metal oxide or hydroxide phase is thermodynamically stable in contact with an aqueous electrolyte. If the pure metal is in contact with an electrolyte, then it corrodes with the formation of the deposit of oxide/hydroxide film, which slows down or stops the corrosion. This means that the metal is *passivated*. The composition of the passive film depends on the conditions of its formation. For the same metal, various Pourbaix diagrams can be plotted for various passive films. For the purpose of the current work, Fig.3.6 shows the most thermodynamically stable oxide phases at the ambient conditions. On the other hand, the metal oxide films formed under the electrochemical control are often amorphous, hydrated, or even non-stoichiometric [108]. The borders of the passivation domains depend on the solubility of the given phase. In Fig 3.6 these borders are plotted for 10⁻⁶ M concentration of the soluble form.

The supercapacitors with an aqueous electrolyte operate in the potential range between lines (a) and (b) on the Pourbaix diagrams. The lines (a) and (b) show the pH-dependent potential of the water electroreduction to H₂ and the electrooxidation to O₂ correspondingly: this is the potential range of water stability. As one can see from Fig.3.6, the diagrams for different metals differ significantly. Below some types of the electrochemical behavior of the selected TMO are considered in details.

### 3.3.1 TMO stable in wide pH range: $RuO_2$, $TiO_2$

Relatively few transition metals are stable in a wide range of pH: for example $TiO_2$ (Fig.3.6 A), $RuO_2$ (Fig3.6 H). In particular, $RuO_2$ is the most extensively studied TMO as supercapacitor electrode material. Several comprehensive reviews on the applications of the $RuO_2$-based electrodes in supercapacitors are available [109], including a recent work [110]. Besides its chemical and electrochemical stability (Fig.3.6(H)), $RuO_2$ has a high electrical conductivity $\sigma = 2.8 \cdot 10^4$ S.cm⁻¹ [111], resulting from the position of the Fermi level within the DOS of the 2d band [112]. Due to the high electronic conductivity, the $RuO_2$ electrodes can be easily obtained by the cathodic deposition with the fine control on the conditions and amount of deposited material [113, 114, 115].

Assuming 1e- charge storage per Ru atom, the maximal charge storage $Q_C$ for $RuO_2$, according to (3.2), is $Q_C = 721.8$ C/g which translates into the specific capacitance $C_g \approx 800$ F/g for the potential interval of 0.9V. Indeed, the values
of $C_g$ above 700 F/g are routinely obtained for $RuO_2$, indicating the high values of the intercalation parameter $\delta \approx 1$ in (1.8). This provides the ultimately high efficiency of charge storage in $RuO_2$ in acids due to the $H^+$ insertion. High values of the specific capacitance and the insertion parameters are detected even for relatively thick deposits of the $RuO_2$ oxide, up to 1.4 mg/cm$^2$ without the addition of a conductive component [110]. The $RuO_2$/carbon composites can be prepared by the cathodic electrodeposition of the oxide on the carbon substrates with desired morphology. The optimization of structure and composition of these electrode materials provides high specific capacitance values $C_g = 1176$ F/g for ca. 1 mg/cm$^2$ loading of the 12%$_{wt.}$ $RuO_2$/CNT composite [114]. The high electrochemical stability of $RuO_2$ allows applying a wide potential window up to 1V both in acid [113] and alkaline [116] electrolytes.

One may conclude that $RuO_2$ demonstrates optimal electrochemical properties for the supercapacitor applications. However, $RuO_2$ wide commercialization in the industrial supercapacitors is hurdled due to its high price and toxicity.

