A Study of MnO2 with Different Crystalline Forms for Pseudocapacitive

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A Study of MnO2 with Different Crystalline Forms for Pseudocapacitive Desalination Zhi Yi Leong, and Hui Ying Yang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b20880 • Publication Date (Web): 14 Feb 2019 Downloaded from http://pubs.acs.org on February 16, 2019

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ACS Applied Materials & Interfaces

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A Study of MnO2 with Different Crystalline Forms

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for Pseudocapacitive Desalination

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Zhi Yi Leong and Hui Ying Yang*

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*Pillar of Engineering Product Development (EPD), Singapore University of Technology and

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Design, 8 Somapah Road, 487372, Singapore. E-mail: [email protected]

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KEYWORDS: MnO2, hollandite, birnessite, spinel, capacitive desalination, desalination

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ABSTRACT: Recent research in materials for capacitive deionization (CDI) has shown that

8

intercalation materials are capable of salt removal capacities (>40 mg g-1) much higher than that

9

of carbon materials (~15 mg g-1). However, little work has been done to elucidate the relationship

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between the microstructure of an intercalation material and its desalination performance. Herein,

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we report the desalination performances of various crystalline forms of MnO2 in a hybrid CDI

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(HCDI) setup without the use of ion-exchange membranes. MnO2 materials used in our

13

experiments were either poorly crystalline or crystalline forms of 1D hollandite α-MnO2, 2D

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birnessite δ-MnO2 and 3D spinel λ-MnO2. X-ray photoelectron spectroscopy (XPS) performed on

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electrochemically cycled MnO2 showed redox changes associated with intercalation processes in

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crystalline MnO2 whereas poorly crystalline MnO2 showed no such changes. It was further shown

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that surface adsorption is dominant in poorly crystalline MnO2 and that poorly crystalline forms

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of α- and δ-MnO2 exhibited the highest salt removal capacities of 0.17 and 0.16 mmol g-1 (9.93

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and 9.35 mg g-1) respectively. These performances are comparable to state-of-the-art carbon

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materials and are remarkable considering the low surface areas (< 400 m2 g-1) of MnO2 materials.

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INTRODUCTION

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Capacitive deionization (CDI) is an ion-removal technique first developed in the 1960s

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economic desalination of saline waters. A typical CDI cell operates by passing saline water through

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a channel created between two oppositely charged porous electrodes and ions in water (typically

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Na+ and Cl-) are removed via the formation of electrical double layers (EDLs) 2. Once ion

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saturation is reached, the electrodes are discharged and ions are released into a waste stream. The

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discharging of electrodes causes the release of stored electrical energy and this can be harvested

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to offset the initial investment in energy. Furthermore, infrastructure required for CDI is minimal

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when compared to conventional desalination technologies such as reverse osmosis (RO)

1

3

for

or

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thermal distillation 4; both of which require either costly membranes or thermal generators. Despite

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its supposed advantages, CDI has not been widely adopted for large-scale desalination due to its

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relatively low salt adsorption capacity and energy inefficiency at high salt concentrations. These

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limitations can be traced back to the traditional method of operating CDI, which is to prepare

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highly porous carbon electrodes for adsorption of ions. There are two main problems with this

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method. The first is that ion adsorption does not always scale with surface area due to the fact that

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not all surface area is accessible to ions 5. The second is that co-ion expulsion inhibits optimal ion

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adsorption and reduces overall charge efficiency 2, 6.

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Researchers attempt to overcome these limitations by using materials which can 7-9.

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preferentially remove ions based on redox reactions

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motivation that higher surface areas tend to lead to higher salt removal capacities and represents a

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transition from materials dependent on EDL adsorption to bulk intercalating or conversion

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materials 8-11. One of the earliest works validating this concept was done by Lee et al. using tunnel

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structured sodium manganese oxide (NMO)

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This strategy deviates from the original

in a hybrid CDI (HCDI) cell. NMO was used in 3

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conjunction with an activated carbon (AC) paired with an anion-exchange membrane (AEM) to

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yield a salt adsorption capacity of 31.2 mg g-1 and was considered the highest at that time.

