at.%
34.60 42.81 46.42
8.13 11.54 14
24.67 23.54 25.1
8.82 9.67 11.55
40.73 33.65 28.43
83.05 78.79 74.45
La2O3 and ZrO2 form (confirmed from the La3d5/2 peak at 834 eV and Zr3d5/2 peak at 181.5 eV) in the films deposited at higher RF power which is in agreement with the data obtained from EDX analysis. Finally, a peak at 532.7 eV corresponding to oxygen in water molecules was observed which does not follow any such increasing or decreasing trend with deposition power. It may be due to exposure of the thin films to the atmosphere during experimental work [44]. From the above observations it can be inferred that deposition at a lower RF-power increases the deposition of Li which would lead to a higher ionic conductivity as it is not involved in the formation of network or acts as network modifier [41]. Investigation already accomplished by a number of research groups unveiled that network modifiers, basically alkali oxide such as Na2O, when added, result in partial intervention of the network resulting in higher non-bridging oxygen (NBO) [45]. These introduced Na+ ions are bounded to the non-bridging oxygen weakly, permitting to contribute to the ionic conductivity only at high temperature not at the room temperature like Li does in our study. It is undisputable to refer Li2O as the soft part that is primarily responsible for the ionic conduction as Li+. Fig. 5 shows the impedance spectrum of the 40 W, 60 W and 80 W samples with a SUS/Li–La–Zr–O/SUS/SiO2–Si cell structure at room temperature has been carried out with Zahner IM6 Impedance Analyzer. The appearance of the semicircle with a spike in the high frequency region clearly shows a bulk conduction mechanism as there are no grain boundaries present in the amorphous Li–La–Zr–O thin films. However, a significant increase in the semicircle radius had been observed with a higher deposition power which indicates a higher ionic resistance. The ionic conductivity (σ) of the Li–La–Zr–O films was calculated according to equation [30–35] σ¼
d RA
ð1Þ
where, d, R and A represent the thickness of solid-state electrolyte, electrolyte resistance, and electrode area respectively. The thin film deposited at 40 W gives the highest ionic conductivity of 4 ×10−7 S/cm when compared to both the 60 W and 80 W samples. This conductivity measurement reveals that in comparison to earlier investigated amorphous electrolytes such as Li–B–W–O (2.15× 10−7 S/cm) [23], LiPON– LLT–LiPON (1.29× 10−7 S/cm) [17] amorphous Li–La–Zr–O has higher ionic conductivity subsequently rising its possibility as solid electrolyte. This increase in the ionic conductivity with a lower deposition power can't be a structural effect like grain boundary imposed resistance to ionic conductivity for the films deposited at higher RF power. As from the SEM analysis, it is seen that the surfaces of the thin films are almost
Fig. 4. XPS spectrum of 0 1s and Li 1s (in the inset) peaks for 40, 60, and 80 W samples.
smooth, uniform, and without substantial grain formation with increasing deposition power. However, an explanation for this trend in increasing the ionic conductivity with lower deposition power can be drawn from the XPS analysis which shows a higher concentration of Li for the film deposited at 40 W RF. The mobility of Li ions has a direct effect on the ionic conductivity of the Li–La–Zr–O thin films. Unlike aqueous electrolyte, the mobility of ion retarded the mobility governed by the equation λi = μi (zi·F), in which solid electrolyte follows a different transport mechanism and induces an increase in ionic conductivity with increasing Li concentration [46]. Again, in amorphous oxy electrolytes Li2O acts as a Li soft short range order part and thin film composition with a higher concentration of Li soft part (Li2O in our electrolyte) gives a higher conductivity and lower activation energy for Li+ ion transport [47,48] which is similar to our experimental results. Activation energy (Ea) was calculated by the following equation [30–49] using the Arrhenius equation and measured conductivity at various temperatures. The plots are shown in Fig. 6 σT ¼ ðσt Þ exp
−Ea kT
ð2Þ
where, σ is the ionic conductivity, k is the Boltzmann constant, and T is the temperature. The activation energy of Li–La–Zr–O solid-state electrolyte deposited at 40 W shows lower activation energy of 0.70 eV than that of 60 W and 80 W samples with an activation of 0.81 eV and 0.87 eV, respectively. A minimum electrical conductivity with a high
Table 2 The molar ratio of each component of Li–La–Zr–O thin films deposited with RF-powers (40, 60, and 80 W) from XPS analysis. Deposition power
40 W 60 W 80 W
Elements (molar ratio) Li
La
Zr
O
2.9 2.6 2.2
0.68 0.73 0.81
1 1 1
8 11 15
Fig. 5. Electrochemical impedance spectroscopy of the Li–La–Zr–O thin films deposited at 40 W, 60 W and 80 W RF-power.
