MnO2 batteries: Effect of manganese dioxide properties on electrochemical performance

Electrochimica Acta 105 (2013) 305–313

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Heat treated electrolytic manganese dioxide for primary Li/MnO2 batteries: Effect of manganese dioxide properties on electrochemical performance Wesley M. Dose, Scott W. Donne ∗ Discipline of Chemistry, University of Newcastle, Callaghan, NSW 2308, Australia

a r t i c l e

i n f o

Article history: Received 4 February 2013 Received in revised form 27 March 2013 Accepted 4 April 2013 Available online 12 April 2013 Keywords: Lithium battery Manganese dioxide Heat treatment Energy storage

a b s t r a c t The primary capacity of heat treated manganese dioxide in Li/MnO2 batteries is directly related to the numerous physical properties which characterize this material. A statistical model is employed to isolate the influence of material structure, composition and morphology on the electrochemical performance at the discharge rates 2, 5, 10 and 20 mA g−1 . Among the most influential parameters are the pyrolusite content, Mn(IV) percentage, cation vacancy fraction and surface area. Some of the materials investigated show higher specific capacities than literature materials, with further improvement to the performance of these materials anticipated through the intelligent selection of heat treated manganese dioxide which exhibit the ideal properties outlined in this study. © 2013 Published by Elsevier Ltd.

1. Introduction The majority of portable consumer batteries produced worldwide use manganese dioxide as the cathode material. While the aqueous Zn/MnO2 system holds the dominant position in the portable battery market, since the 1980s non-aqueous Li/MnO2 battery systems have found popular application in high energy density and high power consumer electronics [1]. The Li/MnO2 battery system has many advantages over its counterparts, most commonly lithium cobalt oxides or lithium nickel oxides. Based on cost and abundance, manganese dioxide is a much more viable material in supporting the growing consumer power demands. Additionally, manganese dioxide is a much safer material to use in these cells, while also providing reasonable electrochemical output [2]. The structure of manganese dioxide commonly used in the Zn/MnO2 system is prepared via electrodeposition, and hence termed electrolytic manganese dioxide (EMD) [3]. EMD must be dehydrated by thermal treatment (usually between 250 and 400 ◦ C) prior to use in the Li/MnO2 system. This process not only removes problematic water content, but is also thought to create a structure more suitable for lithium intercalation [4–6]. Whether EMD or heat treated EMD (HEMD), the structure of these -MnO2 compounds

∗ Corresponding author. Tel.: +61 2 49215477; fax: +61 2 49215472. E-mail address: [email protected] (S.W. Donne). 0013-4686/$ – see front matter © 2013 Published by Elsevier Ltd. http://dx.doi.org/10.1016/j.electacta.2013.04.011

consists of a random intergrowth of pyrolusite and ramsdellite, where the fraction of pyrolusite is denoted Pr [7]. The structure is riddled with defects in the form of manganese vacancies, lower valent manganese ions (Mn(III)), and structural water (present as protons) which are associated with these defects [8–10]. The effect of heat treatment on EMD has been studied extensively [7,5,11–13], but in brief, heat treatment converts most of the ramsdellite domains in the initial ramsdellite/pyrolusite intergrowth to pyrolusite (i.e., Pr increases), while also annealing defects and the associated structural water. Since the first successful report of manganese dioxide as a cathode material in primary lithium batteries, a number of authors have considered the discharge mechanism of these cells [4,14–16,5]. The most recent model was outlined by Bowden et al. [6], who found the discharge process could be divided into three regions. The first, and shortest, step is said to relate to the homogenous lithiation of residual ramsdellite regions in the HEMD. In the second region of the discharge they proposed that lithium insertion into the pyrolusite domains results in a slow transformation to an expanded ramsdellite lattice structure, which is then rapidly reduced in the third and final discharge process. The XRD data presented in the work of these authors suggests the HEMD investigated was highly pyrolytic. The mechanism for a sample with a greater retention of the parent -MnO2 structure (achieved by lithiation of the EMD prior to heat treatment) has been reported to take place in two reversible steps [17]. First, lithium is inserted into the ramsdellite structure to form a mixed valence Mn(III/IV) product, which is

