Molybdenum-Suboxide Thin Films as Anode Layers in Planar Lithium Microbatteries

In this paper, we investigate the effects of operational conditions on structural, electronic and electrochemical properties on molybdenum suboxides (MoO3-δ) thin films. The films are prepared using pulsed-laser deposition by varying the deposition temperature (Ts), laser fluence (Φ), the partial oxygen pressure (PO2) and annealing temperature (Ta). We find that three classes of samples are obtained with different degrees of stoichiometric deviation without post-treatment: (i) amorphous MoO3-δ (δ < 0.05) (ii) nearly-stoichiometric samples (δ ≈ 0) and (iii) suboxides MoO3-δ (δ > 0.05). The suboxide films 0.05 ≤ δ ≤ 0.25 deposited on Au/Ti/SiO2/flexible-Si substrates with appropriate processing conditions show high electrochemical performance as an anode layer for lithium planar microbatteries. In the realm of simple synthesis, the MoO3-δ film deposited at 450 °C under oxygen pressure of 13 Pa is a mixture of α-MoO3 and Mo8O23 phases (15:85). The electrochemical test of the 0.15MoO3-0.85Mo8O23 film shows a specific capacity of 484 µAh cm−2 µm−1 after 100 cycles of charge-discharge at a constant current of 0.5 A cm−2 in the potential range 3.0-0.05 V.


Introduction
Thin-film technology opens the way for the development of micro-electrochemical power sources powering miniaturized devices such as credit cards, chip units, medical implants and devices, stand-alone sensors, etc. It is also useful for understanding the intrinsic properties of the active materials of lithium batteries (cathode, solid electrolyte, and anode) free of carbonaceous additive and polymeric binder [1]. Thin film electrodes have the advantage of shortening the ionic and electronic pathways, which enhances high-rate cycle ability.
Molybdenum oxides (MoO 3 and its MoO 3−δ suboxides with stoichiometric deviation δ < 1) belong to the class of materials that offer a tunability of their intrinsic electronic properties from wide bandgap semiconductor MoO 3 to metallic MoO 2 [2]. The varieties of oxidation states from Mo 6+ to Mo 4+ in MoO 3−δ films makes them highly attractive in the field of energy storage and conversion and they have found application in all-solid state thin film microbatteries (TFµBs) [3,4], electrochemical  [20]. Copyright 2016 American Chemical Society. Figure 2 shows a schematic diagram of the electronic states of molybdenum oxides as a function of the oxygen vacancies. The electronic character varies from insulator (MoO3) to semiconductor (MoO3-δ) and finally to metal-like (Mo9O26), in agreement with conductivity measurements [21]. The electron distribution in pure MoO3 follows the ionic model (Mo 6+ and O 2− ) corresponding to the Mo 4d 0 configuration. For non-stoichiometric MoO3 oxides, extended Mo 4d states, which then lie in the bandgap as gap states due to the occurrence of Mo 5+ and Mo 4+ ions, are filled by electrons donated from oxygen vacancies. The formation of oxygen deficiency (termed also as sub-stoichiometric) not only produces an increase of the electrical conductivity σe [23] owing to the additional gap states, but also increases the surface energy of the particles and promotes electrochemical reactions [21,24].  [20]. Copyright 2016 American Chemical Society. Figure 2 shows a schematic diagram of the electronic states of molybdenum oxides as a function of the oxygen vacancies. The electronic character varies from insulator (MoO 3 ) to semiconductor (MoO 3−δ ) and finally to metal-like (Mo 9 O 26 ), in agreement with conductivity measurements [21]. The electron distribution in pure MoO 3 follows the ionic model (Mo 6+ and O 2− ) corresponding to the Mo 4d 0 configuration. For non-stoichiometric MoO 3 oxides, extended Mo 4d states, which then lie in the bandgap as gap states due to the occurrence of Mo 5+ and Mo 4+ ions, are filled by electrons donated from oxygen vacancies. The formation of oxygen deficiency (termed also as sub-stoichiometric) not only produces an increase of the electrical conductivity σ e [23] owing to the additional gap states, but also increases the surface energy of the particles and promotes electrochemical reactions [21,24]. Magnetic susceptibility measurements showed that MoO 3 , MoO 2 and Mo n O 3n−1 shear suboxides in between are all feebly paramagnetic.
The MoO 3−δ suboxide phases (with 0 < δ < 0.25) exhibit Li-insertion capacity much higher than that of stoichiometric MoO 3 phase, which justifies the efforts to fabricate them under the form of thin films.
For the development of these technologies, oxygen vacancies are easily generated in the MoO 3−δ materials prepared in a thin film architecture by optimizing the deposition conditions [25]. Like stoichiometric MoO 3 , sub-stoichiometric MoO 3−δ thin films can be obtained by evaporation in a vacuum at relatively low temperature (T < 500 • C) or in reducing atmosphere. However, few works explore the effect of oxygen deficiency and electrical conduction enhancement on energy storage properties [26,27].
It is well known that PLD, based on the process of the transportation of material (laser ablation), is a powerful technique to fabricate multicomponent oxide dense films with high purity. Generally, the stoichiometry can be easily tuned by a control of growth rate (R g ), laser fluence (Φ), substrate temperature (T s ), oxygen partial pressure (P O 2 ) and morphology of the substrate [55]. Two strategies are used to fabricate well-crystallized MoO 3−δ thin films: (i) the deposition is carried out in an O 2 atmosphere and at elevated T s , which provides Mo oxidation inside the chamber during the growth process [53], (ii) the film oxidation is prepared in a vacuum with post-annealing at a moderate temperature (T a ) without the use of reactive O 2 and substrate heating [50,51]. In view of the numerous studies of MoO 3−δ thin films prepared using the PLD technique, it is important to compare the structural properties of films obtained under different experimental conditions [56]. Hussain et al. optimized the deposition conditions (T s = 200 • C, P O2 = 13 Pa) for crystalline α−MoO 3 , while at an oxygen partial pressure less than 13, Pa sub-stoichiometric films with α−β mixed phases were found [43]. Camacho-Lopez investigated the structural transition in PLD films; the as-deposited MoO 1.99 is oxidized to MoO 2.77 by heat treatment at 200 • C for 14 h in air [43]. Torres et al. [46] studied MoO 3 films fabricated using a thermal laser deposition method with a CO 2 laser (λ = 10.6 µm) at an intensity of 0.73 W mm −2 under oxygen pressure of 0.5 Pa. For a substrate temperature T s ≥ 350 • C, the XRD patterns displayed both (0k0) and (0kl) orientations representing α-β mixed MoO 3 phases. Robinson-Azariah et al. studied the effect of the repetition rate (R rl ) on the composition of MoO x PLD films deposited at a fluence of 150 mJ cm −2 on fluorine-doped tin-oxide-coated glass substrate kept at 25 • C under O 2 gas at a pressure of 5 Pa [49]. While XRD patterns exhibited the orthorhombic α-MoO 3 phase for all films, the XPS studies showed the presence of Mo 5+ oxidation state giving oxygen-deficient films MoO 2.98 and MoO 2.75 at R rl of 2 and 10 Hz, respectively. PLD MoO 3−δ suboxide films (~100 nm thick) with different compositions were deposited on fused silica slides using two excimer laser sources, XeF (λ ex = 351 nm, Φ = 85 mJ) and KrF (λ ex = 248 nm, Φ = 200 mJ), without introduction of oxygen in the chamber and subjected to post-annealing at a temperature, T a , in the range of 300-500 • C for 4 h in air [50]. The as-prepared films were amorphous (dark color) and started to be partially crystallized at T a ≈ 400 • C; the chemical surface state studied using XPS exhibited an O/Mo ratio of 2.95, which decreased to 2.78 upon heating at 500 • C. Subsequent annealing at 300 • C in vacuum reduced the deviation from stoichiometry to MoO 2.65 [57].
In this work, we investigate the growth of molybdenum-oxide thin films (oxygen deficient MoO 3−δ and Mo n O 3n-1 suboxides) deposited on Au/Ti/SiO 2 /flexible-Si substrates using a PLD technique with appropriate processing conditions. Relationships are established between structure, texture, cationic environment and deposition conditions such as substate temperature, ablation power and oxygen partial pressure. The electrochemical performance of as-prepared and heat-treated films are examined as anode in TFµBs. Tuning the oxygen vacancies in MoO 3−δ suboxide thin films appears to be one of the best engineering strategies to obtain an extended cycling stability and a decrease of the overpotential. The introduction of oxygen vacancies accelerates the electron transport, which can play a positive role in boosting the Li + diffusivity. To the best of our knowledge, this is the first report on the correlation between the nanoscale structure of MoO 3−δ thin films and their electrochemical properties. The reversible lithium insertion/deinsertion reaction at ambient temperature makes them a candidate for a negative electrode in lithium microbatteries.
This paper is organized as follows. Section 2 presents the experimental techniques (substrate preparation, growth and characterization). In Section 3, detailed results describe the fundamental properties of films including structure, morphology and composition. The next section, Section 4, reports the electrochemical properties of MoO 3−δ thin films (charge-discharge profiles and Li + ions kinetics). Finally, a general discussion is provided in Section 5, comparing the properties of as-prepared PLD MoO 3−δ films with those reported in the literature.

Preparation of Au/Ti/SiO 2 /flexible-Si Substrates
Au/Ti/SiO 2 /Si stacking multilayers were used as substrates, in which the Si chip is the flexible mechanical support (12 µm thick, ULTRATHIN, Virginia Semiconductors Inc., Fredericksburg, VI, USA), the SiO 2 film (100 nm thick) has a unique role of electrical insulation between the current collector (Au/Ti) and Si wafer. The gold film acts as current collector, while the titanium layer enhances the mechanical adhesion to the SiO 2 layer. The multilayered substrates were fabricated according to the following sequence. (i) Silicon chips were treated according the RCA-type cleaning procedure [58]. (ii) The SiO 2 layer (about 100 nm thick) was produced by cleaning the Si chips in H 2 SO 4 /H 2 O 2 solution then rinsed in hydrogen fluoride. (iii) the current collector was formed via metallization with a 20-nm thick Ti layer deposited using radio-frequency (RF) magnetron sputtering then metallized with a 20-nm thick Au film. A scheme representing the stack of the different layers is shown in Figure 3a.