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Following the success of this pioneering work were studies on materials with Na+ ion intercalation

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properties. Manganese dioxides 13-16, sodium iron pyrophosphate 17, titanium disulphide 9, tunnel

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structured manganese oxides 18, Prussian blue analogues 11, 19-20 were all investigated for their Na+

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removal properties. Prussian blue analogues in particular received significant attention and one

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innovative application of sodium nickel hexacyanoferrate (NaNiHCF) and sodium iron

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hexacyanoferrate (NaFeHCF) yielded an impressive adsorption capacity of 59.9 mg g-1 11. This

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result was more than double that of the highest salt adsorption capacities produced by graphene 21-23.

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composites

While the improvement in desalination performance is undeniable, the

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experimental formats of these hybrid cells require one AEM to manage anion transport. This

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inherent requirement makes it difficult to characterize an intercalation material’s desalination

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performance based on its own merit and it is unfair to compare adsorption capacities with materials

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employing the traditional CDI format.

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While most of the recent work concentrated on NMO or Prussian blue analogues, there was

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also considerable interest in MnO2 materials. MnO2 is a pseudocapacitive material which stores

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ions based on a single-electron transfer redox reaction involving the Mn4+/ Mn3+ couple 24-25. This

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reaction can be conducted through two possible mechanisms, bulk intercalation of cations (𝐶 + ),

19 20 21

𝑀𝑛𝑂2 + 𝐶 + + 𝑒 ― ⇌ 𝑀𝑛𝑂𝑂𝐶

(1)

and/ or surface adsorption of cations 25,

(𝑀𝑛𝑂2)𝑠𝑢𝑟𝑓𝑎𝑐𝑒 + 𝐶 + + 𝑒 ― ⇌ (𝑀𝑛𝑂2― 𝐶 + )𝑠𝑢𝑟𝑓𝑎𝑐𝑒


(2)

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. Thin films of MnO2 have been shown to possess an extremely high specific capacitance of 1380

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F g-1 25 whereas bulk electrodes of MnO2 can reach more than 290 F g-1 26-28 depending on their

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crystal structure. The huge difference in specific capacitance is not only related to how MnO2

4

electrodes are fabricated but is also the result of a number of factors including electrical contact

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between material/ current collector interfaces, mean manganese oxidation states, crystal defects,

6

morphological effects and electrolyte permeability etc

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exists in a variety of polymorphs which assume different tunnel structures and sizes. Even within

8

the same phase of MnO2, there can be more than one type of tunnel structure

9

suitable MnO2 materials for CDI requires a multi-faceted approach.

24.

To further complicate things, MnO2

29.

Thus, finding

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In this work, we investigate the electrochemical and desalination properties of MnO2 across

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two areas, (1) crystal structure and (2) crystallinity. These two areas are notable for their direct

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impacts on ion storage. Crystal structure in MnO2 can be classified as 1D, 2D or 3D based on the

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orientation of tunnels formed by interlinking chains of MnO6 octahedra units. Different tunnel

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structures and arrangement result in different tunnel sizes and crystallographic volumes for ion

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storage. To investigate these effects, we synthesized three types of MnO2 polymorphs, 1D

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hollandite α-MnO2, 2D birnessite δ-MnO2 and 3D spinel λ-MnO2. Regarding the crystallinity of

17

MnO2, a previous study

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crystalline MnO2 compounds and surface adsorption in amorphous MnO2. We investigate this by

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synthesizing poorly crystalline and crystalline forms of α- and δ-MnO2 and evaluate the efficacy

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of each in electrochemical and desalination experiments. Desalination experiments were

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performed using a HCDI cell without the use of ion-exchange membranes to obtain base values of

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salt removal capacity. An AC electrode served as the anode and was carefully fabricated with a

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mass percentage determined by the ratio of gravimetric capacitances between MnO2 and AC.

30

had shown that intercalation of cations most likely occurred in



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Through the experimental evidence provided in this work, we hope to elucidate the relationships

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between crystal microstructure, electrochemical behaviour and desalination performance.

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RESULTS AND DISCUSSION

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Figure 1. Crystal structures of (a) α-MnO2, (b) δ-MnO2 and (c) λ-MnO2. Green spheres: Mn.

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Red spheres: O. Pink spheres: interstitial cations (K+, Na+ or Li+). Blue spheres: interstitial water.

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Table 1. Tunnel size and features of MnO2 crystals 24, 27, 31-32.

