Magnetron Sputtering Preparation of Nitrogen-Incorporated Lithium

Article pubs.acs.org/JPCC

Magnetron Sputtering Preparation of Nitrogen-Incorporated Lithium−Aluminum−Titanium Phosphate Based Thin Film Electrolytes for All-Solid-State Lithium Ion Batteries Guoqiang Tan,† Feng Wu,†,‡ Li Li,*,†,‡ Yadong Liu,† and Renjie Chen*,†,‡ †

School of Chemical Engineering and Environment, Beijing Institute of Technology, Beijing Key Laboratory of Environmental Science and Engineering, Beijing 100081, China ‡ National Development Center of High Technology Green Materials, Beijing 100081, China ABSTRACT: We report for the first time a new lithium ion conducting Li−Al−Ti−P−O−N thin film solid electrolyte for allsolid-state lithium ion batteries. It was prepared by radio frequency (RF) magnetron sputtering deposition using a NASICON structural Li−Al−Ti−P−O target in a N2 atmosphere at various temperatures. XRD and SEM test results showed that the thin film was composed of an amorphous structure and that its surface was smooth, dense, and homogeneous. FTIR and XPS analyses indicated that nitrogen atoms were actually incorporated into the Li−Al−Ti−P−O matrix framework. The substitution of nitrogen for oxygen in the thin film created more abundant crosslinking structures and decreased the electrostatic energy, which favored the higher mobility of lithium ions. A high Li ionic conductivity of 1.22 × 10−6 S/cm was obtained for the thin film deposited at room temperature. Moreover, the higher value of 1.22 × 10−5 S/cm for the thin film deposited at 500 °C indicated that some crystallites in the amorphous film might be beneficial in improving Li ionic conductivity. Therefore, different conductivity values are correlated with structural differences. The temperature dependence of the ionic conductivities fit the Arrhenius relation and the thin film deposited at 500 °C possessed of the lowest activation energy. Electrochemical analyses suggest that the high Li ionic conductivity is attributed to the reduced activation energy by the control of composition and structure. These properties make this thin film electrolyte a promising candidate material for use in all-solid-state thin film lithium ion batteries.

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INTRODUCTION In the past decade, portable electronic devices, smart cards, implantable medical devices, microelectromechanical systems (MEMS), and nanoelectromechanical systems (NEMS) have played an important and continuously increasing role in daily life.1−4 All of these devices need a power supply to serve as an energy storage and conversion device like a battery. Of the several battery technologies available, lithium ion batteries (LIBs) are considered the most promising because they provide the highest voltages and the largest energy storage densities.5 They have been used as power sources in various electronic devices and in environmentally friendly vehicles (fully electric and hybrid cars).6 However, conventional LIBs containing organic electrolytes have problems that limit their scalability, particularly in regard to safety, lifetime, and cost. All-solid-state LIBs containing nonflammable solid electrolytes offer a fundamental solution for the safety of conventional LIBs. Moreover, with their increased cycle life and energy density they provide diversity for battery design and facilitate battery miniaturization.7 Their main problem is a lower power density, and this is mostly attributed to the low Li ionic conductivity of the solid electrolytes and the high resistance to lithium ion transfer across the electrode/electrolyte interface.8 The search © 2012 American Chemical Society

for good solid electrolytes is a major goal in the development of all-solid-state LIBs. High-performance solid electrolytes with high Li ionic conductivity but negligible electronic conductivity, wide electrochemical stable window, good thermal stability, high resistance to shock and vibration, and excellent compatibility with electrodes are especially useful for high energy and power densities as well as for long-term stability.9 Solid electrolytes for all-solid-state LIBs comprise solid polymer electrolytes, inorganic lithium ion conductions, and hybrid organic−inorganic solid electrolytes. Solid polymer electrolytes are currently the most popular solid electrolytes, but their poor ambient-temperature conductivity and thermal stability need to be improved.10,11 Hybrid solid electrolytes based on polymers mixed with nanoscale inorganic fillers, such as Al2O3, SiO2, TiO2, and ZrO2, have been proved effective in enhancing the mechanical strength and ionic conductivity.12,13 However, the reason for the improvement in ionic conductivity upon the addition of nanosized ceramic particles is still debated, and even contradictory reports have been published. Inorganic Received: July 26, 2011 Revised: January 5, 2012 Published: January 9, 2012 3817