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then further reduced to Mn(III). This is followed by the lithiation of pyrolusite domains in a mechanism essentially the same as that suggested for HEMD. The dependence of the discharge mechanism on the material properties highlights the importance of selecting the most suitable EMD/HEMD to achieve optimum performance. This can only be made possible once concrete relationships have been forged between material characteristics and electrochemical performance. Despite the significance of such a result, surprisingly few reports of this nature appear in the literature. The initial report by Ikeda et al. [4] found the /ˇ-MnO2 and ˇ-MnO2 structures most suitable for lithium intercalation. Following on from this, Pistoia [18] showed that the cathode efficiency increases until the /ˇ-form produced upon moderate heating is transformed into a ˇ-like form, although the latter was found to be still more active than the -form. Ilchev et al. [19] later determined the most common optimum heating temperature for a selection of EMD samples was 300 ◦ C for 8 h, while longer times led to poorer utilization. A connection between high specific surface area and porosity and high specific capacity has also been reported by Ilchev and Banov [20]. Research by Sarciaux, Guyomard and co-workers [21,22] represents the most comprehensive reported study of this kind to date. They considered the effects of oxygen content (i.e., x in MnOx ), pyrolusite fraction (Pr ) and microtwinning on the material capacity and structural modification during lithium insertion. The structural transformation is found to be more difficult for materials with higher Pr and microtwinning, indicated by a lower thermodynamic potential and slower first order kinetics. Intuitively, HEMD with higher oxygen content (indicating a low defect concentration and higher manganese oxidation state) were found to have a greater capacity. However, the diversity that exists among materials which can be classified as EMD is brought about by considerably more material properties than those considered in these reports. EMD can vary in: (i) structure – through pyrolusite content, unit cell volume, and crystallite size, (ii) composition – oxygen content by varying amounts of Mn(IV), Mn(III) and cation vacancies, and (iii) morphology – quantifiable through BET surface area. All these characteristics will affect the electrochemical performance of the manganese dioxide cathode in a Li/MnO2 cell and therefore must be considered in a thorough analysis of this system. Consequently, there is considerable scope for battery optimization by the informed selection of precursor EMD materials that result in superior performing heat treated materials in Li/MnO2 batteries. This first requires two key relationships to be understood; i.e., (i) how HEMD properties are affected by the properties of their precursor EMD materials, and (ii) how the properties of HEMD influence the electrochemical performance. To address the first relationship, we have demonstrated that the properties of HEMD materials can be linked to key properties in the precursor EMD used [23]. However, to our knowledge there has been no comprehensive study relating the properties of HEMD to the electrochemical performance, despite this battery system being used commercially for the last 30 years. The initial findings from this research have been reported in communication form in [24], in which relationships between the electrochemical capacity of MnO2 and pyrolusite content, percentage Mn(IV) and BET surface area are briefly outlined. This report will discuss in detail how the numerous other material properties (pyrolusite content, unit cell volume, average crystallite size, Mn(IV) percentage, Mn(III) percentage, cation vacancy fraction, and BET surface area) contribute to the performance of manganese dioxide in Li/MnO2 batteries. This report is, to our knowledge, the first of its kind in the literature and as such represents a fundamental step in the optimization of HEMD for use in these batteries.

2. Experimental 2.1. Synthesis of -MnO2 The samples of -MnO2 used in this work were prepared by anodic electrodeposition, and hence termed EMD. The cell used for electrolysis is based on a temperature controlled 2 L glass beaker in which two 144 cm2 (72 cm2 on either side) titanium sheets were used as the anode substrate, and three similarly sized copper sheets were used as the cathode substrate. The electrodes were arranged alternately so that each anode was surrounded on both sides by a cathode. The electrolyte was an aqueous mixture containing 0.05–0.5 M MnSO4 (APS Chemicals, 98%) and 0.01–1.0 M H2 SO4 (Lab-Scan, 98%) maintained at an elevated temperature. Electrodeposition of the manganese dioxide was conducted with an anodic current density within the range 20–100 A m−2 according to the reactions: Anode(Ti) : Cathode(Cu) : Overall :

Mn2+ + 2H2 O −→ MnO2 + 4H+ + 2e− 2H+ + 2e− −→ H2

Mn2+ + 2H2 O −→ MnO2 + 2H+ + H2

(1) (2) (3)

The overall process was carried out for three days, during which time the electrolyte Mn2+ concentration was of course depleted, while the H+ concentration increased. To counteract this, and hence maintain a constant electrolyte composition over the duration of the deposition, a concentrated (1.5 M) MnSO4 solution was added continually at a suitable rate to replenish Mn2+ and dilute any excess H2 SO4 formed. Under these conditions control of the solution conditions was typically maintained to within ±2%. After deposition was complete, the solid EMD deposit was mechanically removed from the anode and broken into chunks ∼0.5 cm in diameter and then immersed in 500 mL of Milli-Q water to assist in the removal of entrained plating electrolyte. The pH of this chunk suspension was adjusted to pH 7 with the addition of 0.1 M NaOH (Sigma–Aldrich, 98%). After ∼24 h at a pH of 7 the suspension was filtered and the chunks then dried at 110 ◦ C. After drying, the chunks were milled to a −105 ␮m powder (mean particle size ∼45 ␮m) using an orbital zirconia mill. The powder was then suspended in ∼500 mL of Milli-Q water and its pH again adjusted to 7 with the further addition of 0.1 M NaOH. When the pH had stabilized, the suspension was filtered and the collected solids dried at 110 ◦ C. When dry the powdered EMD was removed from the oven, allowed to cool to ambient temperature in a desiccator and then transferred to an airtight container for storage. 2.2. Material characterization The structure of the materials used in this work were all confirmed to be -MnO2 by X-ray diffraction using a Phillips 1710 diffractometer equipped with a Cu K˛ radiation source ( = 0.15418 nm) and operated at 40 kV and 30 mA. The scan range was from 10 ◦ to 80◦ 2, with a step size of 0.05◦ and a count time of 2.5 s. Morphology was examined by gas adsorption using a Micromeritics ASAP 2020 Surface Area and Porosity Analyser. A representative 0.10 g sample of the manganese dioxide material was degassed under vacuum at 110 ◦ C for 2 h prior to analysis. An adsorption isotherm was then determined over the partial pressure (P/P0 ) range 10−7 –10−1 using N2 gas as the adsorbate at 77 K. The specific surface area was extracted from the gas adsorption data using the linearized BET isotherm in the range 0.05 < P/P0 <0.30, while the pore size distribution was determined using a Density Functional Theory-based approach (Micromeritics DFTPlus V2.00).