$TiO_2$ is a wide-gap semiconductor (the band gap $E_g=3.2$ eV for anathase) with a relatively low electronic ($5.5 \cdot 10^{-6}$ S·cm$^{-1}$ [117]) and ionic conductivities, which limits its application in supercapacitors. The annealing in a reductive atmosphere (in vacuum or $H_2$) creates the oxygen vacancies which results in higher electronic conductivity. When a sufficient amount of regular oxygen vacancies is accumulated, the Magnélli phases of $Ti_nO_{2n-1}$ are formed having the conductivity of $Ti_4O_7$, $\sigma = 1.5 \cdot 10^3$ S·cm$^{-1}$, comparable to the graphite [101] (Fig.3.3). However, according to the Pourbaix diagram (Fig.3.6A) the reduced states of $Ti^{3+}$ and $Ti^{2+}$ are unstable in aqueous solutions and are oxidized to $Ti^{4+}$, which leads to the conductivity loss. Due to the semiconducting nature of the bulk $TiO_2$, the interfacial capacitance is predominant by the charge depletion layer within the oxide. Therefore, the capacitance of $TiO_2$ is smaller than the interfacial capacitance of well-conducting materials [118]. Thus, the bulk or powdered $TiO_2$ are not considered as suitable electrode material for supercapacitors.

On the other hand, the development of 2D-nanostructured $TiO_2$ phases is an interesting and well explored approach towards the tuning of the properties of titania. The nanotubes of titania can be routinely prepared by the hydrothermal synthesis [117] or by the anodization of the Ti electrode [119, 120]. Both approaches provide the control of the geometry of nanotubes, namely, their length (up to few µm), wall thickness (from few to few tens of nm) and inner diameter
(ten or tens of nm) can be controlled. The composition of the nanotubes can also be varied by changing the composition of synthesis media or by post-treatment [121]. The advantages of the TiO₂-nanotubes in comparison to the conventional nanopowders are (i) a lower contact electrical resistance and (ii) a shorter electron and/or ion diffusion path. Due to these advantages, the TiO₂-nanotubes are widely studied as electrode materials in electrocatalysis, especially in the oxygen electroreduction and the oxygen evolution reactions due to the high corrosion resistance of TiO₂) [122]. The TiO₂-nanotubes are often used in the photoelectrocatalysis, in particular, in the water splitting reaction [121, 123]. Moreover, the impedance spectroscopy analysis of the arrays of vertically aligned TiO₂-nanotubes demonstrated a significant increase in the charge carriers density in comparison to the bulk TiO₂, resulting in much higher interfacial capacitance [124]. It was concluded that the interfacial capacitance of the arrays of nanotubes, in contrast to the bulk TiO₂, is predominant by the EDL capacitance and the pseudocapacitance related to the proton insertion [125, 126]:

$$TiO_2 + e^- + H^+ \rightleftharpoons TiOOH$$

More impressive results were obtained for the layered titanates, prepared by the hydrothermal synthesis. Here the sheets of anathase TiO₂ were separated by the interlayer space, which was stabilized by the insertion of proton or alkaline cations. The specific capacitances above 300 F/g were reported for the hydrated H₂TiO₃ [127] and the sodium layered titanates Na₂Ti₂O₄(OH)₂ [128]. This was attributed to the reversible insertion of the electrolyte cations. One may conclude that despite a low electronic and ionic conductivities of the bulk TiO₂, the development of the 2D-nanostructured layered titanate materials is a promising approach to the development of new efficient electrode materials for supercapacitors.

### 3.3.2 TMO stable in acidic pH: oxides of W and Mo

The most of transition metal oxides (for example, the oxides of Co, Ni, Mn) is more stable in neutral and weakly alkaline solutions, while being thermodynamically unstable in acids. The exceptions are the oxides of W (Fig.3.6 I) and Mo, which are stable in acids and dissolving in alkaline solutions.

In particular, WO₃ is a wide-gap (2.6 eV) semiconductor with the n-type conductivity at the ambient temperature [129]. The maximal theoretical charge stor-
age for WO₃, according to (3.2), is $Q_c = 451.3$ C/g. The monoclinic WO₃ has a relatively low intrinsic conductivity: below $10^{-5}$ S.cm⁻¹, making this phase un-practical for the supercapacitor applications. However, WO₃ can be easily reduced thermally (in a reducing atmosphere) or electrochemically, thus, forming the sub-stochiometric oxide WO₃-x due to the oxygen vacancies. The intrinsic conductivity of the reduced oxide is as high as $4.8 \cdot 10^{-2}$ S.cm⁻¹ [129]. If the synthesis of the tungsten oxide is adapted to avoid the interparticles resistance, then the WO₃-based electrodes shows a significant charge capacitance due to the H⁺ insertion phenomena occurring in acids, similar to the reaction 1.8 with RuO₂:

$$ \delta H^+ + WO_3 + \delta e^- \rightarrow H_\delta WO_3 \quad (3.10) $$

In order to avoid the interparticles resistance, various synthetic strategies are applied, namely, the template synthesis [130], the hydrothermal synthesis of nanorods [131], or the direct synthesis of graphene / oxide composites [132], [133].