dx.doi.org/10.1021/jp207120s | J. Phys. Chem. C 2012, 116, 3817−3826

The Journal of Physical Chemistry C

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sputtering) was mixed in a porcelain crucible and melted at 700 °C for 2 h in an electrical furnace. The obtained ceramics were then ball-milled with dehydrated ethanol at a rotating speed of 420 rpm for 8 h to obtain a fine powder using a planetary ball mill (Fritsch Puluerisette 7, Germany). After ball milling, the powder was dried at 80 °C for 24 h. The Li−Al− Ti−P-O target was prepared using a conventional cold-pressing method. 22 g of ultima powder compound was pressed into a 60 mm diameter pellet, and the pellet was sintered at 900 °C for 5 h. The composition of the target was estimated by XPS to be Li1.62Al0.74Ti1.55P3O12.5. Deposition of the Thin Films. Li−Al−Ti−P−O−N thin films were prepared by RF magnetron sputtering using the Li− Al−Ti−P−O target under a high-purity N2 atmosphere with a specific pattern, as shown in Scheme 1. The distance between

lithium ion conduction as a true solid electrolyte in which only lithium ions are mobile can be divided into oxide or non-oxide (e.g., sulfides, halides, nitrides) by element and amorphous or crystalline by phase.14−17 In general, oxide materials are believed to be superior to non-oxide materials because of easy handling and mechanical, chemical, and electrochemical stability. To date, most of the discovered inorganic lithium ion conductions have had either high ionic conductivity or wide electrochemical window stability, but not both.18 Some oxides are excellent lithium ion conductors; for example, the NASICON structural material LiTi2(PO4)3 exhibits a relatively high conductivity of 2.0 × 10−6 S/cm at room temperature, and this gives rise to further investigations into conductivity optimization by chemical substitution and structural modifications. An increase in conductivity was observed when Ti4+ was partially substituted by various trivalent cations in Li1+xMxTi2−x(PO4)3 (M = Al3+, In3+, Ga3+, Sc3+) systems and/or when the P5+ ion was substituted by Si4+.19−21 Compared with crystalline electrolytes, amorphous electrolytes are widely investigated and applied in all-solid-state LIBs with a resultant denser microstructure, less grain boundary effect, and better chemical stability. Many amorphous electrolytes even exhibit higher ionic conductivity than their mother materials.22 Nitride glassy oxide-based electrolytes offer improved ionic conductivity and other properties, and these are attributed to the increased cross-linking structures and reduced electrostatic energy from the substitution of doubly −N= (Nd) and triply coordinated −N⟨ (Nt) nitrogen for bridging (BO) and nonbridging oxygen (NBO) ions in the glass network.23,24 To promote the use of electrolytes in practice, their properties and their fabrication processes are very important. Thin film fabrication can reduce the thickness of the electrolyte layer, resulting in a reduction in the internal resistance of devices. A large surface area is also necessary to achieve a large contact area for a large output current. Thin film electrolytes have been prepared using different kinds of vapor deposition techniques such as vacuum evaporation, radio frequency (RF) sputtering, and pulsed laser deposition. Compared with other fabrication technologies, RF magnetron sputtering can easily be used to fabricate large-scale and dense films with outstanding performance. Herein, we successfully prepared nitrogen-incorporated lithium−aluminum−titanium phosphate (Li−Al−Ti−P−O− N) thin film electrolytes using RF magnetron sputtering. This constitutes a new chemical composition with an amorphous glassy structure and predominantly ionic conduction. We report here for the first time a mixed effect of chemical substitutions and structural modifications in an amorphous Li1+xAlxTi2−x(PO4)3 system and its effect on structure and ionic conductivity. The high Li ionic conductivity together with an acceptable electrochemical stability window, good thermal stability, chemical stability under ambient conditions, environmental benign nature, readily available starting materials, and low cost suggest that this glassy Li−Al−Ti−P−O−N thin film is a promising electrolyte for all-solid-state LIBs.

Scheme 1. Schematic Representation of the RF Magnetron Sputtering System with the Li−Al−Ti−P−O Target