Film Characterization
The phase and structure of as-prepared MoO3-δ films was characterized using X-ray diffraction (XRD), Raman scattering (RS) and Fourier transform infrared (FTIR) spectroscopy. XRD patterns were investigated (X-ray diffractometer 3003 TT, Siefert, Leuven, Belgium) using a CuKα radiation source (λ = 0.15406 nm) in the 2θ-range of 10°-70° with a scan speed of 0.03 degree s −1 . The vibrational studies were recorded using Raman spectroscopy at a spectral resolution of ≈ 1 cm −1 (Jobin Yvon HR800UV, Longjumeau, France) using an excitation wavelength of 632.8 nm (He:Ne laser) and using Fourier transform infrared (FTIR) spectroscopy (VERTEX 80v, Bruker, Karlsruhe, Germany) in the spectral range 100-1000 cm −1 . The optical absorption studies were carried out using a UV−visible spectrophotometer (UV-VIS-NIR, Hitachi U3400, Tokyo, Japan). The chemical valence state of Mo and surface chemical composition were probed using X-ray photoelectron spectroscopy (XPS, K-Alpha-Thermo Scientific spectrometer, Dreieich, Germany) with a monochromatic Al Kα X-ray source (λ = 1486.68 eV). The surface morphological characteristics of the films were studied using scanning electron microscope (SEM, Carl Zeiss, EVO-MA15, Oberkochen, Germany). The surface crystallography was investigated using reflection high-energy electron diffraction (RHEED, EFZ4 device, Carl Zeiss, Germany). The surface topography was analyzed using atomic force microscopy (AFM, Park NX10, Park Sytems, Suwon, Korea). Thickness of the films was measured using an optical profilometer (model Profilm3D, Filmetrics, San Diego, CA, USA).
Investigation of the electrochemical properties of MoO3-δ thin film (1 cm 2 in area, 450 nm in thickness) was carried out using CR2477-type coin cells (Renata, Brive, France). Particular attention was paid to the coin-cell assembly by addition of an O-ring to avoid the presence of Li with Si and SiO2 materials (Figure 3b). Non-aqueous Li//MoO3-δ cells were assembled in argon-filled glovebox with Li foil as a combined counter and reference electrode and Whatman GF/D borosilicate glass fiber as separator. The aprotic electrolyte was 1 mol L −1 LiPF6 dissolved in ethylene carbonate and dimethyl carbonate (EC:DMC; 1:1 w/w). Data were collected using an electrochemical analyzer CHI 608C (CH Instruments Inc., USA) at a current density of 10-200 μA cm −2 . For a transfer of 6 moles of electrons per Mo, the current density of 1.17 A g −1 corresponds to 1C rate. Electrochemical impedance spectroscopy (EIS) data were collected in the range 0.1 Hz-1 MHz with a bias voltage of 5 mV.

Structural Properties
The structure of PLD MoO3-δ thin films was investigated using XRD, Raman and FTIR spectroscopy. All the films show uniform thickness, but they vary from dark blue to white in color as a function of the deposition parameters. Figure 4a,b shows the XRD diagrams of PLD MoO3-δ thin films prepared under different conditions of temperature and oxygen partial pressure. When deposited at Ts = 25 °C, at any PO2 the films are amorphous (a-MoO3) in nature. Their XRD spectrum displays a broad band centered at 2θ = 23° characteristic of an amorphous phase. For films deposited at Ts = 100 °C, the XRD spectra display the patterns of the mixed phase with both (0k0) and (0kl) orientations of the α-and β-MoO3, respectively, the β-MoO3 phase being characterized by the (011)

MoO x Thin Films Fabrication
MoO 3−δ thin films were deposited on the Au/Ti/SiO 2 /Si multilayer substrates described above. The target was composed of MoO 3 powders synthesized by solid-state reaction from molybdic acid (H 2 MoO 4 ) as precursor (Merck KGaA, Darmstadt, Germany, 99.99% grade). The product was ground for 3 h and pressed into a pellet at a pressure of 392 MPa. Subsequently, the pellet was heated at 10 • C min −1 and sintered at 650 • C for 3 h in air to yield a target with a density of about 91% of that of dense MoO 3 (4.2 g cm −3 ). MoO 3−δ thin films were deposited using PLD on substrates maintained in the range 25 ≤ T s ≤ 500 • C. The target was ablated using a KrF excimer laser λ = 248 nm (COMPex 201, Coherent, Göttingen, Germany) with laser fluence of 0.2-1.0 J cm −2 and repetition rate of 10 Hz. The target was rotated at 10 rotations min −1 to obtain homogenous thin films and circumvent complete depletion of the material at the same spot. The target-to-substrate distance was 4 cm, and the incident laser beam ablated the target surface with an angle of 45 • . The system chamber was evacuated to a base pressure of 0.4 mPa prior to film deposition. During reactive deposition, the chamber was filled with pure oxygen gas maintained at P O 2 in the range 0.5-13 Pa to obtain MoO 3−δ films with different deviations from stoichiometry. The thickness of the PLD-prepared MoO 3−δ films was~450 ± 5 nm.

Film Characterization
The phase and structure of as-prepared MoO 3−δ films was characterized using X-ray diffraction (XRD), Raman scattering (RS) and Fourier transform infrared (FTIR) spectroscopy. XRD patterns were investigated (X-ray diffractometer 3003 TT, Siefert, Leuven, Belgium) using a CuKα radiation source (λ = 0.15406 nm) in the 2θ-range of 10 • -70 • with a scan speed of 0.03 degree s −1 . The vibrational studies were recorded using Raman spectroscopy at a spectral resolution of ≈1 cm −1 (Jobin Yvon HR800UV, Longjumeau, France) using an excitation wavelength of 632.8 nm (He:Ne laser) and using Fourier transform infrared (FTIR) spectroscopy (VERTEX 80v, Bruker, Karlsruhe, Germany) in the spectral range 100-1000 cm −1 . The optical absorption studies were carried out using a UV−visible spectrophotometer (UV-VIS-NIR, Hitachi U3400, Tokyo, Japan). The chemical valence state of Mo and surface chemical composition were probed using X-ray photoelectron spectroscopy (XPS, K-Alpha-Thermo Scientific spectrometer, Dreieich, Germany) with a monochromatic Al K α X-ray source (λ = 1486.68 eV). The surface morphological characteristics of the films were studied using scanning electron microscope (SEM, Carl Zeiss, EVO-MA15, Oberkochen, Germany). The surface crystallography was investigated using reflection high-energy electron diffraction (RHEED, EFZ4 device, Carl Zeiss, Germany). The surface topography was analyzed using atomic force microscopy (AFM, Park NX10, Park Sytems, Suwon, Korea). Thickness of the films was measured using an optical profilometer (model Profilm3D, Filmetrics, San Diego, CA, USA).
Investigation of the electrochemical properties of MoO 3−δ thin film (5 cm 2 in area, 450 nm in thickness) was carried out using CR2477-type coin cells (Renata, Brive, France). Particular attention was paid to the coin-cell assembly by addition of an O-ring to avoid the presence of Li with Si and SiO 2 materials (Figure 3b). Non-aqueous Li//MoO 3−δ cells were assembled in argon-filled glovebox with Li foil as a combined counter and reference electrode and Whatman GF/D borosilicate glass fiber as separator. The aprotic electrolyte was 1 mol L −1 LiPF 6 dissolved in ethylene carbonate and dimethyl carbonate (EC:DMC; 1:1 w/w). Data were collected using an electrochemical analyzer CHI 608C (CH Instruments Inc., USA) at a current density of 10-200 µA cm −2 . For a transfer of 6 moles of electrons per Mo, the current density of 1.17 A g −1 corresponds to 1C rate. Electrochemical impedance spectroscopy (EIS) data were collected in the range 0.1 Hz-1 MHz with a bias voltage of 5 mV.

