Oxygen vacancies enhance pseudocapacitive charge storage properties of MoO 3-x

In the format provided by the authors and unedited. DOI: 10.1038/NMAT4810 Oxygen vacancies enhance pseudocapacitive charge storage properties of MoO 3-x Hyung-Seok Kim, 1 John B. Cook, 2,3 Hao Lin, 1 Jesse S. Ko, 1 Sarah H. Tolbert, 1,2,3* Vidvuds Ozolins, 1* and Bruce Dunn, 1,3,* 1 Department of Materials Science and Engineering, UCLA, Los Angeles, California 90095-1595, United States 2 Department of Chemistry and Biochemistry, UCLA, Los Angeles, California 90095-1569, United States 3 The California NanoSystems Institute, UCLA, Los Angeles, California 90095 * to whom correspondence should be addressed: bdunn@ucla.edu; tolbert@chem.ucla.edu; vidvuds@ucla.edu NATURE MATERIALS www.nature.com/naturematerials 1

Figure S1. TEM images of (a) reduced MoO3-x, and (b) fully oxidized MoO3. (Inset of (a) shows the blue color of as-prepared reduced MoO3-x) NATURE MATERIALS www.nature.com/naturematerials 2

Figure S2. AFM analysis of reduced MoO3-x. The images (top) and height profile (bottom) show that the thickness of an individual nanobelts is about 15 nm. NATURE MATERIALS www.nature.com/naturematerials 3

Figure S3. TGA analysis of reduced MoO3-x in air and argon atmosphere. (Flow of 100 ml/min and a ramping rate of 10 C/min). The concentration of oxygen vacancies were calculated from the difference in weight decrease between the two TGA traces. 100.0 99.5 Weight (%) 99.0 98.5 98.0 97.5 1.36% 400 C in Argon 97.0 400 C in Air 96.5 50 100 150 200 250 300 350 400 Temperature ( C) NATURE MATERIALS www.nature.com/naturematerials 4

Table S1. Surface area measurement on reduced MoO3-x and fully oxidized MoO3. Reduced MoO3-x Fully oxidized MoO3 BET surface area (m 2 g -1 ) 24 15 NATURE MATERIALS www.nature.com/naturematerials 5

Discussions of HSE Vs. GGA+U in MoO3 It is widely known that local and semi-local xc functionals (such as LDA and GGA) fail in systems with localized electrons electrons, such as transition metal compounds, due to incomplete of cancellation of electron self-interaction effects 1. The +U method is the simplest approach to account for on-site Coulomb interactions in a mean-field manner; GGA+U has been shown to calculate the redox potentials accurately 2. HSE06 is arguably more sophisticated because it is parameter-free and calculates the corrections self-consistently. Nevertheless, it still a rather ad hoc mixture of DFT and Hartree-Fock with screening and has had mixed success in transition metal compounds. For instance, Ong et al. showed that HSE is important when the polaron lives in hybridized transition metal d - Oxygen p orbitals, in which case GGA+U cannot localize the polaron because the +U is applied to the d orbital only 3. However, in the same paper it is shown that HSE predicts the wrong phase diagram for olivines, while GGA+U is better. In MoO3, the polaron lives almost entirely in the d orbital. In our previous work 4, we calculated polaron energies and migration barriers in stoichiometric MoO3 using GGA+U and HSE; the predictions are comparable and found to be in good agreement with experiment. We expect that the same conclusion will hold for reduced MoO3-x because the localization of the polaron is still confined to the Mo d orbitals. Therefore, we applied GGA+U method in the research of this paper. NATURE MATERIALS www.nature.com/naturematerials 6

Discussions of polaron configurations, polaron binding energies and variance of bond lengths in reduced MoO 3-x. Oxygen vacancy was modeled by removing one oxygen atom in a 3 1 3 supercell containing 36 formula units of stoichiometric MoO3. The preference site of oxygen vacancy in our results is consistent with previous GGA+U calculations, but the most stable polaron state is different. Coquet et al. 5 found that the structure with a terminating oxygen vacancy and with the reduced MoO3-x center in the Mo state is the most stable structure, which is the structure with the second lowest formation energy in our calculation, while in another study 6 the polaron state was not mentioned. We argue that our results should be more reliable because at least two polaron configurations were considered in each type of oxygen vacancy structure and previous calculations only showed one polaron configuration. Moreover, our prediction for polaron states is in good agreement with our experimental observation (Observation of Mo states) and our prediction for the polaron orbital (dxz) in the six-fold coordinated Mo ion is consistent with our previous study of fully oxidizedα-moo3 7. Since GGA methods cannot describe the localized d electrons in the transition metals correctly, it is not surprising that the preferred site of oxygen vacancy by GGA methods is the Os site, which contradicts the results from GGA+U methods 5. Furthermore, binding energies of the polarons were investigated by comparing the total energy difference of structures with various polaron configurations. If the bipolaron structure with an Ot vacancy is treated as a reference structure (see the polaron orbitals in Fig. 2c and the structure a in Fig. S4), the energies of the structures with a Mo cation at the defect center (structure b in Fig. S4) and with two adjacent Mo cations within the same bilayer are about 0.20 ev higher (structure c in Fig. S4). However, the separation of the bipolaron requires at least 0.37 ev. As shown in Structure d to f in Fig. S4, isolation of two polarons within the same NATURE MATERIALS www.nature.com/naturematerials 7