the target and the substrate was 6 cm. The base pressure was 1.0 × 10−5 Pa, which could guarantee a clean sputtering chamber. The working pressure was 2.0 Pa, the RF power was 80 W, and a typical deposition rate was estimated to be 200 nm/h. Before the film deposition the target was presputtered for 30 min. Thin films were deposited on several types of substrates such as stainless steel (SS), Si wafer, Cu foil, and Pt layer, as required for various characterization experiments. Thin films were deposited on Si wafer and SS sheet for SEM, XPS, FTIR, and XRD. Properties of ionic conductivity and stable electrochemical window were determined using blocking electrode cells stacked in the form of SS/Li−Al−Ti−P−O− N/SS and Pt/Li−Al−Ti−P−O−N/Pt, respectively. A sandwich structure consisting of a layer of SS, a layer of thin film, and another layer of SS deposited sequentially on a Si wafer substrate was prepared for the electrochemical impedance spectroscopy test. During the deposition substrates were heattreated at various temperatures: room temperature, 200 °C, 300 °C, 400 °C, and 500 °C. Characterization of Properties. After we prepared thin films, we protected them in an argon atmosphere and measured their properties as quickly as possible using various characterization techniques. The crystal structure characterization was carried out by X-ray diffraction (XRD, Japan) using a Rigaku Xray diffractometer with Cu Kα radiation at a scan rate of 8°/ min. Chemical bonds in thin films were examined by Fourier transform infrared spectroscopy (FTIR, Nicolet 6700). The film morphology and thickness were determined by scanning electron microscopy (SEM, JSM-35C, Japan), and energy dispersive X-ray detector analysis (EDX) was also carried out

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EXPERIMENTAL SECTION Preparation of the Target. Solid electrolyte ceramics were synthesized using a conventional melt-quenching method.19 Reagent grade Li2CO3, Al2O3, TiO2, and NH4H2PO4 were used as starting materials. A 42 g batch consisting of 14Li2O− 9Al2O3−38TiO2−39P2O5 (in mol %, with a certain excess of Li2CO3 to compensate for Li loss during heating and 3818



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on this instrument. The chemical composition of the thin films was determined using an X-ray photoelectron spectrometer (XPS, PHI Quantera, Japan). Electrochemical impedance spectroscopy (EIS) was carried out to determine the ionic conductivity using blocking electrode cells in the ac frequency range of 1 Hz−100 kHz with a Zahner IM6e (Germany). The activation energy of Li ionic conduction was calculated using the temperature dependence of the conductivity from 20−120 °C. The Li+ transference number was measured by the dc polarization technology using a symmetrical Li/Li−Al−Ti−P− O−N/Li cell, and the electronic conductivity was evaluated by the Hebb−Wagner polarization method25 using a (−)Li/Li− Al−Ti−P−O−N/Au(+) cell at room temperature. The electrochemical window stability in the range 0−5 V was examined by linear sweep voltammetry (LSV) at a scan rate of 0.1 mV/s at room temperature.

in shape and size, and the size is ranged from 100 to 400 nm. The bulk film is dense and homogeneous, as shown by the cross section of the film (Figure 1d). The composition of thin films was approximately estimated by EDX spectroscopy in Figure 2, and this was further confirmed by the XPS result. For our films were well protected in argon atmosphere before characterization tests, the element content of thin films was almost reliable. The surface morphology of the film without grain boundaries also indicates that our RF sputtering is useful in the formation of the glassy amorphous film, and this is determined by X-ray diffraction. Figure 3a shows the XRD pattern of the Li−Al− Ti−P−O target, which clearly contains sharp diffraction peaks because of the lithium analogue of the NASICON-type structure (i.e., LiTi2(PO4)3) without any other impure phase. It suggests that the Al3+ ion is incorporated into the structure of LiTi2(PO4)3 and forms the Li1+xAlxTi2−x(PO4)3 phase. The diffraction pattern is almost consistent with that of the reported Li2O−Al2O3−TiO2−P2O5 solid ceramic.19 For thin films, the XRD pattern of the film deposited at room temperature is devoid of any dominant peaks and has an almost flat diffraction pattern except for that on the stainless steel substrate (i.e., 44° for SS(111), 51° for SS(100), and 75° for SS(220)), suggesting that the film is an amorphous glassy. Upon deposition at 500 °C the XRD pattern of the film contains a broad peak between 15° and 30° and two visible but weak diffraction peaks at 20.8° and 24.5°. This result indicates that some crystallites form and grow up when the glass is deposited at 500 °C, and the film contains Li1+xAlxTi2−x(PO4)3 as a major crystalline phase. It also suggests that this thin film has an amorphous structure that coexists with small crystallites. These results support the assertion that our thin film electrolytes are mainly composed of an amorphous glassy structure. Infrared Spectroscopy Analyses. FT-IR spectroscopy provides considerable insight into the chemical bonding structure of the target and thin films. From the mid-IR region (4000 to 400 cm−1) strong absorptions were observed and arose from the framework structure Li, Al, Ti, P, O, and N species. Figure 4 shows FTIR absorption spectra of both the crystalline target and glassy thin films. The target exhibits a typical IR spectrum for a well-crystallized NASICON-type Li1+xAlxTi2−x(PO4)3 framework. The spectra are dominated by the internal vibrations of the PO43− units which involve the displacement of oxygen atoms of the tetrahedral PO43− anions and give frequencies closely related to those of the free molecule. The bands at 460, 500, 583, and 634 cm−1 involve the symmetric and antisymmetric bending mode of the O−P−O bands with a small contribution from the P vibration. The highfrequency bands observed at 940 and 1020 cm−1 are the manifestation of a correlation field effect because of the coupling of the PO4 vibrators in the unit cell. The lantern units present in the NASICON phase give rise to infrared bands within 1150−1250 cm−1, and this is attributed to the stretching vibrations of the terminal PO3 units. All these spectral features are fingerprints for NASICON-type structures.26−28 The spectra is dominated by vibrational features due to phosphate ions, and transition metal ions are also present within 400−700 cm−1. The glass spectrum is approximately the envelope of the crystallized phosphate spectrum. For the Li−Al−Ti−P−O glassy film the broad band between 900 and 1140 cm−1 involves the symmetric and antisymmetric stretching mode of the P−O