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Table 1 Material properties of starting EMD and heat treated EMD samples. Synthesis a ◦

M1

M2

M3

M4

M5

a b c d e f

T / C

Time/h

25 200 250 300 350 400 25 200 250 300 350 400 25 200 250 300 350 400 25 200 250 300 350 400 25 200 250 300 350 400

– 13.12 12.38 7.15 1.48 – – 22.62 13.97 3.53 0.33 – – 25.33 12.95 4.68 1.65 0.60 – 14.33 7.12 3.68 2.02 0.68 – 25.90 21.38 5.90 1.50 0.53

Pr b

UCVc /nm3

td /nm

wt% Mn(IV)

wt% Mn(III)

CVFe

BET SAf /m2 g−1

0.39 0.45 0.49 0.78 0.80 – 0.39 0.66 0.73 0.74 0.73 – 0.43 0.65 0.71 0.73 0.76 0.77 0.35 0.65 0.72 0.70 0.75 0.76 0.41 0.64 0.69 0.76 0.81 0.80

0.1194 0.1177 0.1173 0.1150 0.1143 – 0.1199 0.1177 0.1168 0.1159 0.1159 – 0.1185 0.1163 0.1148 0.1148 0.1142 0.1134 0.1200 0.1182 0.1164 0.1159 0.1142 0.1134 0.1200 0.1173 0.1168 0.1155 0.1135 0.1135

22.12 20.85 20.60 16.97 18.74 – 16.89 15.28 16.64 17.35 17.11 – 17.28 16.97 14.15 15.95 17.58 19.98 14.16 12.40 16.28 11.41 14.10 14.26 15.36 17.46 15.82 14.14 23.23 22.90

56.3 57.3 60.0 54.9 54.0 – 53.4 51.2 53.9 54.5 53.2 – 57.6 56.5 60.5 53.6 56.6 56.7 59.0 58.0 59.7 57.9 58.2 57.2 58.6 61.2 61.0 61.7 59.8 59.7

3.07 3.99 2.43 5.60 6.21 – 5.58 6.28 8.97 6.97 9.20 – 2.14 2.67 0.81 0.82 4.18 4.43 0.05 0.95 2.74 3.18 3.34 3.91 1.85 0.13 0.09 0.63 1.48 1.89

0.069 0.040 0.026 0.001 0.000 – 0.075 0.044 0.014 0.015 0.000 – 0.087 0.052 0.048 0.041 0.016 0.000 0.100 0.061 0.043 0.030 0.025 0.028 0.082 0.058 0.038 0.025 0.011 0.019

19.0 19.8 18.8 18.3 14.7 – 45.4 42.7 44.3 39.1 42.9 – 43.1 42.8 39.5 44.6 39.0 39.2 66.1 68.6 67.3 62.7 57.6 58.4 31.0 24.9 24.8 21.6 20.9 24.9

Heating temperature used for thermal synthesis. Pr : pyrolusite fraction. UCV: unit cell volume. t: average crystallite size. CVF: cation vacancy fraction. BET SA: BET surface area.

The composition of the materials was determined using two consecutive potentiometric titrations as outlined in Vogel [25]. This procedure and the method for determining the values for (i) x in MnOx , (ii) percentage of Mn(IV) and (III), and (iii) cation vacancy fraction (CVF), has been described in considerable detail in our previous work [26] and so will not be repeated here. 2.3. Thermogravimetric analysis (TGA) Kinetic analysis was based on TG experiments conducted using a PerkinElmer Diamond TG/DTG controlled by Pyris software. Approximately 10 mg of -MnO2 sample was added to an aluminium sample pan and placed into the analyzer. The same mass of ˛-Al2 O3 in a similar aluminium pan was used as the reference material for DTA measurements. The heating profile applied to the sample under a static air atmosphere was a linear ramp from ambient temperature to 600 ◦ C at 1.0 ◦ C min−1 . 2.4. Thermal treatment of EMD Approximately 10 g of EMD was heated in an alumina boat crucible by a Eurotherm HTC1400 furnace with a static air atmosphere set at the required temperature. After the elapsed isothermal heating time, the sample was removed from the oven and allowed to cool to room temperature. 2.5. Electrochemical characterization The electrochemical properties of the samples were determined using CR2032 coin-type cells constructed in a dry argon filled glove