Structural Properties
The structure of PLD MoO 3−δ thin films was investigated using XRD, Raman and FTIR spectroscopy. All the films show uniform thickness, but they vary from dark blue to white in color as a function of the deposition parameters. Figure 4a,b shows the XRD diagrams of PLD MoO 3−δ thin films prepared under different conditions of temperature and oxygen partial pressure. When deposited at T s = 25 • C, at any P O 2 the films are amorphous (a-MoO 3 ) in nature. Their XRD spectrum displays a broad band centered at 2θ = 23 • characteristic of an amorphous phase. For films deposited at T s = 100 • C, the XRD spectra display the patterns of the mixed phase with both (0k0) and (0kl) orientations of the αand β-MoO 3 , respectively, the β-MoO 3 phase being characterized by the (011) and (022) reflections (JCPDS card No. . For films prepared in the range 150 < T s < 380 • C, the XRD reflections are those of the unique α-MoO 3−δ phase indexed to the (0k0) plane reflections. They give evidence of a highly preferred orientation, i.e., basal plane parallel to the substrate surface. These lines match well with the standard patterns of the orthorhombic α-MoO 3 phase (JCPDS card No. 05-0508). Moreover, in contrast to the crystalline MoO 3 , the broadened diffraction peak located at 2θ ≈ 25 • is related to the nano-size of the crystallites and the distorted layered structure that is in accordance with published results [59,60]. For the single α-MoO 3−δ phase, the ratio between intensities of the (020) and (060) lines increases with the increase of the substrate temperature. Their intensities are almost identical for T s = 350 • C. Similar trends were observed by Julien et al. at a lower temperature of 250 • C for flash-deposited films [31]. The XRD patterns of the MoO 3−δ films deposited at T s > 400 • C contain two distinct sets of reflections attributed to the coexistence of the α-MoO 3−δ and Mo 8 O 23 phases. This latter phase is well defined with reflections at 2θ ≈ 19.  . In summary, several remarks are worth making: (i) the deposition temperature and the partial oxygen pressure are the preponderant parameters tuning the film stoichiometry, while the laser fluence governs mainly the deposition rate, (ii) the best crystallinity of the single α-MoO 3−δ phase is obtained for films deposited at moderate temperature, (iii) with the increase of the adatom mobility, the thermodynamics favor the hierarchy α-MoO 3 , α-β-MoO 3−δ < α-MoO 3−δ < Mo 8 O 23 , and (iii) an increase in T s up to 400 • C favors the formation of films with high oxygen deficiency. The change in the oxygen deficiency has also been observed on films fabricated using different techniques. Guerfi et al. reported the growth by electrodeposition of MoO 2.8 (Mo 5 O 14 ) film at room temperature. This metastable tetragonal phase with a strong preferred orientation along the plane (10,10,0) transformed to orthorhombic MoO 3 after heat treatment at 260 • C [61]. This behavior has been investigated for MoO 3−δ films fabricated using reactive RF-magnetron sputtering [62], electron beam evaporation [37], plasma-assisted activated reactive evaporation [63] and CVD [64]. Note that the deviation from stoichiometry can be evidenced by studying the fundamental absorption edge, which can be evaluated in terms of direct inter-band transitions. The optical band-gap depends on oxygen vacancies as donor centers close to the valence band. Julien et al. showed that the optical band-gap of MoO 3−δ films deposited at 25 • C is 3.37 eV (large δ value) and decreases to 2.80 eV for the films prepared at 300 • C (small δ value) [31]. Sivakumar et al. [37] reported a band-gap of 2.8 eV for films deposited on glass substrate maintained at 25 • C (light gray in color), decreasing to 2.35 eV when deposited at 200 • C (deep blue color).
Electrochem 2020, 2, FOR PEER REVIEW 7 and (022) reflections (JCPDS card No. 47-1081). For films prepared in the range 150 < Ts < 380 °C, the XRD reflections are those of the unique α-MoO3-δ phase indexed to the (0k0) plane reflections. They give evidence of a highly preferred orientation, i.e., basal plane parallel to the substrate surface. These lines match well with the standard patterns of the orthorhombic α-MoO3 phase (JCPDS card No. 05-0508). Moreover, in contrast to the crystalline MoO3, the broadened diffraction peak located at 2θ ≈ 25° is related to the nano-size of the crystallites and the distorted layered structure that is in accordance with published results [59,60]. For the single α-MoO3-δ phase, the ratio between intensities of the (020) and (060)   . In summary, several remarks are worth making: (i) the deposition temperature and the partial oxygen pressure are the preponderant parameters tuning the film stoichiometry, while the laser fluence governs mainly the deposition rate, (ii) the best crystallinity of the single α-MoO3-δ phase is obtained for films deposited at moderate temperature, (iii) with the increase of the adatom mobility, the thermodynamics favor the hierarchy α-MoO3, α-β-MoO3-δ < α-MoO3-δ < Mo8O23, and (iii) an increase in Ts up to 400 °C favors the formation of films with high oxygen deficiency. The change in the oxygen deficiency has also been observed on films fabricated using different techniques. Guerfi et al. reported the growth by electrodeposition of MoO2.8 (Mo5O14) film at room temperature. This metastable tetragonal phase with a strong preferred orientation along the plane (10,10,0) transformed to orthorhombic MoO3 after heat treatment at 260 °C [61]. This behavior has been investigated for MoO3-δ films fabricated using reactive RFmagnetron sputtering [62], electron beam evaporation [37], plasma-assisted activated reactive evaporation [63] and CVD [64]. Note that the deviation from stoichiometry can be evidenced by studying the fundamental absorption edge, which can be evaluated in terms of direct inter-band transitions. The optical band-gap depends on oxygen vacancies as donor centers close to the valence band. Julien et al. showed that the optical band-gap of MoO3-δ films deposited at 25 °C is 3.37 eV (large δ value) and decreases to 2.80 eV for the films prepared at 300 °C (small δ value) [31]. Sivakumar et al. [37] reported a band-gap of 2.8 eV for films deposited on glass substrate maintained at 25 °C (light gray in color), decreasing to 2.35 eV when deposited at 200 °C (deep blue color).  Vibrational spectroscopies, Raman and FTIR, are powerful probes for the analysis of the short-range order in solids. They are used to confirm the presence of O-defects in MoO 3−δ films [65,66]. Figure 5a displays the Raman scattering spectra of PLD MoO 3−δ thin films deposited at temperature in the range 25-300 • C. Data were collected using the laser excitation wavelength of 632.3 nm. Analysis of spectral features is based on stretching and bending modes of octahedral MoO 6 entities building the MoO 3 framework. In Figure 5a, there are four intense modes in the range of 600-1100 cm −1 , which are the characteristic bands inherent to the fully oxygenated α-MoO 3 . They are located at 995, 900, 820 and 660 cm −1 corresponding to the asymmetrical (A g , ν as Mo=O) and symmetrical (A g , ν s Mo=O) stretching vibrations of the terminal Mo=O double bonds and to the asymmetrical (B 2g , B 3g , ν as O-Mo-O) stretching vibrations of O-Mo-O bonds, respectively [41,67]. The Raman active modes in the low-frequency region at 283 and 340 cm −1 are attributed to the Mo-O-Mo bending vibrations [68]. The deposition temperature-dependent vibrational properties and changes in the chemical bonds were successfully analyzed based on the concept of group frequency developed by Cotton and Wing [69]. The relative intensities of vibration modes and Raman shifts of MoO x compounds are significantly impacted by the size and shape of crystallites and by the crystal orientation as well [70]. Furthermore, even for a small stoichiometric deviation in MoO 3−δ (δ < 0.04) a frequency shift of the vibrational modes is observed in the Raman spectrum, indicating a weaker Mo-O bonding due to oxygen deficiency [71]. Vibrational spectroscopies, Raman and FTIR, are powerful probes for the analysis of the shortrange order in solids. They are used to confirm the presence of O-defects in MoO3-δ films [65,66]. Figure 5a displays the Raman scattering spectra of PLD MoO3-δ thin films deposited at temperature in the range 25-300 °C. Data were collected using the laser excitation wavelength of 632.3 nm. Analysis of spectral features is based on stretching and bending modes of octahedral MoO6 entities building the MoO3 framework. In Figure 5a, there are four intense modes in the range of 600-1100 cm −1 , which are the characteristic bands inherent to the fully oxygenated α-MoO3.  [68]. The deposition temperature-dependent vibrational properties and changes in the chemical bonds were successfully analyzed based on the concept of group frequency developed by Cotton and Wing [69]. The relative intensities of vibration modes and Raman shifts of MoOx compounds are significantly impacted by the size and shape of crystallites and by the crystal orientation as well [70]. Furthermore, even for a small stoichiometric deviation in MoO3-δ (δ < 0.04) a frequency shift of the vibrational modes is observed in the Raman spectrum, indicating a weaker Mo-O bonding due to oxygen deficiency [71]. The Raman spectra of MoO3-δ thin films show clearly the change in morphology and composition from amorphous a-MoO3 and α-MoO3-δ to single crystal α-MoO3. These structural modifications are visualized by a shift of the three high-frequency stretching modes at 995, 820, and 660 cm −1 for orthorhombic α-MoO3 to 976, 800 and 648 cm −1 for oxygendeficient MoO3-δ. The presence of Mo 5+ defects can be identified by the appearance of new Raman bands compared with the spectrum of stoichiometric MoO3. The Raman peaks located at 771 cm −1 confirm the presence of a localized Mo 5+ intermediate oxidation state produced by oxygen vacancies in α-MoO3-δ. The same set of peaks is observed in the Raman spectrum of Mn4O11 (MnO2. 75), in which the average oxidation state of Mo is +5.5. Chen et al. [72] investigated the presence of oxygen vacancies in MoO3 nanobelts (20 nm thick) using Raman spectroscopy, which was evidenced by the redshift of the main Raman peaks. Moreover, Lee et al. have assigned the peak at 400 cm −1 to vibrational bonds between Mo 5+ and oxygen [73]. The Raman spectrum of films deposited at 450 °C presents the superposition of the α-MoO3 and Mo8O23 vibrational modes. The peaks at 225, 255, 377, 675, 902 and 958 cm −1 are assigned to the m-Mo8O23 phase [62,74]. The FTIR absorbance spectra of PLD MoO 3−δ thin films grown at various temperatures (25 ≤ T s ≤ 400 • C) are presented in Figure 5b. In the high-wavenumber region (500-1000 cm −1 ), the broad absorption band is resolved into several sub-bands at 570, 610, 818, 871, 993 and 1008 cm −1 corresponding to the stretching modes involving (O-Mo 2 ), (O-Mo 3 ) and (Mo=O) bonds. The broader peak at~650 cm −1 is attributed to bridging Mo-O-Mo bonds of Mo 6+ [75,76]. The splitting of the Mo=O stretching at 993 and 1008 cm −1 reflects the good crystallinity of the films. It corresponds to the vibration of the zig-zag rows, which is a typical feature of the layered α-MoO 3 phase. The bending infrared modes are observed in the low-frequency region (200-500 cm −1 ). The presence of Mo 5+ defects can be identified by the appearance of new IR bands at 732 cm −1 compared with the spectrum of stoichiometric MoO 3 . This mode is assigned to the stretching mode of Mo 5+ . Dun et al. observed the mode characteristic of the bridging O-Mo 2 of Mo 5+ at 700 cm −1 [77]. Similarly to the Raman patterns, the FTIR spectra of sub-stoichiometric MoO 3−δ films display a frequency shift upon oxygen deficiency. This is consistent with the literature reports. Sun et al. reported the ν(O-Mo 2 ) band shifts towards lower wave numbers (868 → 721 → 683 cm −1 ) for sub-stoichiometric α-MoO 3−δ due to the reduction of the MoO 3 lattice through the formation of oxygen vacancies in the doubly coordinated oxygen. In contrast, the ν(O=Mo) band shifts towards higher wavenumbers (948 → 973 → 975 cm −1 ) [78].