bilayer needs 0.37 ev, while separation of two polarons into two different bilayers demands 0.47 ev. Bipolaron structures are favored in reduced MoO3-x, because adjacent two polarons can minimize the total energy by relaxing the lattice distortion locally and this was confirmed by the analyzing the variance of bond lengths. Compared to the same kind of Mo-O bond lengths in stoichiometric MoO3, only 5 Mo-O bonds lengths in the reference structure are changed about 0.1 Å, while in structure 3 to 6, at least 8 Mo-O bond lengths are changed approximately 0.1 Å (See Fig. S5). As for structure c, although the distance of the two polarons is close, the two MoO6 octahedra containing the two polarons are edge sharing in the different layers, which is believed to be less flexible than the two polarons in the two adjacent corner-sharing octahedra in reference structure a. We point out that in the Mo structure, only one Mo-O bond length is changed more than 0.1 Å, but the repulsion between the two polarons in the same Mo cation center might lead to the increase of energy. NATURE MATERIALS www.nature.com/naturematerials 8

Figure S4. Binding energy of polarons with different polarons configuration (a-f).the configuration (a) used as the reference. Mo, Mo, Mo ions are highlighted in green, blue and white, respectively. NATURE MATERIALS www.nature.com/naturematerials 9

Figure S5. The difference of corresponding bond lengths between the reduced MoO3-x structures (a to f in Figure S5) and the stoichiometric MoO3. (Unit: Å) NATURE MATERIALS www.nature.com/naturematerials 10

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Figure S6. Long-term cycling of R-MoO 3-x. Long-term cycling measurements were performed using a three-electrode configuration. The working electrode was prepared as a nanoparticle thinfilm containing roughly 40 μg of active material (R-MoO3-x) with a coverage of 1 cm 2 onto an O2 plasma treated stainless steel current collector. The counter and reference electrodes were lithium metal foil and the electrolyte was 1M LiClO4 in propylene carbonate. The R-MoO3-x electrodes contained no carbon or binder additives so that the fundamental long-term cycling behavior of the material was evaluated. The electrodes were cycled between cut-off voltages of 2 2.8 V vs. Li/Li + using galvanostatic charging-discharging (at 30C) and cyclic voltammetry (at 10 mvs -1 ) for a total of 10000 cycles. The reduced MoO3-x exhibited excellent cyclability. At 30C, galvanostatic (a, b) charge-discharge curves indicate 70% capacity retention, decreasing to a very respectable value of 90 mah g -1 at 10,000 cycles. Cyclic voltammetry (CV) experiments (c, d) indicate capacity retention of 76% over 10,000 cycles. These scans also show that the capacity loss with cycling is associated with a decrease in the diffusion limited current peaks in the CV (compare Fig. S6c with Fig. 3e). This suggests that the capacitor-like currents in the CV cycle well while the diffusion-controlled currents cannot keep up with high rate cycling (30C or 10mVs -1 ). For both the potentiostatic and galvanostatic measurements, the coulombic efficiency was over 99%. NATURE MATERIALS www.nature.com/naturematerials 12

2.8 a 800 b 100 Voltage (V vs. Li/Li + ) 2.6 2.4 2.2 2.0 Galvanostatic Cycling 30C; 2-2.8 V vs. Li/Li + Cycle 10 Cycle 100 Cycle 1000 Cycle 10000 0 20 40 60 80 100 120 140 Capacity (mah/g) Capacity (C/g) 700 600 500 400 300 200 100 0 Charge Discharge Galvanostatic Cycling 0 2000 4000 6000 8000 10000 Cycle Number (n) 80 60 40 20 0 Coulombic Efficiency (%) 0.4 0.22 V c Cycle 10 700 d 100 Current (ma) 0.3 0.2 0.1 0.0-0.1-0.2-0.3 Cyclic Voltammetry 10 mv/s; 2-2.8 V vs. Li/Li + 0.05 V Cycle 100 0.03 V Cycle 10000 Cycle 10 Cycle 100 Cycle 1000 Cycle 10000 1.8 2.0 2.2 2.4 2.6 2.8 3.0 Voltage (V vs. Li/Li + ) Capacity (C/g) 600 500 400 300 200 100 0 Charge Discharge 0 2000 4000 6000 8000 10000 Cycle Number (n) Cyclic Voltammetry 80 60 40 20 0 Coulombic Efficiency (%) Capacitance (Fg -1 ) 500 400 300 200 100 0 Reduced MoO 3-x Fully Oxidized MoO 3 0 20 40 60 80 100 Sweep rate (mvs -1 ) Electrochemical analysis of thin film MoO 3 electrodes. To complement the results shown in Figure 3b, which considers the sweep rate dependence on capacity, we have plotted here the specific capacitance (Fg -1 ) as a function of sweep rate. The outcome is the same as we observe the R-MoO3-x (red) having a much greater specific capacitance than that of F-MoO3 (black) NATURE MATERIALS www.nature.com/naturematerials 13