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RESULTS AND DISCUSSION Morphology and Structure of the Thin Film. The surface morphology and cross-section thickness of Li−Al−Ti− P−O−N thin films deposited on Si-wafer substrates at room temperature and 500 °C are shown in Figure 1. From Figure 1a,

Figure 1. SEM images of the surface and cross section of Li−Al−Ti− P−O−N thin films deposited at room temperature (a, b) and 500 °C (c, d).

the surface of the thin film deposited at room temperature is obviously smooth, dense, uniform, and without pinholes, cracks, or aggregates of particles. This suggests that the thin film is homogeneous and has a flat surface morphology. The smooth surface enables the thin film to decrease the contact resistance between the thin film and electrodes, which is very important to avoid shortcuts and safety problems. The uniform surface without any grain boundaries also illustrates the amorphous structure. The thickness of the thin film measured in the cross-section direction by SEM (Figure 1b) was about 2.1 μm, and this can be controlled by adjusting the sputtering rate and the sputtering time. It is noteworthy that the surface of the thin film deposited at 500 °C is somewhat nonuniform (Figure 1c). A few grain boundaries may be formed as a result of small crystals in the glass at high temperature, and this is confirmed by XRD. During the heat treatment, the temperature is near the crystallization temperature of the Li−Al−Ti−P−O− N; small crystals form and grow up in the film. They are varied 3819



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Figure 2. EDX of Li−Al−Ti−P−O−N thin films deposited at room temperature (a) and 500 °C (b).

similar spectrum feature contrast to the Li−Al−Ti−P−O film. However, in Li−Al−Ti−P−O−N film, the band centered at 960 cm−1 involving P−O−P stretching mode is obviously weakened while the band centered at 1050 cm−1 representing terminal oxygen P−O− stretching mode is increased. That is to say, the ratio NBO/BO increases when nitrogen is incorporated. This indicates that the percent of P−O groups with the NBO increase significantly after nitridation. During the nitridation, the BO (P−O−P) is mainly substituted by nitrogen and converts to the P−N= and P−N⟨; meanwhile, some NBO sites are formed for more cross-linked structure. Another reason might be the homogenization of Li ions during sputtering and heat treatment.29 Some of Li in the form of an amorphous oxide will react with other oxides during the sputtering and heat treatment, such as the addition of Li2O can break up PO bridges, creating NBO. Also, there is a phase reaction between amorphous phases and crystalline phases to form another amorphous phases. These homogenizations will distribute the Li ions and incorporate them into the glass network, which will produce a large number of NBO. From these results we conclude that Li ionic conductivity increased upon a significant increase of NBO sites. FT-IR failed to detect obvious potential differences in the microstructures of the thin films that were deposited at different temperatures. Further microstructural analysis is needed on thin films to identify various atomic-scale microstructures such as bonding configurations, structural relaxation, point defects, and bond-angle distributions. All these parameters may control lithium ion conduction in thin films for enhancement or retardation. X-ray Photoelectron Spectrometry. To further characterize the chemical nature of the glassy film, the target and thin films were characterized by XPS. Figure 5 shows a set of parallel XPS spectra of the target and thin films. All the XPS spectra consist of peaks assigned to Al, Ti, P, O, and C. C 1s was used as an internal reference (284.80 eV). An obvious peak around 398.00 eV was observed in thin films, but not the target, and this is attributed to the nitrogen substituted oxygen during RF sputtering. Another obvious decrease in the peak intensity of P in thin films upon increasing the deposition temperature was also consistent with the EDX results. In thin films, the surface element concentrations represented a decrease or increase

Figure 3. XRD patterns of the Li−Al−Ti−P−O target (a) and Li− Al−Ti−P−O−N thin films deposited at room temperature (b) and 500 °C (c).