box. The cathode was fabricated by compressing a 1:8:1 mixture (by weight) of EMD:graphite:binder (polyvinylidene fluoride) into a pellet 1 cm in diameter and ∼1 mm thick, which was subsequently dried at 110 ◦ C under vacuum for 12 h and then transferred to the glove box. Lithium foil served as the anode, while a Celgard 2400 sheet was employed as a separator. The electrolyte was a solution of 1 M LiPF6 in 1:1 (w/w) ethylene carbonate:dimethyl carbonate. The cells were tested at room temperature by a PerkinElmer VMP2 multi-channel potentiostat using galvanostatic mode. 3. Results and discussion 3.1. Material properties Five EMD samples (M1–M5) exhibiting suitable variation in composition, structure and morphology were prepared. The properties of these materials are listed in Table 1. As demonstrated in this table, there exists significant variation (although still within the bounds for EMD [27]) between the five samples selected in structure (Pr , 0.35–0.43; unit cell volume, 0.118–0.120 nm3 ; average crystallite size, 14.2–22.1 nm), composition (Mn(IV), 53.4–59.0%; Mn(III), 0.05–5.6%; cation vacancy fraction, 0.07–0.10), and morphology (surface area, 19–66 m2 g−1 ). Using the method of kinetic analysis based on thermogravimetric experiments outlined in detail in our earlier reports [28,26,29], the isothermal heating time necessary to theoretically completely remove water from the -MnO2 structure was calculated at a set of pre-defined temperatures for each EMD sample. These heat treatment regimes and the properties of the 5 EMD and resultant 23 HEMD samples are listed in Table 1. As expected, there exists

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Fig. 1. Representative discharge curves for HEMD with respect to (a) thermal treatment temperature (M5 discharged at 2 mA g−1 ), and (b) discharge rate (M5 heat treated at 250 ◦ C).

considerable variation of the HEMD properties, not only with respect to heat treatment temperature, but also depending on the properties of the starting EMD. 3.2. Electrochemical performance Representative discharge curves of the HEMD materials are shown in Fig. 1. The primary capacity of the 28 EMD/HEMD materials in Li/MnO2 2032 coin cells was used as the indicator for electrochemical performance. This value was determined for each material at four rates; 2, 5, 10 and 20 mA g−1 of MnO2 , with a cutoff cell potential of 2.0 V. Table 2 lists the capacities measured (in mAh g−1 ) for the HEMD materials at the four discharge rates. On average the capacity for an unheated EMD was 105 mAh g−1 (or 34% Table 2 Electrochemical performance of starting EMD and heat treated EMD samples. The lowest and highest performing materials at each discharge rate are indicated by italic and bold font, respectively. Temperature/◦ C

M1

M2

M3

M4

M5

25 200 250 300 350 400 25 200 250 300 350 400 25 200 250 300 350 400 25 200 250 300 350 400 25 200 250 300 350 400

Primary capacity/mAh g−1 2 mA g−1

5 mA g−1

10 mA g−1

20 mA g−1

81.1 112.5 189.0 212.6 211.4 – 110.8 199.5 243.1 235.2 234.4 – 72.1 155.6 213.1 217.3 204.9 204.8 88.4 164.6 217.3 224.7 234.4 235.6 177.2 251.5 265.1 258.5 245.3 248.6

62.0 99.0 161.7 214.4 203.3 – 91.1 187.3 222.8 226.0 223.6 – 66.9 139.6 211.9 216.5 201.7 217.1 82.2 159.6 188.1 213.6 231.5 218.6 135.5 253.3 251.0 256.3 251.9 265.8

42.6 59.4 98.3 142.0 149.7 – 84.2 198.7 220.8 234.6 231.4 – 56.5 89.1 202.3 206.2 167.2 202.9 70.9 130.4 183.8 208.3 221.9 227.0 99.8 240.1 243.4 244.2 246.8 248.5

40.2 33.2 79.6 145.0 116.5 – 74.0 171.1 198.5 218.2 223.2 – 48.0 97.3 180.2 179.0 179.0 207.8 58.5 107.3 152.4 180.2 196.8 209.0 93.1 202.8 239.9 223.8 219.2 220.3