Crystallographic form

Mineral

α-MnO2

hollandite

δ-MnO2

birnessite

Features (1 x 1), (2 x2) tunnels (1 x ∞) layer

Size (Å) 2.3 x 2.3, 4.6 x 4.6 7

8

λ-MnO2 spinel (1 x 1) tunnel 2.3 x 2.3 Structural and chemical characterization. The crystallographic structure of each type of MnO2

9

is shown schematically in Figure 1 and a summary of their approximate tunnel sizes are given in

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Table 1. Each MnO2 polymorph is derived from a series of edge, corner or face sharing MnO6

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octahedra units which in turn form an infinite number of interlinked octahedra chains

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Tunnels constructed from MnO6 octahedra units are classified in a m x n format, where m and n

13

represent the number of MnO6 octahedra units used for tunnel height and width respectively 29, 33.


24, 27, 32.

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The formation of these tunnels during synthesis usually results in the entrapment of ions and water

2

thus, inductively coupled plasma optical emission spectrometry (ICP-OES) and thermogravimetric

3

analysis (TGA) were used to determine the ionic composition and amount of crystal water.

4

Empirical formulae determined from ICP-OES are presented in Table S1 and TGA curves are

5

shown in Figure S7. Briefly, the results show small amounts of cationic species in the as-

6

synthesized materials and crystal water loss between 0.65 – 2.34 % above 190 °C. A detailed

7

discussion is provided in Supplementary Note S1.

8 9 10

Figure 2. XRD patterns of (a) α-MnO2, (b) δ-MnO2 and (c) λ-MnO2. (Red lines: crystalline phase. Black lines: poorly crystalline phase. Blue indexes: impurity peaks.) 7 ACS Paragon Plus Environment


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Powder XRD patterns of all MnO2 samples are shown in Figure 2. XRD patterns for both

2

crystalline and poorly crystalline samples of α-MnO2 are given in Figure 2a and an immediate

3

contrast in peak resolution can be observed. Peaks for α-MnO2-C were well-defined and almost all

4

of them could be indexed to pure phase tetragonal α-MnO2 (PDF#00-044-0141). On the contrary,

5

only two broad peaks centred at 2θ = 37.1 and 66.3 ° were observed for the poorly crystalline

6

sample. Figure 2b shows the XRD pattern for birnessite δ-MnO2-C sample and some impurity

7

peaks of cryptomelane (2θ = 17.9, 28.5, 37.5 and 53.2 °) were present. Cryptomelane impurities

8

could had formed via two possible ways, the first was during calcination of birnessite powders 24

9

and the second was during H2SO4 washing

26.

Nevertheless, the rest of the peaks were clearly

10

indexed to the monoclinic system (PDF#00-043-1456) of pure phase δ-MnO2. The low angle peak

11

of 12.3 ° was particularly strong due to the basal spacing of the layered structure

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MnO2-C, δ-MnO2-PC only had two strong broad peaks at 2θ = 12.3 and 36.9 ° and two smaller

13

peaks at 2θ = 24.7 and 65.7 °. Spinel λ-MnO2-C obtained after the delithiation process showed

14

sharp, well-resolved peaks that could be indexed to a cubic symmetry with space group Fd3m (no.

15

227) (PDF#00-042-1169). We note that evacuation of lithium is incomplete as some peaks of

16

LiMn2O4 (2θ = 32.9, 38.7 and 55.1 °) were present in the final sample.

17 18 19


33.

Unlike δ-

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Figure 3. SEM images of poorly crystalline (a, b) α-MnO2 and (c, d) δ-MnO2.



1 2

Figure 4. SEM images of crystalline (a, b) α-MnO2, (c, d) δ-MnO2 and (e, f) λ-MnO2.


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1 2

Figure 5. HRTEM images of (a) α-MnO2-PC, (b) α-MnO2-C, (c) δ-MnO2-PC, (d) δ-MnO2-C and

3

(e) λ-MnO2-C. Inserts are the corresponding SAED patterns.



1

Both low and high magnification photomicrographs of MnO2 are shown in Figure 3 and 4.