Composition
The composition of as-deposited MoO 3−δ thin films was characterized by high-resolution XPS. Figure 6a-c shows the XPS spectra of the Mo 3d core level in MoO 3−δ thin films deposited on substrate maintained at 25, 300 and 450 • C under oxygen pressure of 6 Pa. The XPS spectra were analyzed by evaluating the peak area of elements using Gaussian profiles after removing the secondary electron background. All XPS spectra can be deconvoluted using two Mo 3d doublets (3d 5/2 and 3d 3/2 ). The first doublet centered at 236 and 232.8 eV is typical of the Mo 6+ state with a spin-orbit separation of~3.2 eV, whereas the second one located at 234.6 and 231.9 eV is due to Mo 5+ [79,80]. The film deposited at 25 • C is highly disordered and shows a large stoichiometric deviation (Figure 6a). When deposited at 300 • C, the XPS spectrum is dominated by the Mo 6+ doublet that corresponds to nearly MoO 3 films (Figure 6b), while the increase of the Mo 5+ doublet is observed for films prepared at T > 400 • C (Figure 6c). In the case of film deposited at 450 • C, the combination of XRD and XPS data provides the composition 0.85Mo 8 O 23 -0.15MoO 3 . From the variation of the O/Mo ratio ( Figure 6), we can distinguish three regions related with the morphology and composition of the films, i.e., in region (I) the MoO 3−δ films are highly disordered and off-stoichiometric, region (II) includes near-stoichiometric films, and region (III) encompasses oxygen deficient lattices. These results show that Mo cation exists in the highest oxidation state in MoO 3−δ thin films grown in the range 200 ≤ T s ≤ 400 • C, which suggests an efficient gas phase reaction in the plume of the laser ablation. When the temperature increases beyond 400 • C, the films contain Mo in lower oxidation state +6 < n < +5 due to dissociation. The re-evaporation from the film surface making the incorporation of oxygen into the lattice more difficult cannot be ruled out.
Electrochem 2020, 2, FOR PEER REVIEW 10 vacancies. However, it is unlikely that oxygen vacancies remain as point defects, as ordering of oxygen vacancies has been observed at very low concentrations [82].  Taking k as the ratio of the quantity of Mo 5+ /Mo 6+ calculated from the ratio of the integrated peak area, the stoichiometric deviation can be calculated by the relation [81]: Table 1 reports the composition of MoO 3−δ thin films as a function of the deposition temperature calculated using Equation (1) from data in Figure 7. It is well known that the α-MoO 3 phase retains the orthorhombic structure until the O/Mo ratio is reduced to about 2.89 with the formation of the shear phase like Mo 9 O 26 or Mo 8 O 23 (monoclinic, P2/c space group) for higher concentrations of oxygen vacancies. However, it is unlikely that oxygen vacancies remain as point defects, as ordering of oxygen vacancies has been observed at very low concentrations [82]. Electrochem 2020, 2, FOR PEER REVIEW 10 vacancies. However, it is unlikely that oxygen vacancies remain as point defects, as ordering of oxygen vacancies has been observed at very low concentrations [82].   This finding is consistent with previous investigations [45,83,84]. For all films, the minor 3d doublet for Mo 5+ in the XPS spectra indicates that the perfectly stoichiometric MoO3 was not reached. These results suggest that the growth of MoO3-δ thin films with tunable oxygen vacancies favors their electrical properties as anode materials for electrochemical micro-systems. The MoO3-δ films thermally deposited at 25 °C are sub-stoichiometric with an O/Mo atomic ratio of 2.73 increasing to 2.91 after annealing at 200 °C, while heat treatment at 400 °C reduces the oxygen vacancy defects [83]. Fernandes-Cauduro et al. investigated the effects of in situ annealing in MoO3-δ thin films prepared using reactive sputtering. The deviation from stoichiometry decreases from 0.125 to 0.11 for film treated at 200 °C but it then increases (δ = 0.20) for the film formed at 500 °C, which is explained by the appearance of the 6% Mo 4+ state at the surface [85]. A similar trend was observed by Han et al. [86] for MoO3-δ films prepared using thermal evaporation. It is suggested that the relative amount of Mo 5+ states exhibits a minimum for films annealed at 200 °C, i.e., composition MoO2.75. Note that, in most cases, the oxidation state of Mo 4+ was not detected except for films in which the MoO2 phase coexists with α-MoO3.
Recently, Novotny and Lamb [84] prepared MoOx films deposited on α-Al2O3 (0001) at 580 °C using a conventional molecular-beam epitaxy Knudsen cell. The films deposited in a vacuum and in This finding is consistent with previous investigations [45,83,84]. For all films, the minor 3d doublet for Mo 5+ in the XPS spectra indicates that the perfectly stoichiometric MoO 3 was not reached. These results suggest that the growth of MoO 3−δ thin films with tunable oxygen vacancies favors their electrical properties as anode materials for electrochemical micro-systems. The MoO 3−δ films thermally deposited at 25 • C are sub-stoichiometric with an O/Mo atomic ratio of 2.73 increasing to 2.91 after annealing at 200 • C, while heat treatment at 400 • C reduces the oxygen vacancy defects [83]. Fernandes-Cauduro et al. investigated the effects of in situ annealing in MoO 3−δ thin films prepared using reactive sputtering. The deviation from stoichiometry decreases from 0.125 to 0.11 for film treated at 200 • C but it then increases (δ = 0.20) for the film formed at 500 • C, which is explained by the appearance of the 6% Mo 4+ state at the surface [85]. A similar trend was observed by Han et al. [86] for MoO 3−δ films prepared using thermal evaporation. It is suggested that the relative amount of Mo 5+ states exhibits a minimum for films annealed at 200 • C, i.e., composition MoO 2.75 . Note that, in most cases, the oxidation state of Mo 4+ was not detected except for films in which the MoO 2 phase co-exists with α-MoO 3 .
Recently, Novotny and Lamb [84] prepared MoO x films deposited on α-Al 2 O 3 (0001) at 580 • C using a conventional molecular-beam epitaxy Knudsen cell. The films deposited in a vacuum and in O 2 atmosphere at 580 • C were oxygen deficient with an average formula MoO 2.67 . Their remarkable thermal stability is due to oxygen vacancies. For these films the XPS binding energies are recorded at 530.6 ± 0.1 eV for O 1s and 232.6, 231.0 and 229.2 ± 0.1 eV for Mo 3d 5/2 peaks of Mo 6+ , Mo 5+ and Mo 4+ oxidation states, respectively. Similar binding energies were obtained by Bhosle et al. [45] for MoO 2.75 films grown using the PLD method on sapphire substrates at 500-600 • C under O 2 pressure of 0.1 Pa. MoO x films prepared using hot wire the oxidation-sublimation deposition (HWOSD) technique at room temperature under an oxygen pressure of 0.2 Pa were amorphous. XPS analysis showed that the composition MoO 2.94 and the oxygen vacancies led to an electrical conductivity of 1.6 × 10 −6 S cm −1 [87]. XPS studies of the MoO x (x = 2.64-2.73) films fabricated using RF reactive magnetron sputtering with various oxygen flow rates (1.6-3.6 sccm) under pressure fixed at 0.4 Pa indicated the presence of Mo 5+ and Mo 6+ oxidation states and excluded the presence of Mo 4+ states [88]. These films are semiconductors (n-type) in nature.

Morphology
The crystallite size (or coherent length) was determined from the full-width at half-maximum β 0k0 of the (0k0) XRD reflections, using the Scherrer formula L c = Kλ/β 0k0 cosθ 0k0 . The average crystallite size varied between 23 and 69 nm in the T s range 200-450 • C. Since surface roughness is an important physical parameter that affects the electrochemical performance, the surface topography was investigated using atomic force microscopy (AFM). Figure 8a shows the SEM image of the substrate surface. The typical surface morphology of a PLD MoO 3−δ films grown under P O2 = 13 Pa was investigated using AFM imaging (Figure 8b-d). The AFM results showed that the root-mean-square (RMS) roughness varies from 1.2 to 4.8 nm for the film deposited in the T s range 100-400 • C. The growth temperature dependence of the grain size is plotted in Figure 9. It is observed that, in the range 100-500 • C, the grain size follows an exponential law with T s . The surface roughness of films formed at emphT s = 450 • C, i.e., 0.85Mo 8 O 23 -0.15MoO 3 , is about 7.2 nm, which is consistent with the XRD data. This is in contrast with the results reported by Hussain et al. [43]. They determined a surface roughness of 3.4 nm for crystalline MoO 3 films deposited using PLD at 200 • C under P O2 of 10 Pa and concluded that the surface roughness of the films decreased and the grain size increased with the increase in T s . The different behavior is attributed to high stress in the films generated by the growth of films on substrate with different textures (amorphous, crystallized, textured, metallic) [31]. Fernandes-Cauduro et al. [85] studied the effect of oxygen partial pressure and sputtering power on amorphous DC-sputtered MoO x films. It was found that the surface roughness (R rms ) was calculated from AFM relative height profiles for the as-deposited films on commercial ITO decreased from 2.2 to 1.0 nm, when the P O2 increased from 0.1 to~0.3 Pa.  [88]. These films are semiconductors (n-type) in nature.