Figure S7. Kinetic analysis of reduced MoO3-x and fully oxidized MoO3 using the power law relationship. The b-values of reduced 0.5 MoO3-x (red) and fully oxidized MoO3 (black) are calculated from the slope of the log (i) vs. log ( ) plot using the anodic current response at ~ 2.5 V (vs. Li/Li + ). log (peak current) (ma) 0.0-0.5-1.0-1.5-2.0 Anodic peak b = 0.85 b = 0.67 Reduced MoO 3-x Fully oxidized MoO 3-2.5-1.0-0.5 0.0 0.5 1.0 1.5 2.0 log (sweep rate) (mvs -1 ) NATURE MATERIALS www.nature.com/naturematerials 14

Figure S8.The intercalation energies of all symmetry-distinct lithium insertion sites were calculated using repeated random initialization of the polaron states. (a) Calculated lithium intercalation voltages at interlayer sites (cross) and intralyer sites (triangle) of reduced MoO3-x. The polarons introduced by an oxygen vacancy are shown in blue. A, B, C, D sites indicate the positions of an additional polaron introduced by lithium intercalation. For example, the A site means that the additional polaron is located at the A planes shown in (b). (a) (b) NATURE MATERIALS www.nature.com/naturematerials 15

Discussions regarding polaron configurations after lithium intercalation. As shown in Fig. S8, our results indicate that the additional polaron introduced by lithium intercalation does not favor the B plane, where the oxygen defect induced polarons are located. In other words polarons are unlikely to remain together as the polaron concentration increases from 0.22 to 0.33 within the same plane. The highest intercalation energy corresponds to the structure in which the lithium ion at the interlayer site is far from the oxygen vacancy center (6.3 Å) and the additional polaron introduced along with the lithium ion is not in the same layer as those polarons introduced by an oxygen vacancy (see Fig. S9). An explanation regarding the additional peak during the cathodic sweep The calculated voltage difference between reduced MoO3-x and fully oxidized MoO3 explain the new peak at 3.0 V in the CV of reduced MoO3-x (see Fig. 3e). Compared to lithium intercalation in fully oxidized MoO3, oxygen vacancies in the reduced MoO3-x might release the structural strain energy during the lithium intercalation and lead to higher voltages. Therefore, the new peak is only observed in the reduced MoO3-x. In the discharge process, this new peak is less distinctive than that in the charge process, because lithium ions might not migrate directly from the electrolyte to the high voltage sites in reduced MoO3-x.Instead, the weak signal of the new peak in the discharge process is contributed by the diffusion of lithium from the electrolyte to the high voltage sites existing near the interface between electrode and electrolyte. NATURE MATERIALS www.nature.com/naturematerials 16

Figure S9. The lithium intercalation sites for the highest lithium intercalation voltage. NATURE MATERIALS www.nature.com/naturematerials 17

Table S2. Powder XRD derived d-spacing values for the (020), (040), and (060) reflections. A silicon internal standard was used to calibrate the position of the peaks. Sample (020) d (Å) (040) d (Å) (060) d (Å) Reduced MoO 3-x 7.031 3.505 2.335 Fully Oxidized MoO 3 6.943 3.462 2.307 Table S3. Lattice parameters of reduced MoO3-x and fully oxidized MoO3 calculated by DFT when one oxygen vacancy is introduced. a(å) b(å) c(å) Fully oxidized MoO3 11.686 13.904 11.357 ReducedMoO 2.972 11.650 13.930 11.393 The deviation between the b lattice parameter used in the DFT calculations from the measured value for the fully reduced sample ( b=0.13å out of b=14å, or less than 1%) is small and will not affect the calculated intercalation energies significantly. At any rate, this difference is smaller than the uncertainty in the calculated van der Waals gap for stoichiometric MoO3 using different exchange-correlation functionals 8. NATURE MATERIALS www.nature.com/naturematerials 18