utilization), which was less than half that of HEMD, with an average of 215 mAh g−1 (70% utilization) at the 2 mA g−1 discharge rate. This clearly demonstrates the importance of the heat treatment, as has been reported in the literature [4]. This vast difference in performance also suggests that EMD based cells are failing in a different way to HEMD cells. The high water content of the EMD likely leads to destructive side reactions with the electrolyte and anode causing cell failure. HEMD’s, on the other hand, are much more likely to be limited by their intrinsic material properties. We therefore excluded the EMD materials from our subsequent analysis. Considerable variation in performance was noted between the HEMD materials, with capacities ranging from 113 to 265 mAh g−1 (37–86% utilization) at the 2 mA g−1 rate. The maximum and minimum capacities of the HEMD at each discharge rate have been highlighted in italics and bold text, respectively, in Table 2. The minimum capacities at each discharge rate all originate from a single EMD (M1), as do those with highest performance (M5). This highlights the importance of the properties of the precursor material on the performance of the heat treated material, relationships which we have investigated in detail in [23]. Further, the lowest performance is found for materials which have been heated at only mild temperatures, namely 200 ◦ C, indicating these temperatures are insufficient (even with the much longer heating time used) to generate HEMD materials with good electrochemical properties. Heating temperatures of 250 ◦ C and above, however, are able to create high performing HEMD materials. The differences in electrochemical performance can be attributed to the different structure, composition and morphology characteristics of the HEMD materials. 3.3. Statistical analysis To relate the electrochemical performance of each HEMD material to the numerous properties which could potentially be influencing it, a statistical model was employed. This method involved fitting an expression of the form: Performance = a0 +

n 

ai xik

(4)

i=1

where performance is the primary specific capacity, a0 is a constant, ai are a series of coefficients describing the extent of contribution of the variable xi to the performance, and k is a power term, either 1 or 2. The r2 value was used to determine the fitting as a result of the summation in Eq. (4). To represent the material in terms of structure, composition and morphology, the variables (xi ) used in this analysis were: pyrolusite content (Pr ), unit cell volume,

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Table 3 Coefficient of determination for data fitting. Discharge rate/mA g−1

r2

2 5 10 20

0.93 0.90 0.87 0.91

average crystallite size, Mn(IV) percentage, Mn(III) percentage, cation vacancy fraction, and BET surface area. The expression in Eq. (4) was used for each of the four discharge rates in order to ascertain how the contribution to performance varied with discharge rate for a given material parameter. Prior to use in our model, the variables, xik , were normalized (from −1 to 1 using

2(xik − min)/(max − min)−1) from the measured data to give an unbiased contribution for each parameter. Using this model involving the seven parameters and their squared values, a good fit of the calculated performance to the experimental data was achieved for each of the four discharge rates tested (r2 values are listed in Table 3). The individual contribution of a variable, xi , to the performance at a given discharge rate was then isolated by taking the sum of the related terms from Eq. (5), i.e.: Performance = ai xi + ai+1 xi2

Fig. 3. Influence of the unit cell volume on electrochemical performance.

(5)

where the symbols here have the same meaning as in Eq. (4). The individual contribution of the variable xi at the four discharge rates was then visualized by plotting the performance term from Eq. (5) with respect to the measured values for this parameter. The results of this analysis for the 23 HEMD materials are represented in Figs. 2–8, which show the relative contribution of a parameter to the electrochemical capacity of that material. It should be noted that the scales on the y-axis in Figs. 2–8 are equal and represent the relative contribution to performance. Further, performance values above the x-axis indicate a positive contribution to the capacity, while negative values show a detrimental effect. It is important to note that the method of analysis employed here has allowed us to determine the effect of an individual parameter on the electrochemical performance. The relationship for a given parameter is therefore independent of the other material features, unless two parameters are directly related in the way that the Mn(IV) and Mn(III) percentages are (i.e., inversely proportional). It follows then that combining all the relationships found in this study should point to the properties of HEMD which will exhibit superior capacity. It should be recognized however, that this ideal

Fig. 2. Influence of the pyrolusite content on electrochemical performance. Inset shows the optimum pyrolusite content with respect to discharge rate.

Fig. 4. Influence of the average crystallite size on electrochemical performance.

combination may in fact be intrinsically impossible. It is well known that the synthesis conditions in the production of EMD and the heat treatment conditions used to form HEMD both affect the properties of the final material [27,30]. These properties can often be seen to have some connection. For instance, a high deposition current

Fig. 5. Influence of the manganese(IV) content on electrochemical performance.

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3.4. Influence of structural parameters

Fig. 6. Influence of the manganese(III) content on electrochemical performance.

Fig. 7. Influence of the cation vacancy fraction on electrochemical performance.

for EMD forms a more disordered deposit, with smaller crystallites and high BET surface area. To electrodeposit a material with both high crystallinity and BET surface area is therefore an unlikely prospect.

Fig. 8. Influence of the BET surface area on electrochemical performance.