Figure 4. FTIR spectra of the Li−Al−Ti−P−O target (a), the Li−Al− Ti−P−O film (b), and the Li−Al−Ti−P−O−N film (c) both deposited at room temperature.

groups.29 The broad band between 550 and 700 cm−1 is attributed to the symmetric P−O−P bonding that occurs in glassy structures.27 The Li−Al−Ti−P−O−N film shows a 3820



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eV). This value suggests a monovalent chemical state for the Li ion.30 An obvious reduced Li 1s XPS peak intensity was observed in thin films compared with that of the target, and this was attributed to Li loss during sputtering and heating. Figure 6b gives the XPS spectra of Al 2p, and the increased peak intensity in the XPS of Al 2p corresponds to an increased concentration of Al in thin films. Figure 6c shows the Ti 2p core level spectrum of all samples. The target reveals a Ti4+2p3/2 peak at 458.88 eV with a full width at half-maximum (fwhm) of 1.20 eV. In addition, a small shoulder appears on the high binding energy side, and it is separated from the main peak by about 5.57 eV; we ascribe this to Ti4+2p1/2. However, in thin films the Ti4+2p3/2 and Ti4+2p1/2 binding energies shift slightly to lower binding energies. In addition, two small shoulders appear on the low binding energy sides of the respective Ti4+2p3/2 and Ti4+2p1/2 signals, and they are separated from the main peak by 0.96 and 0.92 eV, respectively. We believe that they are attributable to the Ti3+2p3/2 and Ti3+2p1/2 binding energies, and one reason for this phenomenon is the reduction of tetravalent Ti4+ to trivalent Ti3+ ions, meanwhile generating high Li+ concentrations.30 Guillemot et al.31 reported that the binding energy shift of 1.40 eV from the main peak to the low binding energy side is a characteristic of Ti3+ defects. The binding energy of Ti3+ observed in this study is higher than that reported for Ti2O3, and this is due to the high ionic character of Ti bonds in phosphate glasses and confirms the localization of Ti4+/Ti3+ near the PO bond. These results may be explained by the substitution of nitrogen atoms into the Ti−O bands. Titanium Ti4+ coordinated by four oxygen acts as glass-former oxide, while Ti3+ is a modifier of the glass structure.32 Figure 6d,e shows the P 2p and O 1s XPS spectra. In thin films, the P 2p binding energies shift to higher values and the O 1s binding energies shift to lower values compared with those of the target. These results indicate that the concentration of BO in these films is obviously reduced after nitride. This shows that the nitrogen partially substitutes the oxygen, especially BO, in the glass. In the N 1s spectrum of thin films (Figure 6f), an asymmetric peak was observed around 399.00 eV, which decomposes into two Gaussian−Lorentzian mixed components located at 397.80 and 399.20 eV. These two binding energy values can be attributed to two different bonding states of nitrogen atoms.33 The nitrogen bonds of doubly coordinated nitrogen P−N= (Nd) are at 397.80 eV and triply coordinated nitrogen P−N⟨ (Nt) are at 399.20 eV. A shift of the peak toward higher energy could be interpreted as a slight increase in the Nt content in the film with an increase in the heat-treatment temperature. In the N 1s spectrum of the film (500 °C), the larger area of the Nt structural unit indicates its relatively higher content in the structure. The Nt structural units provide more abundant crosslinking structures and more free space for ion transfer, and their contribution to a drop in electrostatic energy from the Nt structure unit is more than that from a Nd structure unit.34 In summary, the Nt cross-linked structure unit is more responsible than the Nd structure unit for the increase in Li ion mobility. It is noteworthy that there is an additional weak feature in the N 1s region at 403.5 eV for the thin film deposited at 500 °C. This is believed to be due to the 1s state of nitrogen in different chemical environments,35 and we infer that it is the N 1s characteristic spectrum of the polycrystalline structure. This result also indicates a partial crystal of the film electrolyte that was deposited at 500 °C.