Morphology
The crystallite size (or coherent length) was determined from the full-width at half-maximum β0k0 of the (0k0) XRD reflections, using the Scherrer formula Lc = Kλ/β0k0 cosθ0k0. The average crystallite size varied between 23 and 69 nm in the Ts range 200-450 °C. Since surface roughness is an important physical parameter that affects the electrochemical performance, the surface topography was investigated using atomic force microscopy (AFM). Figure 8a shows the SEM image of the substrate surface. The typical surface morphology of a PLD MoO3-δ films grown under PO2 = 13 Pa was investigated using AFM imaging (Figure 8b-d). The AFM results showed that the root-mean-square (RMS) roughness varies from 1.2 to 4.8 nm for the film deposited in the Ts range 100-400 °C. The growth temperature dependence of the grain size is plotted in Figure 9. It is observed that, in the range 100-500 °C, the grain size follows an exponential law with Ts. The surface roughness of films formed at Ts = 450 °C, i.e., 0.85Mo8O23-0.15 MoO3, is about 7.2 nm, which is consistent with the XRD data. This is in contrast with the results reported by Hussain et al. [43]. They determined a surface roughness of 3.4 nm for crystalline MoO3 films deposited using PLD at 200 °C under PO2 of 10 Pa and concluded that the surface roughness of the films decreased and the grain size increased with the increase in Ts. The different behavior is attributed to high stress in the films generated by the growth of films on substrate with different textures (amorphous, crystallized, textured, metallic) [31]. Fernandes-Cauduro et al. [85] studied the effect of oxygen partial pressure and sputtering power on amorphous DC-sputtered MoOx films. It was found that the surface roughness (Rrms) was calculated from AFM relative height profiles for the as-deposited films on commercial ITO decreased from 2.2 to 1.0 nm, when the PO2 increased from 0.1 to ~0.3 Pa.   Figure 10 shows the galvanostatic discharge-charge profiles of the MoO3-δ thin-film electrodes vs. Li + /Li at a current rate of 1 A g −1 (~1C) in the potential window 3.0-0.05 V. For all electrodes, the voltage dropped from the open-circuit voltage (OCV) of ~3.0 V during lithiation to reach firstly a small pseudo-plateau at ca. 2.4 V and secondly a large plateau at ca. 0.4-0.5 V. During the charge process, the voltage curve displays a smooth slope region with a loss of capacity. The chargedischarge curves of MoO3-δ thin films display a similar S-shape. For all samples, we observe that the first charge capacities are lower than the first discharge capacities. After the first cycle of cell formation, the charge capacities stabilized. A large irreversible capacity decay of the first cycle is commonly observed for the transition-metal oxides because of the formation of the solid electrolyte interphase (SE)I layer and the irreversible phase transformation [89][90][91][92].

Electrochemical Properties
The electrochemical activity for Li + storage of the oxygen-deficient oxides occurs with various   Figure 10 shows the galvanostatic discharge-charge profiles of the MoO 3−δ thin-film electrodes vs. Li + /Li at a current rate of 1 A g −1 (~1C) in the potential window 3.0-0.05 V. For all electrodes, the voltage dropped from the open-circuit voltage (OCV) of~3.0 V during lithiation to reach firstly a small pseudo-plateau at ca. 2.4 V and secondly a large plateau at ca. 0.4-0.5 V. During the charge process, the voltage curve displays a smooth slope region with a loss of capacity. The charge-discharge curves of MoO 3−δ thin films display a similar S-shape. For all samples, we observe that the first charge capacities are lower than the first discharge capacities. After the first cycle of cell formation, the charge capacities stabilized. A large irreversible capacity decay of the first cycle is commonly observed for the transition-metal oxides because of the formation of the solid electrolyte interphase (SE)I layer and the irreversible phase transformation [89][90][91][92].

Electrochemical Properties
The electrochemical activity for Li + storage of the oxygen-deficient oxides occurs with various degrees of Li + -ion uptake; the initial discharge specific capacity varied between 390 and 484 µAh cm −2 µm −1 . For thin-film electrodes, the specific capacity is generally expressed as volumetric capacity (µAh cm −2 µm −1 ) due to the inaccuracy of the weight and density. Table 2 summarizes the electrochemical characteristics of the different MoO 3−δ thin film electrodes, i.e., specific discharge capacity of the 1st and 2nd discharge, Coulombic efficiency (CE) and specific capacity after 100 cycles. The initial CE of MoO 3−δ thin film anodes remains almost at~75% irrespective of the deposition temperature. The high reversible capacity of the MoO 2.894 electrode (~484 µAh cm −2 µm −1 ), which corresponds to the insertion of~5. 78 Li + /Mo, is close to the theoretical maximum (5. 79 Li + /Mo). After 100 cycles performed at a 1 A g −1 rate, this film delivers a specific capacity of 300 µAh cm −2 µm −1 (~670 mAh g −1 ). The irreversible insertion is largest for the electrode MoO 2.982 , which can be attributed to its low conductivity. Note that the theoretical specific capacity of 1117 mAh g −1 (6 Li per Mo) corresponds to a volumetric capacity of 5026 mAh cm −3 (taking a density of 4.5 g cm −3 ), which is equivalent to 502.6 µAh cm −2 µm −1 for MoO 3 thin film.  Figure 10 shows the galvanostatic discharge-charge profiles of the MoO3-δ thin-film electrodes vs. Li + /Li at a current rate of 1 A g −1 (~1C) in the potential window 3.0-0.05 V. For all electrodes, the voltage dropped from the open-circuit voltage (OCV) of ~3.0 V during lithiation to reach firstly a small pseudo-plateau at ca. 2.4 V and secondly a large plateau at ca. 0.4-0.5 V. During the charge process, the voltage curve displays a smooth slope region with a loss of capacity. The chargedischarge curves of MoO3-δ thin films display a similar S-shape. For all samples, we observe that the first charge capacities are lower than the first discharge capacities. After the first cycle of cell formation, the charge capacities stabilized. A large irreversible capacity decay of the first cycle is commonly observed for the transition-metal oxides because of the formation of the solid electrolyte interphase (SE)I layer and the irreversible phase transformation [89][90][91][92].