2

α-MnO2-PC was shown to be composed of spherical particles ~200 to 400 nm in diameter with a

3

rough surface characterized by nanosized hair-like fibres (Figure 3a and b). No excessive

4

aggregates were observed and clear interparticle boundaries were present. EDS images and

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mapping spectrum (Figure S2) showed typical Mn and O elements but also revealed a small

6

percentage of K. This is likely due to K+ ions introduced during the formation of tunnel structures

7

in α-MnO2-PC

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different (Figure 4a and b). α-MnO2-C exhibited a rod-like morphology of ~500 to 1000 nm in

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length and ~50 to 70 nm in diameter. Individual rods were well-defined and had smooth surfaces.

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Some of these adjacent rods had even fused together to form thick bundles. Elemental mapping of

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α-MnO2-C (Figure S3) was similar α-MnO2-PC. SEM micrographs of δ-MnO2-PC showed a

12

collection of ill-defined particulate aggregates (Figure 3c and d). The surfaces of these aggregates

13

were uneven and comprised of multiple small particles. EDS mapping showed that these particles

14

are composed of Mn, O and some Na. In contrast, δ-MnO2-C consisted of nanosheets ~20 nm in

15

size with fibrous structures extruding from the edges (Figure 4d). A higher magnification (Figure

16

4c) showed several of these platelets stacked together in aggregates. Elemental mapping for δ-

17

MnO2-C shows Mn, O and K. Lastly, λ-MnO2-C shown in Figure 4e and f consisted of spherical

18

particles ~30 to 70 nm in diameter which appeared smooth relative to α-MnO2-PC.

34-35.

On the other hand, hydrothermally synthesized α-MnO2-C were strikingly

19

High-resolution transmission electron microscopy (HRTEM) images of MnO2 samples are

20

given in Figure 5. Consistent with XRD results, poorly crystalline forms of α-MnO2-C and δ-

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MnO2-PC showed a generally disordered microstructure with some degree of localised

22

crystallinity while long-range crystal order can be observed for their crystalline counterparts.

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Figure 6. (a) N2 adsorption-desorption isotherms of MnO2 polymorphs (insert: close-up of

3

isotherms). (b) Cumulative pore volume and (c) pore size distribution curves of MnO2 polymorphs.

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Table 2. Summary of specific surface areas and porosity characteristics of MnO2 polymorphs.

Sample

SSABET (m2 g-1)

External surface area (m2 g-1)

Micropore surface area (m2 g-1)

Micropore area percentage (%)

Total pore volume (cm3 g-1)

α-MnO2-PC

394

283

111

28.2

0.47

α-MnO2-C

28

25

3

10.7

0.67

δ-MnO2-PC

226

180

46

20.3

1.14

δ-MnO2-C

112

112

0

0.00

0.9

λ-MnO2-C

40

40

0

0.00

0.14



1

Surface area and porosity studies using BET. Nitrogen adsorption-desorption isotherms and

2

porosity measurements are presented in Figure 6. A cursory view of the isotherms revealed

3

remarkable similarities between poorly crystalline and crystalline samples of α-MnO2 and δ-MnO2

4

(Figure 6a insert). For poorly crystalline samples of α-MnO2 and δ-MnO2, we observed a

5

combination of Type II and Type IV isotherms accompanied by a H4 hysteresis loop. These results

6

strongly suggested the existence of macroporous and mesoporous structures with slit-like pores 28.

7

The existence of macropores was further supported by a rapid rise in isotherm near the vicinity of

8

P/ P0 ≈ 1. This was especially evident in α-MnO2-PC where a meteoric rise in volume adsorbed

9

was observed at P/ P0 = 0.976. Pore size distributions of these materials also showed similar

10

patterns with peaks in microporous (1.41 and 1.76 nm) and mesoporous regions (2.76 and 3.62

11

nm). Although micropores and mesopores were both present, micropores seemed to contribute a

12

smaller fraction of overall SSA. Only 28.2 % and 20.3 % of total SSA were associated to

13

micropores for α-MnO2-PC and δ-MnO2-PC respectively. Further information regarding the

14

porosity could be inferred from the cumulative pore volume curves in Figure 6b by comparing the

15

increase in volume adsorbed before and after 4 nm. The increase in volume adsorbed before 4 nm

16

was much larger than that of after 4 nm for α-MnO2-PC which was not the case for δ-MnO2-PC.