Electrical conductivity measurements of reduced MoO3-x and fully oxidized MoO3. Improved electrical conductivity in the reduced MoO3-x also contributes to the fast kinetic of reduced MoO3-x. The bulk electrical conductivity measurement is illustrated in Fig. S9.From this setup the resistivity was measured and the electrical conductivity of each sample was calculated from following equation: (1) k is electrical conductivity of material, t is sample thickness, A is contact area, and R is resistivity of sample. Table S4 shows that electrical conductivity of reduced MoO3-x was an order of magnitude higher than fully oxidized MoO3. These measured values are in reasonable agreement with the reported electrical conductivity of a single MoO3 nanobelt (10-4 S/cm for non-lithiated MoO3 nanobelt 9 ). The conductivity of our sample was measured as a thick film with thickness of 10 μm and a 1 cm 2 area, so these values represent the upper bound of the intrinsic conductivity due to contact resistance. NATURE MATERIALS www.nature.com/naturematerials 19

Figure S10. Experimental setup for bulk electrical conductivity measurement of reduced MoO3- x, partly reduced, and fully oxidized MoO3. MoO 3 film by drop casting (thickness: ~10 μm) ITO coated glass Etched area Silver paste (contact area: 0.5 cm 2 ) Table S4. Calculated bulk electrical conductivity using equation (S1) in each sample. Samples Electrical conductivity Reduced MoO3-x 10-4 S/cm Fully oxidized MoO 3 10-5 S/cm NATURE MATERIALS www.nature.com/naturematerials 20

Figure S11. Ex-situ XPS spectra at specific potentials after electrochemical cycling (at 5 mvs -1 ). Cyclic voltammetry in (a) reduced MoO3-xand (b) fully oxidized MoO3and high resolution XPS spectra of Mo 3d region at specific potentials (i v). (a) Current (ma) 10 Reduced MoO 8 3-x 6 4 2 5 mvs -1 (iv) 2.8 V (v) 3.5 V 0-2 (i) 3.5 V -4 (iii) 1.5 V -6 (ii) 2.5 V 1.5 2.0 2.5 3.0 3.5 Potential (V vs Li/Li + ) Intensity (a.u.) Intensity (a.u.) (i) 3.5 V (ii) 2.5 V (iii) 1.5 V (iii) 1.5 V (iv) 2.8 V (v) 3.5 V (b) Current (ma) 6 4 2 0-2 -4 Fully oxidized MoO 3 5 mvs -1 (iii) 1.5 V (iv) 2.8 V (v) 3.5 V (ii) 2.5 V (i) 3.5 V Intensity (a.u.) Intensity (a.u.) (i) 3.5 V (ii) 2.5 V (iii) 1.5 V (iii) 1.5 V (iv) 2.8 V (v) 3.5 V 1.5 2.0 2.5 3.0 3.5 Potential (V vs Li/Li + ) NATURE MATERIALS www.nature.com/naturematerials 21

References 1. Cheng, J., Sulpizi, M., VandeVondele, J. & Sprik, M. Hole localization and thermochemistry of oxidative dehydrogenation of aqueous rutile TiO2(110). Chemcatchem 4, 636-640 (2012). 2. Zhou, F. et al. First-principles prediction of redox potentials in transition-metal compounds with LDA + U. Phys. Rev. B 70, (2004). 3. Ong, S. P., Chevrier, V. L. & Ceder, G. Comparison of small polaron migration and phase separation in olivine LiMnPO4 and LiFePO4 using hybrid density functional theory. Phys. Rev. B 83, (2011). 4. Ding, H. et al. Computational investigation of electron small polarons in alpha-moo3. J. Phys. Chem. C 118, 15565-15572 (2014). 5. Coquet, R. & Willock, D. J. The (010) surface of α-moo3, a DFT + U study. Phys. Chem. Chem. Phys. 7, 3819-3828 (2005). 6. Lei, Y.-H. & Chen, Z.-X. DFT+U study of properties of MoO3 and hydrogen adsorption on MoO3(010). J. Phys. Chem. C 116, 25757-25764 (2012). 7. Ding, H. et al. Computational investigation of electron small polarons in alpha-moo3. J. Phys. Chem. C 118, 15565-15572 (2014). 8. Ding, H., Ray, K. G., Ozolins, V. & Asta, M. Structural and vibrational properties of alpha-moo3 from van der Waals corrected density functional theory calculations. Phys. Rev. B 85, (2012). 9. Mai, L. et al. Lithiated MoO3 nanobelts with greatly improved performance for lithium batteries. Adv. Mater. 19, 3712-3716 (2007). NATURE MATERIALS www.nature.com/naturematerials 22