Electrochemical Properties
The electrochemical activity for Li + storage of the oxygen-deficient oxides occurs with various degrees of Li + -ion uptake; the initial discharge specific capacity varied between 390 and 484 μAh cm −2 μm −1 . For thin-film electrodes, the specific capacity is generally expressed as volumetric capacity (μAh cm −2 μm −1 ) due to the inaccuracy of the weight and density. Table 2 summarizes the electrochemical characteristics of the different MoO3-δ thin film electrodes, i.e., specific discharge capacity of the 1st and 2nd discharge, Coulombic efficiency (CE) and specific capacity after 100 cycles. The initial CE of MoO3-δ thin film anodes remains almost at ~75% irrespective of the deposition temperature. The high reversible capacity of the MoO2.894 electrode (~484 μAh cm −2 μm −1 ), which corresponds to the insertion of ~5. 78 Li + /Mo, is close to the theoretical maximum (5. 79 Li + /Mo). After 100 cycles performed at a 1 A g −1 rate, this film delivers a specific capacity of 300 μAh cm −2 μm −1 (~670 mAh g −1 ). The irreversible insertion is largest for the electrode MoO2.982, which can be attributed to its low conductivity. Note that the theoretical specific capacity of 1117 mAh g −1 (6 Li per Mo) corresponds to a volumetric capacity of 5026 mAh cm −3 (taking a density of 4.5 g cm −3 ), which is equivalent to 502.6 μAh cm −2 μm −1 for MoO3 thin film.   In the potential window 3-0.05 V, the lithiation in MoO 3 is believed to take place in two stages. Stage I occurs up to a potential of 1.5 V, in which Li + ions intercalate with the orthorhombic framework to form a solid solution according the relation: In Li x MoO 3 (Li bronze phase), the inserted lithium content ranges between 1.0 and 1.5, up to a potential of 1.5 V [93]. The Li + ions are accommodated in the interlayer spacing between octahedral basal Mo−O layers and can be reversibly extracted. The oxidation state of Mo in Li x MoO 3 decreases with x from +6 to about +4.5 leading to a corresponding fall in enthalpy of insertion. The stage II corresponds to potentials below 0.7 V, for which the further lithiation of Li x MoO 3 occurs by a mechanism of conversion. The full reduction is expressed by the relation [94]: During charge following the full conversion reaction, the extracted Mo metal is re-oxidized by Li 2 O. In the potential range below 0.7 V vs. Li + /Li, the discharge process is a conversion reaction occurring at ca. 0.2 V (Equation (3)); the reduction action of Mo 4+ to metallic Mo 0 occurs near 0.2 V. Li reacts with the Li x MoO 3−δ solid solution to form nano-scaled metallic particles nano-dispersed in a Li 2 O matrix. The reversibility of Li 2 O is due to the presence of nano-scaled Mo 0 particles [95]. Note that MoO 3 nanoparticles and films present similarly continuous and smooth discharge curves, unlike the bulk material which does not show significant capacity in the low-potential range. The primary reason for Coulombic efficiencies of~77% is the reversibility of Li 2 O in lithiated MoO 3−δ thin films studied here. Figure 11a,b shows the incremental capacity (IC) profiles, i.e., differential capacity (dQ/dV) vs. cell voltage, of the first lithiation-delithiation process in α-MoO 3−δ and 0.85Mo 8 O 23 -0.15MoO 3 thin films deposited at 450 • C. This analysis can be considered as an efficient tool to determine the electrochemical spectroscopy of electrodes [96]. For instance, IC has been successfully applied to analyze the behavior of doped or blended cathodes [97]. The IC curves were extracted from the galvanostatic charge-discharge (GCD) profiles (Figure 8c) [42,98,99]. Iriyama et al. [42] reported a TEM analysis of the two-phase reaction in electrochemical lithium insertion within α-Li x MoO 3 (0 ≤ x ≤ 0.25). The cyclic voltammogram exhibited a sharp peak at 2.77 V in the cathodic potential scan, which has no corresponding redox peak in the anodic scan, indicating the irreversibility of this reaction. As shown in Figure 9, the cathodic peak of the two-phase reaction occurs at lower potential (~2.30 V) due to the sub-stoichiometric composition of our films. The potential peaks in the first anodic scan (delithiation process) occurs at 1.20 and 1.72 V for α-MoO 3−δ thin films. The conversion peak shifts slightly toward a lower voltage for the mixed 0.85Mo 8 O 23 -0.15MoO 3 thin film, which results in more reduced MoO 3−δ . This oxygen-deficiency effect was also reported by Jung et al. [94]. Note that the anodic peaks are broad, which is indicative of the nano-size effect of pulverized Mo metallic particles after the full conversion process.
Electrochem 2020, 2, FOR PEER REVIEW 14 δ. This oxygen-deficiency effect was also reported by Jung et al. [94]. Note that the anodic peaks are broad, which is indicative of the nano-size effect of pulverized Mo metallic particles after the full conversion process.  Figure 12 compares the cycling performance of oxygen deficient thin films. The rate capability (Figure 10a) was investigated at different current densities in the range 10-1000 mA g −1 . The 0.85Mo8O23-0.15MoO3 thin film exhibits significantly enhanced capacity retention. Figure 10b Figure 12 compares the cycling performance of oxygen deficient thin films. The rate capability (Figure 12a) was investigated at different current densities in the range 10-1000 mA g −1 . The 0.85Mo 8 O 23 -0.15MoO 3 thin film exhibits significantly enhanced capacity retention. Figure 12b presents the discharge capacity with cycling at a current rate of 1 A g −1 over 100 cycles in the potential window 0.05-3.0 V for the three PLD thin films. For the amorphous a-MoO 3−δ film, the discharge capacity drops rapidly at the rate of 0.43% per cycle, while the discharge capacities of α-MoO 3−δ and 85Mo 8 O 23 -0.15MoO 3 films decrease slowly at the rate of 0.19% per cycle. The remarkable cycling performance of the mixed 85Mo 8 O 23 -0.15MoO 3 film is not only due to the higher electronic conductivity induced by oxygen vacancies but also benefits from the blend material composed of the layered MoO 3 and the quasi-1D Mo 8 O 23 suboxide. In comparison with the ratio of the discharge capacity at 0.01C and 1C rate, Q 1C /Q 0.01C (Figure 12a) is found to be 67%, 78% and 76% for a-MoO 3 , α-MoO 3 and 0.85Mo 8 O 23 -0.15MoO 3 thin films, respectively, which suggests a better Li + ion kinetics in the layered α-MoO 3 material.
The evolution of the surface morphology of α-MoO 3−δ thin films after electrochemical cycling has been investigated using scanning electron microscopy. Figure 13 represents the SEM images of the structural properties of the film at the 5th and 50th cycle. It is remarkable that the film surface is maintained after 50 charge-discharge cycles. This picture corroborates the good cyclability of the α-MoO 3−δ thin films deposited at 300 • C, showing a small capacity decay upon cycling.  Figure 12 compares the cycling performance of oxygen deficient thin films. The rate capability (Figure 10a) was investigated at different current densities in the range 10-1000 mA g −1 . The 0.85Mo8O23-0.15MoO3 thin film exhibits significantly enhanced capacity retention. Figure 10b presents the discharge capacity with cycling at a current rate of 1 A g −1 over 100 cycles in the potential window 0.05-3.0 V for the three PLD thin films. For the amorphous a-MoO3-δ film, the discharge capacity drops rapidly at the rate of 0.43% per cycle, while the discharge capacities of α-MoO3-δ and 85Mo8O23-0.15MoO3 films decrease slowly at the rate of 0.19% per cycle. The remarkable cycling performance of the mixed 85Mo8O23-0.15MoO3 film is not only due to the higher electronic conductivity induced by oxygen vacancies but also benefits from the blend material composed of the layered MoO3 and the quasi-1D Mo8O23 suboxide. In comparison with the ratio of the discharge capacity at 0.01C and 1C rate, Q1C/Q0.01C (Figure 10a) is found to be 67%, 78% and 76% for a-MoO3, α-MoO3 and 0.85Mo8O23-0.15MoO3 thin films, respectively, which suggests a better Li + ion kinetics in the layered α-MoO3 material. The evolution of the surface morphology of α-MoO3-δ thin films after electrochemical cycling has been investigated using scanning electron microscopy. Figure 13 represents the SEM images of the structural properties of the film at the 5th and 50th cycle. It is remarkable that the film surface is maintained after 50 charge-discharge cycles. This picture corroborates the good cyclability of the α-MoO3-δ thin films deposited at 300 °C, showing a small capacity decay upon cycling. The chemical diffusion coefficients of Li + ions in the thin film frameworks were evaluated using the potentiostatic intermittent titration technique (PITT). The theoretical and experimental aspects of the kinetic behavior of metal-like anode materials, i.e., Li-alloys LixAl or Li1+ySb, have been extensively studied by Weppner and Huggins [100,101] and further applied to guest ions in insertion electrodes [102,103]. In mixed-conducting electrodes, the apparent chemical diffusion for Li + ions, DLi, is the product (DLi = D0 W) of the component diffusion D0 by the enhancement factor (or thermodynamic factor) W = (∂ ln aLi/∂ ln cLi). Note that D0 is an intrinsic property of the material and WTF of thin film differs from WB of the bulk due to the difference in the activity of intercalant species [104]. From the point of view of fundamental studies, the use of thin film electrodes has the advantage of evaluating accurately intrinsic ion transport with known electrode surfaces by eliminating organic binders and conductive additives. The thermodynamic factor can be estimated from the gradient of the opencircuit voltage vs. composition ∂E/∂x as The chemical diffusion coefficients of Li + ions in the thin film frameworks were evaluated using the potentiostatic intermittent titration technique (PITT). The theoretical and experimental aspects of the kinetic behavior of metal-like anode materials, i.e., Li-alloys Li x Al or Li 1+y Sb, have been extensively studied by Weppner and Huggins [100,101] and further applied to guest ions in insertion electrodes [102,103]. In mixed-conducting electrodes, the apparent chemical diffusion for Li + ions, D Li , is the product (D Li = D 0 W) of the component diffusion D 0 by the enhancement factor (or thermodynamic factor) W = (∂ ln a Li /∂ ln c Li ). Note that D 0 is an intrinsic property of the material and W TF of thin film differs from W B of the bulk due to the difference in the activity of intercalant species [104]. From the point of view of fundamental studies, the use of thin film electrodes has the advantage of evaluating accurately intrinsic ion transport with known electrode surfaces by eliminating organic binders and conductive additives. The thermodynamic factor can be estimated from the gradient of the open-circuit voltage vs. composition (∂E/∂x) as: In PITT, a transient current (I t ) is recorded within a very small potential step. The local current depends on the concentration gradient at the particle surface using Fick's first law. The time-dependent electric transient current I t at each potential step obeys the following relation [b]: where F is the Faraday constant, A the surface area of the electrode, L the film thickness and ∆c Li = c s −c 0 the difference of Li + concentration at the surface at time t and at time t = 0 during each potential step. The primary diffusion parameter is the characteristic diffusion time τ, which is defined for the one-dimensional case as τ = L 2 /D Li . At short time (τ << L 2 /D Li ), the exponential term reduces to unity and the I t obeys the Cottrell's equation: In Figure 14a, the curve I t vs. t −1/2 shows a linear part with a slope proportional to the square root of the diffusion coefficient. This part shows primarily a finite-length diffusion behavior due to the interfacial charge-transfer and ohmic potential drop [105,106]. In the long-term domain, Equation (5) can be linearized by taking its logarithmic form: Electrochem 2020, 2, FOR PEER REVIEW 16 Which allows the determination of the apparent diffusion coefficient from the slope of the linear region of the ln(It) vs. t plot as shown in Figure 13b [96]. Figure 14a presents the typical variation of the thermodynamic factor against the cell voltage (W vs. E) calculated from the 2nd discharge response of a Li//α-MoO3-δ cell using Equation (4). The average value of W at E = 0.5 V is listed in Table 3. As expected for a metal oxide, W is in the range 20-30 [104]. Figure 14b compares the DLi values for PLD MoO3-δ films in direct contact with the aprotic electrolyte. Note that the curves DLi vs. cell voltage have the same features as two minimum values at 0.5 and 1.9 V. The results summarized in Table 3 show that the α-MoO3-δ thin film deposited at Ts = 300 °C exhibits the higher diffusion coefficient in the range from 4 × 10 −14 to 6 × 10 −12 cm 2 s −1 . This high value is due to the layered-like structure of α-MoO3-δ, inducing larger open Li sites and favorable transport paths during the discharge-charge processes.
which allows the determination of the apparent diffusion coefficient from the slope of the linear region of the ln(I t ) vs. t plot as shown in Figure 14b [96]. Figure 15a presents the typical variation of the thermodynamic factor against the cell voltage (W vs. E) calculated from the 2nd discharge response of a Li//α-MoO 3−δ cell using Equation (4). The average value of W at E = 0.5 V is listed in Table 3. As expected for a metal oxide, W is in the range 20-30 [104]. Figure 15b compares the D Li values for PLD MoO 3−δ films in direct contact with the aprotic electrolyte. Note that the curves D Li vs. cell voltage have the same features as two minimum values at 0.5 and 1.9 V. The results summarized in Table 3 show that the α-MoO 3−δ thin film deposited at T s = 300 • C exhibits the higher diffusion coefficient in the range from 4 × 10 −14 to 6 × 10 −12 cm 2 s −1 . This high value is due to the layered-like structure of α-MoO 3−δ , inducing larger open Li sites and favorable transport paths during the discharge-charge processes. region of the ln(It) vs. t plot as shown in Figure 13b [96]. Figure 14a presents the typical variation of the thermodynamic factor against the cell voltage (W vs. E) calculated from the 2nd discharge response of a Li//α-MoO3-δ cell using Equation (4). The average value of W at E = 0.5 V is listed in Table 3. As expected for a metal oxide, W is in the range 20-30 [104]. Figure 14b compares the DLi values for PLD MoO3-δ films in direct contact with the aprotic electrolyte. Note that the curves DLi vs. cell voltage have the same features as two minimum values at 0.5 and 1.9 V. The results summarized in Table 3 show that the α-MoO3-δ thin film deposited at Ts = 300 °C exhibits the higher diffusion coefficient in the range from 4 × 10 −14 to 6 × 10 −12 cm 2 s −1 . This high value is due to the layered-like structure of α-MoO3-δ, inducing larger open Li sites and favorable transport paths during the discharge-charge processes.