17

Combined with our previous knowledge of micropore contribution, we deduced that mesopores

18

below 4 nm were primary contributors to SSA for α-MnO2-PC whereas mesopores above 4 nm

19

were primary contributors to SSA for δ-MnO2-PC. This difference was likely due to interparticle

20

spacings present in α-MnO2-PC (Figure 3a and b) which were not found in δ-MnO2-PC. Both α-

21

MnO2-PC and δ-MnO2-PC possessed relatively high surface areas of 394 m2 g-1 and 226 m2 g-1.

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Similarities observed in poorly crystalline samples of α- and δ-MnO2 did not extend to their

23

crystalline counterparts. α-MnO2-C exhibited a Type IV isotherm with a small H3 hysteresis loop 14 ACS Paragon Plus Environment

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terminating at P/ P0 ≈ 0.5. This isotherm was characteristic of α-MnO2 nanorods

and a small

2

SSA of 28 m2 g-1 was obtained. Pore size distribution was unremarkable with no distinct peaks

3

whereas cumulative pore volume curve showed a linear increase in volume adsorbed. The isotherm

4

of δ-MnO2-C appeared to be a combination of Type II and Type IV with corresponding H3

5

hysteresis loop. Note that the H3 hysteresis loop found in δ-MnO2 was indicative of slit-like

6

mesopores formed by interconnected nanosheets 28, 36. Pore size distribution of δ-MnO2-C showed

7

a broad peak centred at 7.45 nm and judging by the rise in volume adsorbed around this pore width

8

(Figure 6b), we concluded that mesopores in this region were responsible for much of the surface

9

area. This was further supported by t-plot calculations which showed no micropore contribution.

10

Lastly, λ-MnO2-C showed a Type II and Type IV isotherm with H4 hysteresis loop. Pore

11

size distribution had a broad peak centred at 4 nm with a peak width of about 2 nm. Cumulative

12

pore volume curve was characterized by a slight increase in volume adsorbed followed by an

13

almost horizontal plateau. A small SSA of 40 m2 g-1 was obtained. It is notable that samples of δ-

14

MnO2 possessed the highest total pore volumes of 0.9 and 1.14 cm3 g-1 (Table 2). This was possibly

15

due to wide interlayer spacings (~7 Å) present within the structure 24, 27-28.



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Figure 7. (a) Cyclic voltammetry curves at 5 mV s-1 in a 1 M NaCl solution. (b) Galvanostatic

3

charge-discharge curves at 0.1 A g-1 in a 1 M NaCl solution. (c) Specific capacitance under

4

different charge-discharge currents.

5

Table 3. Comparison of specific capacitances and equivalent series resistances (ESRs) at a charge/

6

discharge current of 0.1 A g-1 in a 1 M NaCl solution.

Sample

Specific capacitance (F g-1)

ESR (Ω)

α-MnO2-PC α-MnO2-C δ-MnO2-PC δ-MnO2-C λ-MnO2-C

182.7 100.9 279.0 120.1 66.7

13.3 30.4 11.8 20.2 246.9


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1

Electrochemical studies. CV curves shown in Figure 7a are 5th cycle waves obtained in a 1 M

2

NaCl solution at 5 mV s-1. The curves were generally rectangular without any pronounced redox

3

peaks within an electrochemical window of 0 to 1 V. Among the samples, δ-MnO2-PC possessed

4

the highest specific current and largest area under the curve, thus the highest specific capacitance.

5

The somewhat “leaf-like” voltammogram indicated a less than ideal capacitive behaviour

6

especially when compared to other samples. The second highest specific capacitance belonged to

7

α-MnO2-PC which showed a voltammogram with a prominent current tail end near 1 V. This

8

increase in specific current was likely due to parasitic reactions resulting from a low overpotential

9

for oxygen evolution reaction (OER)

27.

Smaller voltammograms were observed for crystalline

10

forms of MnO2 with α-MnO2-C and δ-MnO2-C almost overlapping each other and λ-MnO2-C

11

enclosed within. The voltammogram of λ-MnO2-C was almost a parallelogram due to a large

12

equivalent series resistance (ESR).

13

Specific capacitances were determined from GCD curves in accordance to standard 37.