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4. Conclusions
Fig. 6. Arrhenius plots for ionic conductivity of Li–La–Zr–O thin films deposited at 40 W, 60 W and 80 W as a function of temperature.
ionic conductivity has been considered as a basic requirement in order to prohibit the self discharge. A high self discharge may lead to a loss of capacity along with internal heat up which may cause severe damage. To ensure the electrical conductivity the I–V curve was measured which depicts the amount of current flow at a particular voltage. The measurement was carried out up to 5 V at 10 mV/s scan rate. The results are shown in Fig. 7. Although an electric conductivity of about 9 μA has been permitted by the Li–La–Zr–O thin films with increasing the voltage, no short circuit pattern of the curves has been obtained. Again, this value is far lower than that shown by LLT sandwich by LiPON (deposited up to 30 min) [50]. This means that the micro-battery with Li–La–Zr–O solid electrolyte will show the low self-discharge during the un-operating period. The amorphous structure and non-equilibrium nature usually present obscureness in conceiving the ionic transport mechanism. An explanation of the observed phenomena can be accounted for by the presence of both the static component like zirconium and lanthanum oxides and the mobile soft component of lithium, whose, presence is confirmed from the XPS data. Ionic transport is most strongly affected by the presence of this soft part. With a decrease in the deposition power a higher content of the Li soft part has been achieved and it leads to a higher ionic conductivity of the films deposited at lower RF-power. These results suggest that Li–La–Zr–O thin films deposited with a lower RF power can serve with better response in solid-state thin film power devices.
Fig. 7. I–V curve of Li–La–Zr–O thin films deposited at 40 W, 60 W and 80 W as a function of temperature.
Li–La–Zr–O thin films were successfully deposited by RF magnetron sputtering at different power and their ionic conductivity study has been investigated. From the XRD patterns, the films deposited at room temperature present amorphous structure. SEM results of thin films show a smooth surface without cracks and pits which is very important for film electrolytes to avoid short-circuiting in the battery application. The ionic conductivity of the films increases with a lower deposition power, up to a value of 4 × 10−7 S/cm with the lowest activation energy of 0.70 eV. This enhancement of ionic conductivity with a lower deposition power can be explained from the XPS analysis. XPS analysis shows a higher concentration of Li2O in the films deposited at 40 W RF power than that of 60 and 80 W samples which act as the Li modifier in oxy electrolytes. The higher content of Li2O should lower the activation energy for Li+ ion transport which is confluent to our experimental results. We propose that Li–La–Zr–O thin films may be a promising solid electrolyte for thin film Li ion batteries (LIBs) of all-solid-state. Acknowledgments We are grateful to