3.4.1. Pyrolusite content The effect of the HEMD structure, in terms of its pyrolusite content (Pr ), on the electrochemical performance is shown in Fig. 2. It is at once evident that this parameter is highly influential on the performance, more than twice that of other parameters at 2 and 5 mA g−1 discharge rates, and at least equally significant for 10 and 20 mA g−1 . This suggests that the material structure plays a more critical role for slower discharge. In addition, for the materials considered, the spread of contribution to performance is much greater at the lower rates than for high. This indicates that the particular structural arrangement is a more critical feature of the material at low rates. We also observe that for a given discharge rate there is an optimum pyrolusite content (shown in the inset of Fig. 2) corresponding to a particular structural arrangement. This begins at Pr = 0.65 for 2 mA g−1 and shifts towards slightly more pyrolytic structures for the higher discharge rates, up to Pr = 0.74 for the 10 and 20 mA g−1 rates. Therefore, structures either side of this, that is more ramsdellite-like structures and highly pyrolytic materials, are less constructive to good performance. This agrees well with the findings of Ikeda et al. [4] and Pistoia [18], but this study provides additional quantitative information as to the most appropriate structure for lithium intercalation, as well the dependance of this material feature on discharge rate. If all crystallites were single-phase (i.e., either ramsdellite or pyrolusite, with the increments in Pr determined by the size and number of each type) irrespective of heating temperature, then the material could only indicate whether a ramsdellite or pyrolusite phase is the preferred structure. Because our current findings indicate that a specific fraction of each phase is most beneficial, we may conclude that the crystallites of HEMD are multi-phase. This finding follows on from the work of Turner and Buseck [31], who used HRTEM to image the defects presents in the natural form of MnO2 , named nsutite. Their results showed some regions of regular alternating single and double chains of manganese octahedra, indicating a multi-phase crystal. Several other structural features were discovered, such as triple chains, semi-coherent intergrowths (creating large seams in the material), and the presence of todorokite, all of which will likely have bearing on the electrochemical features of the material. Considerable scope exists to examine the role of such defects on the mechanism of lithium diffusion into HEMD. At this stage however, the result found here (that a specific fraction of pyrolusite/ramsdellite is ideal) clearly indicates that the phase boundaries and interaction between phases within a crystallite play an important role in improving the electrochemical performance of HEMD for lithium batteries. It is interesting to compare our results with those reported by Balachandran et al. [32], who use a first principles computational study to compare proton diffusion through ramsdellite and pyrolusite structures. Their work demonstrated that protons diffuse through ramsdellite with a lower activation energy barrier than pyrolusite (∼200 meV compared to 575 meV, respectively). However, their model does not reveal how increments of Pr influence the ease of diffusion. While our method does not find the activation energies for the diffusion of lithium through various structures, it does show that neither ramsdellite nor pyrolusite is most beneficial, but rather a given fraction of both, depending on diffusion rate. The differences between proton and lithium ion movement through pyrolusite, ramsdellite or fractions of both could be indicative of: (i) a different activation energy processes for the diffusion of protons and lithium ions through these structures, and/or (ii) a different hop mechanism for lithium ions in these structures than that outlined for the proton by Balachandran et al. [32].

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We can also compare our findings with an investigation of the relative movement of lithium ions and protons into EMD in the aqueous system [33]. In this study EMD was heated at temperatures ranging from 25 ◦ C (unheated) to 300 ◦ C. Although these authors determined that the ratio of Li+ /H+ diffusion kinetics was still less than 1 even for the material heated at 300 ◦ C (i.e., H+ moving quicker through this structure), the ratio was seen to increase with increased heat treatment temperature. While differences can be expected when considering ion movement in the non-aqueous system, this finding also supports the notion of lithium diffusing faster through more pyrolytic structures. The limit of temperatures, and hence pyrolusite fractions, considered in this work likely explains why an optimum value for lithium ion diffusion through HEMD was not found, as has been discovered in this work. Comparing the differences between protons in the aqueous system and lithium ions in non-aqueous for EMD and HEMD respectively, highlights another feature evident from this study. On average across the discharge rates tested, the utilization of the MnO2 was 64%, with a maximum of 86% achieved. However, in the aqueous system almost full utilization of EMD is achieved using discharge rates 20 mA g−1 [27]. Considering the slow rates HEMD has been tested in this study, the comparatively low utilization is surprising. It should be noted, however, that some of the materials produced from this study are in fact superior in terms of specific energy and comparable in terms of specific power to those previously reported in the literature [34]. The lower utilization across the range of materials examined reveals what may be a fundamental limit of the intercalation of lithium into HEMD. 3.4.2. Unit cell volume Fig. 3 shows that the unit cell has only a minor influence on the performance of HEMD. Within the range of structures considered (0.1134–0.1181 nm3 ) most of the unit cell volumes act to the detriment of the material, albeit only slightly. However, an equally slight positive contribution is made when the unit cell volume is below 0.1146 nm3 for 2 mA g−1 discharge. The change from negative to positive contribution shifts to smaller volumes for higher discharge rates. At these high rates however, some benefit is gleaned when the unit cell is larger. From this we can conclude that with increasing discharge rate a smaller cell volume becomes increasingly limiting and a more expanded cell volume is desirable. Considering the material unit cell volume and Pr content together then, we see that at low discharge rate a moderate pyrolusite content (Pr = 0.65) and low unit cell volume (0.1134 nm3 ) are the values which contribute most positively to electrochemical performance. However, in the transition from ramsdellite to pyrolusite, as occurs as Pr is increased, the unit cell volume decreases with it, changing most significantly in the b crystallographic direction [29]. To reap a benefit from a smaller unit cell therefore creates a less desirable crystal structure, or vice versa. Although some compromise is likely possible (i.e., materials could be found with a smaller than typical unit cell for a certain Pr ), to find a material with the optimum Pr and optimum unit cell is unlikely. 3.4.3. Average crystallite size The effect of the average crystallite size on the capacity of HEMD is shown in Fig. 4. Compared to the other parameters considered, the effect of the average crystallite size is minimal. This is perhaps somewhat surprising given that larger crystallites would lead to longer diffusion paths in the solid state and consequently potentially result in whole regions of the crystallite being inaccessible. However, its relatively low effect suggests that the other parameters are more limiting than the range of average crystallite sizes (11.4–23.2 nm) considered here. Further, it could be that taking the average of the crystallite size is not the best measure of the influence of crystallite size on the electrochemical performance. It