14

reporting for supercapacitors

Charge-discharge curves shown in Figure 7b described an

15

essentially linear curve with a minor curvature near 0 V. Slight deviations in linearity indicated

16

pseudocapacitive behaviour and could had resulted from either one or both of the charge storage

17

mechanisms outlined in Equations (1) and (2). Table 3 summarizes the specific capacitances and

18

ESRs obtained from the discharge curve. Capacitances follow the order of λ-MnO2-C < α-MnO2-

19

C < δ-MnO2-C < α-MnO2-PC < δ-MnO2-PC. Poorly crystalline MnO2 outrank all crystalline types

20

of MnO2 and based on their surface areas, it is reasonable to assume that surface adsorption of ions

21

(Equation 2) was responsible for the high capacitance. If that was the case then a higher surface

22

area would had resulted in a greater number of ion adsorption sites and a higher capacitance.

23

However, it is curious to note that the BET area of α-MnO2-PC exceeded δ-MnO2-PC by over 17 ACS Paragon Plus Environment


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70 % yet it possessed a lower capacitance. A more comprehensive explanation is that both surface

2

adsorption and bulk intercalation processes work in conjunction during charge storage. Thus, the

3

higher capacitance of δ-MnO2-PC could be attributed to large interlayer distances between layers

4

of edge sharing MnO6 octahedra combined with favourable textural features.

5

Despite the low capacitances of crystalline MnO2, they are still relatively too high to be

6

characteristic of EDL charging. Thus, we attribute their capacitances to intercalation reactions

7

(Equation 1). Specific capacitances of crystalline MnO2 showed some dependence on the

8

dimensionality of their crystal tunnels with 2D birnessite δ-MnO2-C leading at 120.1 F g-1 followed

9

by 1D hollandite α-MnO2-C and 3D spinel λ-MnO2-C. According to recent studies

29, 38-39,

2D

10

tunnels of δ-MnO2 are able to provide more open space for ion diffusion which could lead to higher

11

ion storage and better electrochemical performance. Figure 7c also shows a low capacitance for

12

spinel λ-MnO2 which could associated to partial occupancy of tunnels by Li+ ions.

13

ESRs for our MnO2 samples exhibited a trend that negatively correlates to specific

14

capacitance and better electrochemical performances were observed for samples with lower ESR.

15

An extremely large ESR of almost 250 Ω was recorded for λ-MnO2-C which was responsible for

16

the heavy distortion in CV curve and the large voltage drop in the discharge curve. Subsequently,

17

we expected λ-MnO2-C to exhibit the worst desalination performance among the samples.

18

XPS analysis of MnO2 electrodes. In order to investigate the charge storage mechanisms further,

19

we performed XPS surface analysis on MnO2 electrodes under three experimental conditions, as-

20

prepared, oxidized at 1 V and reduced at 0 V. Oxidized or reduced electrodes were cycled 10 times

21

before ending at 1 or 0 V respectively for their analyses.


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1

Mean manganese oxidation states were determined from Mn3s core level spectra by

2

examining the energy separation between deconvoluted Mn3s peaks. The difference in energy

3

levels, also known as multiplet splitting, is caused by parallel spin coupling between the 3s electron

4

and 3d electron during photoelectron ejection

5

oxidation state based on a linear relationship established between the energy difference and

6

valence of Mn

7

given in Table S2. The average oxidation state of as-prepared samples for poorly crystalline α-

8

MnO2 and δ-MnO2 is 3.6 and changes slightly depending on the experimental condition. α-MnO2-

9

PC exhibited a lower oxidation state of 3.5 when reduced and δ-MnO2-PC showed a higher

10

oxidation state of 3.8 when oxidized. These observations were attributed to the fast kinetics of

11

surface redox reactions where only MnO2 closest to the surface participate 25. Any changes to the

12

oxidation state were quickly normalized by redox reactions driven by the chemical potential

13

between bulk and surface MnO2 25. On the other hand, crystalline MnO2 compounds possessed an

14

average oxidation state of 3.9 with very different electrochemical characteristics. Manganese

15

oxidation states in both α-MnO2-C and δ-MnO2-C were readily reduced to near +3 and returned to

16

near +4 states upon oxidation. A more apparent change in manganese oxidation state was observed

17

due to stabilization of Na+ ions within the crystal structure. The only exception was λ-MnO2-C

18

which showed minor changes in its oxidation state due to partial stabilization by residual Li+ ions.