Figure 5. XPS survey spectrum of the target and thin films deposited at room temperature and 500 °C.

trend upon increasing the deposition temperature. The Li, P, and N element concentrations all showed a slightly decrease trend because of a loss for these elements during the heating treatment. Contrarily, the Al and Ti element concentrations were increased. The Li content in the electrolyte is largely related to the Li+ ionic conductivity, so we added a certain excess of Li2CO3 for the target preparation to compensate for Li loss during sputtering and heating. Figure 6a illustrates the XPS spectra of the Li 1s core level of the target and thin films. The binding energies of the Li 1s electrons in all samples are present at the same value (55.14

Figure 6. XPS survey spectra of various elements in the target and thin films. Black line: target; blue line: thin film (25 °C); red line: thin film (500 °C). Inset: Gaussian−Lorentzian curves of the thin film (500 °C). 3821



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Figure 7. Electrochemical impedance spectroscopy of the Li−Al−Ti−P−O−N thin film deposited at room temperature was measured at various temperatures.

Impedance Properties. Electrochemical impedance spectra (EIS) measurements of thin films were performed as a function of temperature from 20 to 120 °C. Nyquist plots for the thin film deposited at room temperature and measured at different temperatures are shown in Figure 7. The Nyquist plots show semicircular arcs at high frequencies and spikes at low frequencies. The single semicircle, which is caused by the bulk Li−Al−Ti−P−O−N thin film, is simulated using an equivalent circuit consisting of a geometrical capacitance in parallel with a bulk resistance. The bulk resistances are used to calculate the conductivities of the film electrolytes at different temperatures. It is noteworthy that the diameter of the semicircles decreases with an increase in the measurement temperature. This indicates that the conductivity of the film electrolyte increases with an increase in the measure temperature. A low-frequency sloping spike of the impedance spectrum is assigned to the interfaces between the thin film and the adjacent SS electrodes. In the cases where spikes are observed at low frequencies a constant phase element (CPE) is added in series to the parallel circuit. The electrochemical impedance spectrum of the film electrolyte is characteristic of a thin film conducting dielectric with bulk relaxation processes, which are sandwiched between the blocking contacts. The equivalent circuit is simulated using the Cole−Cole model. The ionic conductivity σ of the Li−Al−Ti−P−O−N films can be calculated according to eq 1. The thin film deposited at 500 °C gives the largest ionic conductivity of 1.22 × 10−5 S/cm at room temperature while the Li−Al−Ti−P−O−N thin film

deposited at room temperature also gives a high ionic conductivity of 1.22 × 10−6 S/cm at room temperature, which is similar to that of the well-known LiPON.36 This result suggests that the Li−Al−Ti−P−O−N thin film solid electrolytes with high ionic conductivity are good enough for application in all-solid-state thin film LIBs

σ = d /(RS)

(1)

where d is the thickness of the film electrolyte, S is electrode/ electrolyte interface area, and R is the resistance of the film electrolyte that is determined from the measured impedance by selecting the value of Z′ when −Z″ goes through a local minimum in the electrochemical impedance spectra.37 The temperature dependence of conductivity for the thin film electrolytes is shown in Figure 8. The conductivity σ follows the Arrhenius equation given as eq 2. The activation energy (Eσ) for lithium ion conduction in the solid electrolyte is also calculated using eq 2 and the measured conductivity at various temperatures

σT = Ae−Eσ /(KBT )

(2)

where A is the pre-exponential factor, KB is the Boltzmann constant, and T is the absolute temperature. Figure 8a shows the conductivities of the film electrolytes deposited at different heat-treatment temperatures as a function of the measured temperature. The conductivity increases rapidly with an increase in the measured temperature as well 3822



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500 °C, and the corresponding activation energy of conduction Eσ is about 0.44 eV, which is smaller than that of LiPON (0.56 eV),38 indicating higher lithium ion mobility for Li−Al−Ti−P− O−N than that for LiPON. This proves that the thin film electrolytes are good lithium ion conductors. Additional insight into the conduction mechanism is expected from the measurements of the Li+ transference number and the electronic conductivity. The typical dc polarization curve of the film electrolyte is shown in Figure 9.

Figure 9. The dc polarization curve of the Li−Al−Ti−P−O−N thin film deposited at room temperature.