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may well be that one particular direction is particularly influential on the performance, while the others play little or no role. At this stage however, there is no indication in the literature as to what crystallographic directions would be best to include in such an analysis.

3.5. Influence of composition parameters 3.5.1. Electrochemically active Mn(IV) Fig. 5 shows the effect that the percentage Mn(IV) has had on the electrochemical performance. As can be expected, higher proportions of electrochemically active Mn(IV) contribute more to higher capacity. Materials with below 60% Mn(IV) however, show a negative contribution to performance. This feature arises because of the inversely proportional relationship between Mn(IV) and Mn(III). A low Mn(IV) suggests high Mn(III), a feature which would act to reduce the specific capacity of the material. It is interesting to note the point at which the proportion of Mn(IV) begins to enhance material performance is approximately 60%, and is essentially independent of discharge rate.

3.5.2. Mn(III) defects We anticipate that the presence of Mn(III) in the material would be detrimental to the capacity primarily because its presence reduces the percentage of the active Mn(IV) species per gram of material. Additionally, the structural water associated with the Mn(III) defect acts to limit the capacity. The way in which Mn(III) percentage has influenced the capacity is displayed in Fig. 6. In general, the trend shown here is as we would expect, that is, lower amounts of Mn(III) are less detrimental to the performance of the material, to an essentially nil contribution when no Mn(III) remains. What is unforeseen however, is that materials with large proportions of Mn(III) (∼9%) are less detrimental to performance than those with lesser amounts. The reason for this relationship is at this stage unclear.

3.5.3. Cation vacancy defects The plot showing the effect of cation vacancies on the performance (Fig. 7) indicates that a high fraction of vacancies is unfavorable for performance regardless of discharge rate. These vacancies lower the ratio of Mn/O, which leads to less active material per gram and therefore lowers the material capacity. In addition, the structural water associated with these vacancies would also act to the detriment of the material. When HEMD is immersed in the non-aqueous environment these protons can diffuse out of the structure into the cell environment where they either react with the electrolyte to form HF, or with the metallic lithium anode forming H2 , both of which are destructive to the electrochemical performance of the cell. The materials with high CVF are typically low temperature HEMD’s, thus highlighting the importance of water removal from MnO2 cathode materials, a task best performed at higher temperatures. It is evident however, that a vacancy fraction below ∼0.05 begins to contribute in a positive way to the overall capacity. The presence of these vacancies opens extra diffusion pathways between the tunnels of the ramsdellite/pyrolusite structure increasing the dispersion of charge through the material, which would enhance the lithium kinetics and penetration. The vacancy fraction leading to optimum performance for most discharge rates is 0.02, barring that for 20 mA g−1 rate, which is lower, at 0.01. Fewer vacancies lower the dispersion of charge and therefore may cut off access to regions of the material. The lower optimum value for the 20 mA g−1 rate could be the result of: (i) faster movement of lithium ions directly along the tunnels rather than through vacancies connecting