19

O1s core level spectra presented in Figure S14 was deconvoluted into three main peaks

20

belonging to Mn-O-Mn (529.8 ± 0.1 eV), Mn-O-H (530.9 ± 0.5 eV) and H-O-H (532.3 ± 0.6 eV).

21

We note that the percentage of structural water present in the as-prepared sample was mostly

22

higher than reduced or oxidized samples (Table S1). Interstitial water was gradually removed as

23

the sample was cycled to an electrochemically stable state. Percentages of Mn-O-H and H-O-H

41-42.

25, 40.

We can calculate the mean manganese

Mn3s core level spectra is shown in Figure S13 and the associated data is



1

remained fairly stable as the electrode was cycled between its reduced and oxidative states. Thus,

2

we concluded that hygroscopic functional groups were unaffected by the effects of electrochemical

3

cycling. Our XPS analyses further confirmed our hypotheses that surface adsorption was dominant

4

in poorly crystalline MnO2 whereas intercalation was the main process in crystalline MnO2.

5

Desalination performance. We evaluated the desalination efficiency of MnO2 cathodes across

6

three different metrics, salt removal capacity, charge efficiency and energy consumption. MnO2

7

cathodes were polarized to negative potentials for two reasons; the first is to prevent the adsorption

8

of Cl- ions on its surface and the second is to ensure that MnO2 was immediately reduced instead

9

of oxidizing it first which will would had resulted in an inverted CDI performance 15.

10 11


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1 2

Figure 8. (a) Representative cycles of MnO2 samples with corresponding current profiles in a 5

3

mM NaCl solution at -1 V. (b) Salt removal capacity against applied potential in a 5 mM NaCl

4

solution. (c) Salt removal capacity against salt concentration at -1 V potential. (d) Salt removal

5

capacity against specific capacitance. Desalination results were obtained based on 5th cycle of ion

6

removal/ release.

7

Figure 8a shows representative cycles of all MnO2 samples in a 5 mM NaCl solution under

8

an applied potential of –1 V. It is evident that MnO2 samples in the poorly crystalline phase

9

experienced more rapid decreases in solution conductivity than crystalline ones. This observation

10

was consistent with that seen in AC and carbon-based composite samples, which indicated that

11

surface adsorption of ions had occurred. Note that we do not differentiate between surface 21 ACS Paragon Plus Environment


1

adsorption due to EDL formation and surface adsorption due to electrochemically active sites

2

(Equation 1). We varied the potential to less reducing potentials (more positive values) to observe

3

changes in salt removal and the results are presented in Figure 8b. Generally, lower reducing

4

potentials resulted in lower salt removal capacities. At a low reducing potential of -0.6 V, samples

5

exhibited salt removal capacities comparable to each other but as the potential became increasingly

6

negative, poorly crystalline samples began to outperform crystalline ones. This better performance

7

is largely attributed to their superior surface area which allowed high surface adsorption. λ-MnO2-

8

C was expected to fare the worst based on its capacitance but had performed modestly well

9

between α-MnO2-C and δ-MnO2-C.

10

We further tested our samples in salt solutions of different concentrations as shown in

11

Figure 8c. Poorly crystalline MnO2 presented much better performance at higher salt

12

concentrations with α-MnO2-PC and δ-MnO2-PC reaching 0.17 ± 0.01 and 0.16 ± 0.004 mmol g-1

13

(9.93 and 9.35 mg g-1) respectively. This result is equal or higher than some carbon materials

14

(mesoporous AC – 4.6 mg g-1, chemically reduced graphene oxide – 1.85 mg g-1, single-wall CNTs

15

– 9.35 mg g-1)

16

crystalline MnO2 was more apparent at higher salt concentrations (≥ 10 mM) which implied that

17

salt removal was largely dependent on surface adsorption. When the highest salt adsorption

18

capacity of each sample was plotted against specific capacitance (Figure 8d), some curious

19

observations could be made. Firstly, the hypothesis that materials with higher specific capacitances

20

exhibited better salt removal performances holds true and we observed significantly higher salt

21

removal for poorly crystalline MnO2 materials. Despite δ-MnO2-PC having better electrochemical

22

properties, it was α-MnO2-PC that showed a slightly better desalination performance. We

23

hypothesized that an advantageous combination of high surface area and micropore percentage

21.