The external current (Iext) largely charged on applying a polarization potential 10 mV. The initial polarization current (Ii, t = 100 min) and the steady state current (Is, t = 300 min) during polarization were measured. The Li+ transference number was calculated to be 0.9999, which is very approximate to 1; this indicates Li+ conduction is absolutely dominant in the electrolyte. The electronic conductivity of the film electrolyte, which controls the self-discharge of a lithium ion battery and therefore is of great interesting, was also measured with a Hebb−Wagner type cell arrangement with one reversible and one blocking electrode. Figure 10 shows the potential

Figure 8. (a) Conductivities of the Li−Al−Ti−P−O−N thin films prepared at different deposition temperatures and measured at various temperatures. (b) Arrhenius plot of the ionic conductivity of Li−Al− Ti−P−O−N thin films deposited at various temperatures.

as an increase in the deposition temperature. The Arrhenius plot indicates similar activation energies for the ionic conduction of the thin films, as shown in Figure 8b. The curves tend to have low activation energies (0.63 to 0.44 eV with an increase in the deposition temperature), indicating high lithium ion mobility. The ionic conductivities and active energies of conduction for the thin films deposited at different temperatures are shown in Table 1, and it is obvious that the Table 1. Lithium Ion Conductivities and Active Energies of the Film Samples Deposited at Various Temperatures ionic conductivity, σ/10−6 S cm−1 film samples

deposition temp (°C)

30 °C

50 °C

75 °C

100 °C

active energy, E(eV)

A B C D E

25 200 300 400 500

1.22 2.76 5.40 8.59 12.24

5.06 9.92 16.92 23.40 30.91

22.82 39.70 56.43 71.12 89.50

91.73 133.04 172.10 202.13 243.26

0.63 0.56 0.51 0.47 0.44

Figure 10. Partial electronic conductivity of the Li−Al−Ti−P−O−N thin film (deposited at room temperature) vs Hebb−Wagner cell voltage at room temperature.

dependence of electronic conductivities measured in the range of typical battery operating voltage. It is noticeable that the electronic conductivity increases with increasing the polarization voltage. As described in detail in the literature, the increasing voltage leads to a corresponding decrease of the

conductivity of the thin films increase with an increase in the deposition temperature. The room-temperature conductivity can reach 1.22 × 10−5 S/cm for the film electrolyte deposited at 3823



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chemical potential of Li at the Au electrode.39 This changes the electron and hole concentrations in the Li−Al−Ti−P−O−N electrolyte. Therefore, the electronic partial conductivity increases with increasing polarization voltages over the potential range of 2.5−4.5 V. The electronic conductivity takes values in the range from 7.53 × 10−12 to 4.92 × 10−11 S/ cm. Compared with the ionic conductivity of 1.22 × 10−6 S/cm, the electronic conductivity is negligible. This clearly shows that the Li−Al−Ti−P−O−N electrolyte is an almost exclusive Li ion conducting material having great potential as an electrolyte for LIBs. Electrochemical Stability. Electrochemical stability is also an essential criterion for a solid electrolyte. It was evaluated by linear sweep voltammetry (LSV) for the Pt/Li−Al−Ti−P−O− N/Pt cell up to 5.0 V, as shown in Figure 11. The irreversible

reason why the electrochemical properties of nitride films are superior to those of non-nitride films may be as follows: According to the equation σ = charge density × mobility, the ionic conductivity is determined by the product of charge density and mobility.41 The Li+ content of nitride films slightly decreases for the loss of Li2O during the deposition and heat treatment. The main reason for the increase in ionic conductivity is related to the ionic mobility. During nitridation, the formation of P−N bonds replace P−O bonds, leading to a more reticulated anionic network for the faster ionic mobility. The oxynitride glass network is then formed by PO4 and the new PO3N and PO2N2 tetrahedra, which increase the crosslinking density. These cross-linking structures create more conduction paths with a low activation energy and give rise to a notable modification of glass properties.42 The nitrogen substitution of oxygen reduces the highly polar bonding nature of oxygen, which in return provides weaker polar bonding between P and N, resulting in the enhanced mobility of lithium ions. Moreover, the substitution of P−O bonds with the more covalent P−N bonds decreases the electrostatic energy, which is helpful for Li ionic conductivity. Besides, in glass networks, NBO sites offer hopping sites for ionic conduction where Li ions easily jump into or out due to relatively weak bonding or shallow energy well. The formation of NBO is also helpful in creating a relatively open network structure with a large free volume for ions transfer.43 The ratio NBO/BO determined by XPS is approximately 7/3 for the film, and this value slightly increases under heat treatment. It is consistent with the IR test results. Therefore, the increased ionic conductivity of films is partly attributable to abundant NBO sites for Li+ transfer. Also, Ti3+/Al3+ chemical substitutions act as a modifier of the glass structure, resulting in more abundant NBO sites. Figure 12 shows the partial structure of the Li−Al−Ti−P− O−N. Both BO and NBO are substituted by Nt and Nd, which make up the chains of the glass network. The Nt/Nd ratio determined from their respective areas in the spectrum is about 3 for the film (500 °C). It has been claimed that a high Nt/Nd ratio would result in a high ionic conductivity of oxynitride phosphate glasses.44 Compared with Nd structural unit, the formation of Nt structural unit will remove more P−O bonds (BO). Consequently, the contribution to the drop of electrostatic energy from Nt structural unit is more than from Nd structural unit. It gives evidence that Nt cross-linking structural unit is more responsible than Nd structural unit for the increase in Li+ mobility. The nitrogen atoms also stabilize the mixed P(O, N)4 tetrahedra through the delocalization of πbonds.45 This is beneficial in improving the chemical and electrochemical durability of the films.