Growth Conditions
These results shed light on the growth process of the oxygen-deficient MoO 3−δ films prepared via different PLD conditions. The ability to generate these films with various compositions and morphologies is the most significant result of this work. The successful manufacture of sub-stoichiometric films has been optimized by the adjustment of the PLD growth parameters in order to obtain the best electronic properties and the best electrochemical properties for their potential application as electrodes for lithium microbatteries; the growth process of MoO 3−δ films has been discussed several times [4,[43][44][45][46]. Ramana and Julien [4] suggested that creation of oxygen vacancies in deposited MoO 3−δ film is due to the re-evaporation from the surface of the substrate. In this process, energy levels in the energy gap are generated close to the valance band (E v ) considered as donor centers. It is widely recognized that α-MoO 3−δ films prepared at low temperature (T s < 100 • C) are ill-textured and highly disordered with an amorphous contribution. Zhang et al. [50] reported that ITO/spin-coated MoO 2.98 films are amorphous and exhibit a deep-blue color. The influence of the partial O 2 pressure was discussed by several workers. Carcia and McCarron showed the influence of the partial O 2 pressure on the structure of magnetron-sputtered MoO 3 films. Using a total argon plus oxygen pressure of 1.3 Pa, α-MoO 3 grew under P O 2 = 50%, whereas β-MoO 3 (monoclinic structure) was formed under P O 2 ≈ 10% [107]. Similarly, Altman et al. [108] fabricated MoO 3 thin films using a molecular beam epitaxy technique assisted by oxygen plasma under a low O 2 pressure of 4 mPa. By manipulating the deposition conditions, i.e., temperature, deposit cycle and deposition rate, epitaxial β-MoO 3 (tetragonal phase) thin films were grown on SrLaAlO 4 substrates at 400 • C, while the decrease of the deposition temperature to 260 • C resulted in the formation of polycrystalline α-MoO 3 (orthorhombic phase) films. To obtain suboxide MoO 2.86 thin films using reactive magnetron sputtering at room temperature, the O 2 ratio of 15% was fixed in the Ar + O 2 gas mixture at a total pressure of~1 Pa [34]. The band-gap E g of 2. 82 [112]. The composition 0.65MoO 3 -0.35Mo 4 O 11 was obtained by calcination of the precursor (ammonium heptamolybdate tetrahydrate as a source of Mo) heat treated with a small fraction of zirconia under reduced atmosphere at 500 • C for 5 h in a 5% H 2 /Ar flow. Defective MoO x films grown on glass substrate at 300 • C using spray pyrolysis have been identified as Mo 9 O 26 phase. After annealing at 500 • C for 20 h in a controlled O 2 atmosphere, the XRD patterns of the heat-treated films showed the presence of two crystallographic phases: monoclinic Mo 9 O 26 and orthorhombic Mo 17 O 47 . The Mo 9 O 26 phase grew preferentially along the (712) plane [113].

Electronic Properties
MoO 3 is known to be an unintentional n-type semiconductor at ambient conditions, due to intrinsic point defects related to oxygen vacancies (V O ) and molybdenum interstitials (Mo i ) [13]. It has been widely recognized that the introduction of oxygen vacancies (V O ) is an efficient strategy to promote the specific discharge capacity and rate capability and to boost the Li + -ion diffusion of electrode materials [15][16][17][18][19][20]. Enhancement of the attainable capacity is attributed to extra valence electrons into the delocalized electron cloud. Similarly, in TiO 2 , due to oxygen vacancies, a narrowed bandgap and change of the Fermi energy level indicates an increase of electron concentration [114]. Recently, it has been demonstrated that anion vacancies in insertion electrode materials can improve the energy storage capacity and reduce the insertion energy and ion diffusion barrier in the host lattice [27,[115][116][117][118][119][120]. Figure 16 presents a diagram of the band structure close to the surface of MoO 3−δ film. Because of the oxygen vacancies, the Fermi level is enhanced and the concentration of electrons in the conduction band is increased [86]. The increase of Mo 5+ (reduction of the O/Mo ratio) induces a decrease of the surface dipole lowering the work function ϕ (ϕ ≈ 5.95 eV for T s = 450 • C). The large work function of MoO 3−δ films is attributed to its closed-shell character and the lowered electrostatic potential of the inner Mo-O units due to the dipole layer created by planes of terminal oxygen (O 1 ) sites. Therefore, the oxygen vacancy is a shallow donor, the reason why MoO 3−δ easily becomes a n-type degenerate semiconductor [121,122]. Julien et al. also suggested a decrease in optical band gap and attributed the decrease to the formation of oxygen-ion vacancies. The energy gap of MoO 3−δ films is located between 2.8 and 3.2 eV depending on the substrate and annealing temperature [31].
decrease of the surface dipole lowering the work function φ (φ ≈ 5.95 eV for Ts= 450 °C). The large work function of MoO3-δ films is attributed to its closed-shell character and the lowered electrostatic potential of the inner Mo-O units due to the dipole layer created by planes of terminal oxygen (O1) sites. Therefore, the oxygen vacancy is a shallow donor, the reason why MoO3-δ easily becomes a ntype degenerate semiconductor [121,122]. Julien et al. also suggested a decrease in optical band gap and attributed the decrease to the formation of oxygen-ion vacancies. The energy gap of MoO3-δ films is located between 2.8 and 3.2 eV depending on the substrate and annealing temperature [31].

Electrochemistry
Orthorhombic α-MoO3 is an attractive transition-metal oxide as an anode material for Li-ion batteries, owing to its unique layered structure and its high theoretical Li uptake consequently to the high oxidation state (+6) of Mo. It has a theoretical specific capacity of 1117 mAh g −1 based on the full conversion reaction in deep discharge voltages, which is higher than MoO2 (838 mAh g −1 ). The empty interlayer channels formed by its double-layered structure built from MoO6 octahedra provide rapid accommodation of Li ions. However, the insulating nature of α-MoO3 (band gap of 3.1 eV) and the structural degradation associated to the conversion reaction prevent its utilization as an anode [123]. To overcome this drawback, the design of a material with an oxygen-deficient lattice is the key issue to increase the electronic conductivity by several orders of magnitudes. Moreover, the use of hierarchical nano-structures also helps. Indeed, the implementation of a carbon-free MoO3-δ electrode is possible due to the enhanced electronic transport in oxygen-deficient oxides and the reduction of the Li + -ion pathway in nano-structured materials.
Few studies report the electrochemical features of Mo-oxide thin film anodes, which in contrast with traditional electrodes did not use conductive additive and binder [124,125]. Amorphous films (a-MoOx, 2 < x <3) were prepared using a reactive magnetron sputtering carried out at power of 60 W under a working pressure of 0.8 Pa and total gas flow rate at 40 sccm (Ar + O2 of 80:20) [125]. a-MoOx films are a mixture of MoO2 and MoO3 phase (O/Mo ≈ 2.75) and have the morphology of cauliflowerlike protuberances, which consist of ultra-fine particles. After 100 discharge-charge cycles at a current density of 90 μA cm −2 (~225 mA g −1 ) in the potential window 0.01-3.0 V, the specific capacity was 315 μAh cm −2 μm −1 (mass loading of 0.45 mg). Porous α-MoO3 films were fabricated using electrodeposition on Ni foam substrates. These films are composed of grains of several tens of nm in size with nanoholes of the same size [123]. After 50 cycles, the Li + insertion/extraction capacity and Figure 16. Band structure diagram near the MoO 3−δ film surface. E F , E g0 , E CBM , E VBM , ϕ, µ e , E V,local and E V,∞ represent the Fermi level, intrinsic band gap, conduction band minimum, valence band maximum, work function, electron chemical potential, local vacuum level and absolute vacuum level, respectively. Reproduced with permission from [86]. Copyright 2018 Elsevier.