The difference in desalination performance between poorly crystalline and


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1

provided α-MnO2-PC with a greater advantage over δ-MnO2-PC at higher salt concentrations as

2

observed in Figure 8c. Secondly, the poor performances of crystalline MnO2 materials were

3

possibly due to low surface areas coupled with weak diffusion kinetics of ions into the crystal

4

structure. Although we expected salt removal to follow a similar trend as capacitance; i.e. λ-MnO2-

5

C < α-MnO2-C < δ-MnO2-C, this was not observed as the desalination performance of λ-MnO2-C

6

was actually higher than that of α-MnO2-C. Further experimental evidence provided by charge

7

efficiency results (Figure S16d) suggest that it was comparatively more efficient for λ-MnO2-C to

8

remove one mole of salt than it was for α-MnO2-C. The better performance of λ-MnO2-C can be

9

attributed to a weaker dependence on intercalation as suggested by XPS results. The cycling

10

performance of MnO2 was evaluated in 10 desalination cycles as shown in Figure S15. α-MnO2-

11

PC, α-MnO2-C and δ-MnO2–PC showed no decrease in ion removal capacity while slight capacity

12

fading can be observed at the end of the 5th cycle for δ-MnO2–C and λ-MnO2-C. This capacity

13

fading could be attributed to structural degradation in MnO2 as suggested in a recent study 38.

14

Charge efficiencies are reported in Figure S16 and poorly crystalline and crystalline

15

samples were analysed separately. In an ideal case, we expect charge efficiency to be unity which

16

corresponds to the removal of one mole of salt with the transfer of one mole of electron as

17

stoichiometrically suggested by Equations (1) and (2) 13. However, low charge efficiencies were

18

observed for all samples of MnO2 which was unsurprising given that bulk electrical conductivity

19

values for MnO2 are in the order of 10-7 to 10-3 S cm-1 24. Figures S16a and b describe the changes

20

in charge efficiency with respect to applied potential. No discernible pattern exists for poorly

21

crystalline MnO2 and charge efficiencies fall between the range of 0.1 to 0.4. On the other hand,

22

charge efficiencies of crystalline MnO2 decreased at high reducing potentials. This result indicates

23

an inherent trade-off between high salt removal and charge efficiency for crystalline MnO2 23 ACS Paragon Plus Environment


1

materials. The salt concentration was also varied while the potential was kept constant at -1 V and

2

a sporadic distribution of charge efficiencies lesser than 0.35 are shown in Figure S4c and d with

3

α-MnO2-PC having the lowest charge efficiency at ~0.05.

4

CONCLUSION

5

We have demonstrated for the first time, the application of MnO2 in a HCDI cell without the use

6

of AEMs. Crystalline forms of 1D hollandite α-MnO2, 2D birnessite δ-MnO2 and 3D spinel λ-

7

MnO2 were synthesized along with poorly crystalline forms of α-MnO2 and δ-MnO2. Specific

8

capacitances obtained from GCD experiments showed that charge storage was not solely due to

9

either surface adsorption or bulk intercalating reactions but a combination of both. XPS results on

10

reduced and oxidised electrodes further showed that intercalation processes dominated crystalline

11

MnO2 whereas surface adsorption was more prevalent for poorly crystalline MnO2. Desalination

12

experiments showed that ion removal was higher for porous poorly crystalline MnO2 than their

13

crystalline counterparts. α-MnO2-PC and δ-MnO2-PC both showed salt removal capacities of 0.17

14

and 0.16 mmol g-1 (9.93 and 9.35 mg g-1) which are competitive against traditional carbon

15

materials. Based on our findings, the surface area of MnO2 is a crucial factor in improving its

16

desalination performance. Thus, MnO2 should be modified in carbon-based composites such as

17

CNTs 43 to increase surface area and improve electrical conductivity.

18

SUPPORTING INFORMATION

19

Experimental details, synthesis methods, electrode fabrication, HCDI setup, XPS, EDS, ICP-OES,

20

TGA and charge efficiency results.

21

ACKNOWLEDGEMENTS


Page 24 of 28

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1

The research project is supported by the Singapore National Research Foundation under its

2

Environmental & Water Technologies Strategic Research Programme and administered by the

3

Environment & Water Industry Programme Office (EWI) of the PUB.

4

REFERENCES

5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40

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Page 28 of 28

A Study of MnO2 with Different Crystalline Forms for Pseudocapacitive

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