Figure 11. Linear sweep voltammogram for the blocking electrode cell with the Li−Al−Ti−P−O−N thin film deposited at room temperature.

onset of the current can be defined as the electrolyte breakdown voltage. A slight current charge occurred above 3.8 V, which is higher than that of the Li−Al−Ti−P−O thin film of 3.5 V, as reported in previous work.40 This value is also higher than that of the LiPON thin film at 3.0 V.38 This suggests that the Li−Al−Ti−P−O−N thin film electrolyte may have slightly improved stability compared with the Li−Al−Ti− P−O thin film and that it remains stable at high potentials of more than 3.0 V where lithium batteries usually operate. This thin film electrolyte is, therefore, predicted to be a promising candidate for all-solid-state thin film LIBs. Mechanism Analysis. An obvious improvement in the electrochemical properties of films has been confirmed, and the

Figure 12. Partial structure of the nitride Li−Al−Ti−P−O−N thin film with the incorporation of −N⟨ and −N= resulting from the reaction of two adjacent phosphate chains with nitrogen. 3824



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The improved ionic conduction is also related to the crystalline structure of the films. The film with partial polycrystalline structure shows higher conductivity than the full amorphous film. It means that the high conductivity of films may also be attributable to the mixed structural effect of a partial polycrystalline glassy structure.46 From the above analysis, the combined approach of the nitrogen incorporation and mixed structure effect is believed to be effective in improving the electrochemical properties of a glassy film electrolyte. Further studies are to be carried out to determine the microstructure and thoroughly investigate the transport mechanism.

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CONCLUSION

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AUTHOR INFORMATION

REFERENCES

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For the first time, Li−Al−Ti−P−O−N thin film electrolytes were successfully obtained by RF magnetron sputtering using the NASCON-type structural target Li−Al−Ti−P−O. From the XRD pattern, the film deposited at room temperature is amorphous, and it appears to become partially polycrystalline upon increasing the deposition temperature to 500 °C. SEM of thin films shows a smooth surface without cracks and pits. This is important for film electrolytes to avoid short-circuiting in the battery application. The FTIR spectra of thin films do not show sharp peaks but do show much broader peaks compared to that of the crystalline target and are consistent with the thin films being amorphous. The FTIR spectra also show the structural and compositional consistency in thin films. XPS analysis confirms that a more cross-linked structure resulted from the nitrogen substitution of oxygen, especially the bridging oxygen in the films. Enhanced ionic conductivity appears to result from the more cross-linked structure by the creation of doubly and triply coordinated nitrogen in phosphate network and the reduced electrostatic energy by the mixed anion effect. The ionic conductivity of the films increases with the deposition temperature, and the highest room temperature ionic conductivity of 1.22 × 10−5 S/cm with the lowest activation energy of 0.44 eV was obtained for the film deposited at 500 °C. We also ascribe this to a mixed structural effect of partial polycrystalline and amorphous glassy electrolyte. Moreover, the incorporated nitrogen atoms stabilize the mixed P(O,N)4 tetrahedra by the delocalization of π-bonds to improve the electrochemical stability of the thin films. In conclusion, a mixed effect of chemical substitution and structural modification gives a great enhancement in the electrochemical properties of thin films. We propose that this thin film electrolyte may be a promising inorganic solid electrolyte for all-solid-state thin film LIBs.

Corresponding Author

*E-mail: [email protected] (L.L.); [email protected] (R.C.).

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Article

ACKNOWLEDGMENTS

This work was supported by the National Key Program for Basic Research of China (No. 2009CB220100), the International S&T Cooperation Program of China (2010DFB63370), the National Science Foundation of China (NSFC, 20803003), New Century Educational Talents Plan of Chinese Education Ministry (NCET-10-0038), and Beijing Novel Program (2010B018). 3825



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Magnetron Sputtering Preparation of Nitrogen-Incorporated Lithium

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