Electrochemistry
Orthorhombic α-MoO 3 is an attractive transition-metal oxide as an anode material for Li-ion batteries, owing to its unique layered structure and its high theoretical Li uptake consequently to the high oxidation state (+6) of Mo. It has a theoretical specific capacity of 1117 mAh g −1 based on the full conversion reaction in deep discharge voltages, which is higher than MoO 2 (838 mAh g −1 ). The empty interlayer channels formed by its double-layered structure built from MoO 6 octahedra provide rapid accommodation of Li ions. However, the insulating nature of α-MoO 3 (band gap of 3.1 eV) and the structural degradation associated to the conversion reaction prevent its utilization as an anode [123]. To overcome this drawback, the design of a material with an oxygen-deficient lattice is the key issue to increase the electronic conductivity by several orders of magnitudes. Moreover, the use of hierarchical nano-structures also helps. Indeed, the implementation of a carbon-free MoO 3−δ electrode is possible due to the enhanced electronic transport in oxygen-deficient oxides and the reduction of the Li + -ion pathway in nano-structured materials.
Few studies report the electrochemical features of Mo-oxide thin film anodes, which in contrast with traditional electrodes did not use conductive additive and binder [124,125]. Amorphous films (a-MoO x , 2 < x < 3) were prepared using a reactive magnetron sputtering carried out at power of 60 W under a working pressure of 0.8 Pa and total gas flow rate at 40 sccm (Ar + O 2 of 80:20) [125]. a-MoO x films are a mixture of MoO 2 and MoO 3 phase (O/Mo ≈ 2.75) and have the morphology of cauliflower-like protuberances, which consist of ultra-fine particles. After 100 discharge-charge cycles at a current density of 90 µA cm −2 (~225 mA g −1 ) in the potential window 0.01-3.0 V, the specific capacity was 315 µAh cm −2 µm −1 (mass loading of 0.45 mg). Porous α-MoO 3 films were fabricated using electrodeposition on Ni foam substrates. These films are composed of grains of several tens of nm in size with nanoholes of the same size [123]. After 50 cycles, the Li + insertion/extraction capacity and the Coulombic efficiency of the thin-film electrode was tested at a current density of 3 A g −1 are 650 mAh g −1 and 97%, respectively. The smaller grain size and widespread nanoholes lead to high Li + insertion kinetics with a diffusion coefficient of 7.1 × 10 −11 cm 2 s −1 , obtained from EIS measurements.
The effect of annealing temperature on electrochemical performance of MoO x nanospheres was studied using binder-rich electrodes. The whitish-blue (indicative for x close to 3) particles deliver a capacity of~800 mAh g −1 after 40 cycles [131]. MoO 3 nanobelts synthesized using a simple hydrothermal route and heat-treated at 265 • C for 3 h in air exhibit an initial discharge capacity of 1250 mAh g −1 being cycled at a current rate of 0.1 C and an initial Coulombic efficiency of 67% [137]. Jung et al. prepared partially reduced Mo oxides, namely MoO 2.929 , MoO 2.903 and MoO 2.895 , by ball-milling as anode materials with 10% acetylene black. It is shown that the specific discharge capacity slightly increases with the increase of oxygen vacancy in the MoO 3 lattice. Remarkably, the MoO 2.895 anode can uptake 8Li per Mo in the potential range 0-3 V vs. Li + /Li [94]. MoO 3 nanobelts synthesized using a simple hydrothermal route and heat-treated at 265 • C for 3 h in air exhibit an initial discharge capacity of 1250 mAh g −1 at a rate of 0.1C and an initial Coulombic efficiency of 67% [137]. Liu et al. reported an initial discharge capacity of 1200 mAh g −1 at 50 mA g −1 for a crystalline MoO 3 anode containing 15% acetylene black, which fell at~320 mAh g −1 at the 10th cycle [138]. Ma et al. [139] investigated the electrochemical performance of suboxide-MoO 3 synthesized via a surfactant-assisted solvothermal route exhibiting an average valence state of 5.58 for Mo cations. The MoO 2.79 anode containing 15% acetylene black delivered a first discharge capacity of 760 mAh g −1 at 0.2 A g −1 rate. Wu et al. [128] reported a specific capacity up to 930 mAh·g −1 over 200 cycles for Mo-O anodes fabricated from slurry spread onto copper foils containing 20 wt.% acetylene black (conducting additive). According to the XRD pattern, the active material was a mixture of three phases, MoO 3 + Mo 4 O 11 + MoO 2 , and the XPS analysis showed the valence state of Mo on the surface as 5.6. In summary, our pulsed laser deposited MoO 3−δ suboxide films exhibit high specific capacity and good stability due to the improved Li + diffusion kinetics [140,141] as compared with films fabricated using electrodeposition [124] and electron-beam evaporation [125].

Li + -ion Kinetics
The potentiostatic intermittent titration technique has been used successfully to study the lithium ion kinetics in planar MoO 3−δ thin films. The chemical diffusion coefficient is found to vary between 10 −12 and 10 −15 cm 2 s −1 in organic electrolyte. D Li + differs by less than one order of magnitude for amorphous film with respect to crystallized films. This gives evidence that the diffusion of Li + ions changes importantly with the oxygen deficiency. For MoO 3 cathode films prepared using flash-and thermal-evaporation techniques at 250 • C, lithium diffusion coefficients are in the range 10 −12 -10 −11 cm 2 s −1 [141]. Halalay et al. determined D Li of 10 −11 cm 2 s −1 in amorphous MoO 3 films (550 nm thick) using an optical method relying on the application of the Beer-Lambert law [142]. In contrast, a low value of 2.88 × 10 −21 cm 2 s −1 was reported for nano-composites formed using ultra-fine MoO 3 anchored in coal-based carbon fibers [143]. Ding et al. [144] fabricated MoO 3−δ nanorods (light blue color) through mechanical grinding of bulk powders, which show a diffusion coefficient of Li + ions of 3.2 × 10 −14 cm 2 s −1 determined from an EIS experiment. The same amplitudes (1.35 × 10 −14 cm 2 s −1 ) were reported for MoO 3 nanobelts as anode materials [145]. MoO 3 /amorphous carbon composite fabricated by calcinating polyaniline with ammonium heptamolybdate tetrahydrate exhibited a low Li diffusion coefficient of 3.4 × 10 −14 cm 2 s −1 at room temperature [146]. Sun et al. [147] stated that oxygen vacancies in MoO 2.86 nanobelts promote a lower energy barrier for Li + diffusion but did not provide values. On the other hand, the oxygen deficient Mo 4 O 11 (MoO 2.75 ) exhibits better transport properties with an average lithium chemical diffusion coefficient of~4.5 × 10 −12 cm 2 s −1 [148]. These data can be compared with the previously reported values for metal-oxide thin films.  [151]. From the results listed in Table 4, listing D Li + observed for MoO 3−δ -based thin film anodes, it is obvious that the apparent diffusion coefficients are strongly dependent on the structure and morphology of MoO 3−δ thin films. Generally, bulk materials are composed of secondary particles (agglomerates), which are interconnected nanoparticles (primary particles) forming a mesoporous architecture. The mesopores (few tens of nm in size) are the intraconnecting voids formed between randomly packed nanoparticles favoring the high surface area between the active material and the electrolyte and enhancing the transport pathway. In contrast, thin films are composed of grains linked through grain boundaries, which are amorphous in nature. The absence of mesoporosity makes Li ion transport into the crystallite (29 nm in size) of the MoO 3−δ film more difficult. Thus, the morphology (mesoporosity, surface area, surface roughness) is the key parameter for transport of Li + ions into the active material. Moreover, the apparent diffusion coefficient D Li in thin films depends on the thermodynamic factor W F , which differs from W B of the bulk [99]. For films deposited at T s < 400 • C, the variation of D Li with the electrode potential exhibits a "W"-type behavior with two minimum regions at ca. 0.5 and 1.9 V, which are related to the strong attractive interactions between Li + ions and the host lattice [152]. This behavior is quite consistent with the one observed in Li x Si anode [153]. As MoO 3−δ thin films are oxygen-deficient materials, the model of charge transport in internal defect material can be applied [154]. Defects are Li-interstitials, Li*, and conduction electrons, e', for example. It has been shown that, in a solid solution where no internal defect reactions occur, the thermodynamic factor is related to the defect concentration (if dilute defects exist). The large increase of W may be also associated with the decrease of the electronic mobility in Li x MoO 3 film. The diffusion coefficients of Li + ions obtained in PLD-prepared MoO 3−δ thin films can be compared with other anode materials [141,[155][156][157][158][159][160][161][162][163] listed in Table 5. For cathode films prepared using flash evaporation (0.5-0.6 µm thick), D Li values are strongly dependent on T s and vary in the range 10 −12-10 −11 cm 2 s −1 [141]. Zhao et al. [124] investigated the Li + -ion kinetics in porous MoO 3 anode films (150 nm thick) prepared using electrodeposition in acid peroxo-polymolybdate electrolyte.
The Li + diffusion of 7.1 × 10 −11 cm 2 s −1 is indeed facilitated by the widespread nanoholes (several tens of nm) and by the small grains on the scale of~100 nm. PLD thin film 6.0 × 10 −14 -6.0 × 10 −12 this work

Conclusions
In this work, we showed that MoO 3−δ thin films were successfully fabricated using a PLD technique with good control of the oxygen deficiency. To the best of our knowledge, this is the first report on the correlation between the nanoscale structure of MoO 3−δ thin films and their electrochemical properties as anode in lithium microbatteries. The structural properties investigated using XRD, RS and FTIR spectroscopies showed that crystallized films can be formed with a wide range of deviation from stoichiometry (δ). The formation of lattice defects can be explained in terms of oxygen re-evaporation from the surface of the substrate during the PLD process. As-deposited Mo-O films at T s = 25 • C are amorphous in nature and highly oxygen deficient. The decomposition of the XPS spectra reveals that the Mo 3d spectra can be well fitted by two 3d doublets in the form of a Gaussian function, corresponding to Mo in +6 and +5 oxidation states. The atomic ratio O/Mo shows a small amount of oxygen vacancies in MoO 3−δ films deposited at T s ≈ 200 • C; the Mo sub-states were suppressed, resulting in a further oxidation of the film. The O/Mo ratio reaches the value of 2.982 when the temperature was increased to 300 • C, and a nearly stoichiometric α-MoO 3 film resulted. Further increase of T s (≈ 450 • C) produces a film with O/Mo = 2.894, which corresponds to the composition 0.85 Mo 8 O 23 -0.15MoO 3 . The electrochemical test of the 0.15MoO 3 -0.85Mo 8 O 23 film shows a specific capacity of 484 µAh cm −2 µm −1 after 100 cycles of charge-discharge at a constant current of 0.5 A cm −2 in the potential range 3.0-0.05 V. So far, it is the best result obtained with MoO 3−δ thin films. Studies of the lithium transport in MoO 3−δ thin film electrodes have shown that D Li is influenced by the nanostructure, morphology, grain size and gain boundaries. The formation of well-defined channels between slabs of the MoO 3 structure provides facile ionic conduction pathways in crystallized films. These channels are ill-defined in films deposited at low temperatures. This is a typical trend between crystalline and disordered networks. For instance, a value D Li = 6 × 10 −12 cm 2 s −1 is obtained for PLD films deposited at 300 • C.
This work has shown the advantage of PLD films free of conducting additive and binder commonly used in the fabrication of electrodes (i.e., 10% carbon black and 10% PVdF), as the use of these electrochemically inactive substances degrades the specific gravimetric capacity of the electrode.
On other hand, the PLD technique has shown unique advantages for the formation of dense films, easy control of the growth rate and production of high purity films with good preservation of the target-phase stoichiometry. The main drawback of PLD films is their limited capacity due to the relative difficulty to grow thick films with good adhesion (as a consequence of the low energy of few eV of molecules falling on the substrate). However, some improvements can be expected by the fabrication of mixed electrodes (for example MoO 3−δ + MoO 2 ) or blended electrodes (for example TiO 2 + MoO 3−δ ), in which the electronic conductivity can be adjusted.