the KARI-University Partnership Programfor their financial support. This work was supported by the Priority Research Centers Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (2009-0093823). This research was supported by the Pioneer Research Center Program through the National Research Foundation of Korea funded by the Ministry of Education, Science and Technology (2010-0019313). We are also thankful to the Hyundae Motor Group for supporting this research. References [1] S.J. Lee, H.K. Baik, S.M. Lee, Electrochem. Commun. 5 (2003) 32–35. [2] F. Albano, M.D. Chung, D. Blaauw, D.M. Sylvester, K.D. Wise, A.M. Sastry, J. Power. Sources 170 (2007) 216–224. [3] P.H.L. Notten, F. Roozeboom, R.A.H. Niessen, L. Baggetto, Adv. Mater. 19 (2007) 4564–4567. [4] F.C. Liu, W.M. Liu, M.H. Zhan, Z.W. Fu, H. Li, Energy Environ. Sci. 4 (2011) 1261–1264. [5] J.M. Tarascon, M. Armand, Nature 414 (2001) 359–367. [6] P. Birke, W. Weppner, Electrochim. Acta 42 (1997) 3375–3384. [7] Y.S. Yoon, S.B. Cho, S.C. Nam, Electrochim. Acta 71 (2012) 86–91. [8] P. Vinatier, P. Vinatier, B. Pecquenard, A. Levasseur, H.J. Sohn, Solid State Ionics 160 (2003) 51–59. [9] X. Yu, J.B. Bates, G.E. Jellison Jr., F.X. Hart, J. Electrochem. Soc. 144 (1997) 524–532. [10] Y. Inaguma, C. Liquan, M. Itoh, T. Nakamura, T. Uchida, H. Ikuta, M. Wakihara, Solid State Commun. 86 (1993) 689–693. [11] T. Abe, M. Ohtsuka, F. Sagane, Y. Iriyama, Z. Ogumi, J. Electrochem. Soc. 151 (2004) A1950–A1953. [12] H. Aono, E. Sugimoto, Y. Sadaoka, N. Imanaka, G. Adachi, J. Electrochem. Soc. 137 (1990) 1023–1027. [13] J. Fu, Solid State Ionics 96 (1997) 195–200. [14] P. Knauth, Solid State Ionics 180 (2009) 911–916. [15] V. Thangadurai, H. Kaack, W. Weppner, J. Am. Ceram. Soc. 86 (2003) 437–440. [16] R. Murugan, V. Thangadurai, W. Weppner, Angew. Chem. Int. Ed. 46 (2007) 7778–7781. [17] F. Vaz, L. Rebouta, P. Goudeau, J.P. Riviere, E. Schaffer, G. Kleer, M. Bodmann, Thin Solid Films 402 (2002) 195–202. [18] M. Bromark, M. Larsson, P. Hedenqvist, M. Olsson, S. Hogmark, E. Bergmann, Surf. Eng. 10 (1994) 205–214. [19] G. Kleer, R. Kassner, E.M. Meyer, M.G. Schinker, W. Do¨ll, Surf. Coat.Technol. 54 (1992) 167–170. [20] R.W. Hoffman, Surf. Interface Anal. 3 (1981) 62–66. [21] M.A. Moske, P.S. Ho, D.J. Mikalsen, J.J. Cuomo, R. Rosenberg, J. Appl. Phys. 74 (1993) 1716–1724. [22] J.K. Ahn, S.G. Yoon, Electrochem.Solid State lett. 8 (2) (2005) A75–A78. [23] Y.-M. Chiang, D.R. Sadoway, Y.I. Jang, B. Huang, H. Wang, Electrochem. Solid-State Lett. 2 (1999) 107–110. [24] T. Ohzuku, T. Yanagawa, M. Kouguchi, A. Ueda, J. Power. Sources 68 (1997) 131–134. [25] S. Buta, D. Morgan, A. Van der Ven, M.K. Aydinol, G. Ceder, J. Electrochem. Soc. 146 (1999) 4335–4338. [26] C. Wolverton, Acta Mater. 49 (2001) 3129–3142. [27] G. Tan, F. Wu, L. Li, Y. Liu, R. Chen, J. Phys. Chem. C 116 (2012) 3817–3821. [28] H. Kitaura, A. Hayashi, T. Ohtomo, S. Hama, M. Tatsumisago, J. Mater. Chem. 21 (2011) 118–124. [29] E. Quartarone, P. Mustarelli, Chem. Soc. Rev. 40 (2011) 2525–2540.
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