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tunnels, and/or (ii) the remaining water content is more influential in limiting the capacity of the material. 3.6. Influence of morphology parameters 3.6.1. Surface area Fig. 8 demonstrates that materials in the range of BET surface areas examined have a detrimental effect on material capacity at the low discharge rates, but this becomes more beneficial at higher rates. The results indicate that at low discharge rates, both low and high surface area materials are less damaging to the capacity, therefore making them more desirable. The opposite is true at the higher rates, where these materials contribute less to good performance. Before discussing these relationships in detail, we note some discrepancies between the findings of Ilchev and Banov [20] (who report a direct relationship between high specific surface area and high specific capacity) and this work. A number of reasons may bring about these differences. Firstly, their studies used chemically prepared -MnO2 derived from the thermal oxidation of a dense MnCO3 , while our materials were deposited electrochemically. While details of the synthesis conditions are not supplied in their work, the materials prepared generally had much higher specific surface area (80–105 m2 g−1 ) than typical EMD (20–80 m2 g−1 ). Further, the scope of materials reported in the study by these authors (6 samples each heated at 300 ◦ C for 2 h) is narrower than that considered here (23 samples heated in a 200–400 ◦ C range for between 0.5 and 26 h). With a material as complex as manganese dioxide the limited sample range considered by these authors, along with the focus on a single material property, may account for the differences between the their findings and those reported here. To understand the effect of surface area on electrochemical performance as found in this work, we must first consider the reduction mechanism for MnO2 . Regardless of discharge rate, lithium ions must travel through the electrolyte medium to the surface of the material. Once ions arrive at the surface they can be intercalated into the structure to compensate for the charge as Mn(IV) is reduced to Mn(III). The lithium ions can then diffuse through the solid state to achieve a homogenous material, which also allows further lithium intercalation at the surface regions. To explain how lower rates have led to a negative contribution to performance, we draw upon some conclusions from a previous study [35] investigating the reduction mechanism of HEMD (350 ◦ C for 1.52 h) in Li/MnO2 batteries. There was some indication in this work that a phase change from /ˇ-MnO2 to LiMn2 O4 , and subsequently Li2 Mn2 O4 , occurred at the surface of the material with increased lithiation. While this surface layer is yet to be characterized in greater detail, at low discharge rates the lithium ion penetration into the core the particle is slow. Thus, only a small fraction of the particle may be accessed before the lithium residing in the surface regions cause a phase change and surface layer formation. If the diffusion of lithium ions is slower through this layer and/or it is less conductive it would render the remaining core of the particle less accessible to lithiation and hence limit full utilization. A less negative contribution for materials with high surface area can be attributed to the higher volume of micropores (those <2 nm wide) which have been found to increase significantly for materials with higher porosity. At low rates, lithium movement along these pores should be possible, hence allowing more of the active material to be utilized. Conversely, at low surface areas a lower proportion of the surface layer is formed which is therefore less deleterious to material performance. For high rate discharge, surface area is seen to contribute in a positive way to performance. Discharge at these rates would be affected later by the surface layer formation since the lithium ions are inserted faster and consequently more of the material may be

accessed before the surface layer becomes limiting. However, a low surface area is of less benefit to performance (as indicated in Fig. 8) since the exposure of the MnO2 surface to lithium ion intercalation is limited. Consequently, to achieve full utilization in these materials, the inserted lithium ions must diffuse a longer path through the solid state to access the core of particles. As lithium ions are driven into the structure more rapidly the long diffusion path length and slow solid state kinetics become severely limiting leading to much poorer material utilization and therefore capacity. A highly porous material also acts to limit the capacity at high rates. This can be explained in terms of the higher volume of micropores found in these materials. As discussed, the higher number and/or longer network of micropores which lead to high porosity may be kinetically limiting to lithium ion diffusion, or simply too small to allow diffusion along them. While the movement of lithium along these pores may be adequately fast for low rates, at high discharge rates these kinetic limitations become more detrimental to electrochemical performance. In a previous study [26] we found that water removal from highly porous materials is retarded, possibly due to slower H+ diffusion through the more extensive pore network. It follows that the diffusion of the larger lithium ion experiences similar limitations, but to a greater degree, in the discharge process of these batteries. Together, the features of low and high surface area materials result in an optimum surface area for 20 mA g−1 discharge rate, found to be 44 m2 g−1 .

4. Conclusions In this work, we have used a statistical model to relate the physical properties of heat treated manganese dioxides and the primary capacity in Li/MnO2 batteries. The relationships formed in this work provide vital insight into the intelligent selection of a HEMD with the ideal properties to give superior electrochemical performance. Specific outcomes from this work include: 1. The optimal HEMD structure had a pyrolusite content of 0.65–0.75, with the higher proportions of pyrolusite more favorable with increasing discharge rates. Small unit cell volumes are also positive, while the average crystallite size was found to have a minimal impact. 2. The composition factors all contributed significantly to the primary capacity. The best values for these parameters were found to be a high Mn(IV) percentage, low amounts of Mn(III), and a low (but non-zero) cation vacancy fraction. 3. Discharge rate was found to dramatically impact what morphology was preferred. At 2 and 5 mA g−1 surface areas of <25 and >60 m2 g−1 were favored, while at the 10 and 20 mA g−1 rates the optimum values were 44 and 54 m2 g−1 , respectively. 4. Physical properties having the most influence over the primary capacity were determined to be pyrolusite content, Mn(IV) percentage, cation vacancy fraction and surface area. Therefore, these should be the primary focus in selecting a HEMD to achieve superior performance in Li/MnO2 batteries. 5. Some of the HEMD materials reported here show superior performance in terms of specific capacity compared to those previously reported in the literature [34], with further improvements anticipated through the intelligent selection of HEMD materials with the ideal properties outlined in this study.

Acknowledgments WMD acknowledges the UoN for the provision of an APA PhD scholarship. Thanks are also extended to the UoN EM-X-ray unit for assistance in obtaining the XRD data.

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