In particular, the mesoporous substochiometric WO₃-x oxide shows the capacitance equal to 199 F/g [131]. This value corresponds to the insertion factor $\delta \approx 0.4$, showing a high efficiency of the charge storage. The reaction 3.10 occurs at low potentials in the wide range below ca. 0.6V RHE. It gives the advantage of using WO₃-based materials as negative electrodes in the assymetric supercapacitors. The operational voltage range of 1.65V was applied for the supercapacitor consisting of the WO₃ nanorods in the negative electrodes and the carbon nanofibers in the positive electrodes [132]. This device demonstrated an outstanding energy density equal to 35.3 Wh/kg at the power density $P = 314$ W/kg.

Besides stability in acidic electrolytes, the advantage of the molybdenum oxide MoO₃ is the accessibility of a wide range of the oxidation states from +6 to +2 [134]. However, the electrical conductivity of MoO₃ at ambient conditions is low: below $10^{-8}$ S.cm⁻¹ at 300K. Thus, a high charge capacitance was observed only in the composites with low oxide loadings [134] or the thin oxide films. In the latter case, one may use the advantage of a relatively facile exfoliation of the layered orthorhombic crystalline structure of MoO₃ [135] to obtain the specific capacitance above 500 F/g. However, these high charge capacitances are observed only at slow charging rates due to the complications with a low electronic conductivity of the oxide phase.
### 3.3.3 TMO stable in neutral and weak alkaline pH: oxides of Mn, Fe, Co, Ni

As mentioned above, the most of transition metals have stable oxides at not very strong alkaline pH, for example, Mn, Fe, Co, Ni, (Fig. 3.6 (D)-(G)).

#### Oxides of Fe

Comparing to the other transition metals mentioned above, the iron oxides have the widest pH of stability (Fig. 3.6(E)). Another remarkable advantages of the iron oxide electrode materials are their relative abundance (since iron is the most abundant element among the transition metals (Fig.3.2)) and non-toxicity. The oxides of Fe are the most interesting electrode materials from the viewpoint of sustainable development. Thus, considerable efforts have been taken for the development of the iron oxide electrode materials for supercapacitors [136]. The most common iron oxides are α-Fe₂O₃ (hematite), Fe₃O₄ (magnetite), and α- and γ-FeOOH (goetite and lepidocrocite correspondingly).

α-Fe₂O₃ (hematite) is the iron oxide with the widest range of thermodynamic stability according to the Pourbaix diagram (Fig.3.6(E)). It has the rhombohedral (corundum) crystalline structure. According to (3.2), the maximal charge capacitance for this oxide ($M = 159.69$ g/mol) is $Q_c=1209$ C/g (assuming $1e^{-}$ per 1 Fe atom storage). However, the bulk pure α-Fe₂O₃ is an insulator. The presence of a structural defect can transform it into a semiconductor [137] with a low conductivity reported in the range $10^{-4} – 10^{-14}$ S.cm⁻¹. Thus, the hematite is used as electrode material in the form of thin films [138], yet demonstrating the specific capacitance equal to 178 F/g (in 1M NaOH in the potential window 0.7V), which is much lower than the theoretically expected value. Nevertheless, more than a 6-fold increase in the surface specific capacitance was detected for the graphite foam after it was decorated by the atomic layer deposition with the uniform layer of the Fe₂O₃ nanoparticles of the size 10 nm [139]. This electrode material demonstrated a remarkable capacitance retention more than 95% after $5·10^4$ cycles in 1M NaOH. The assymetric supercapacitor with the Fe₂O₃/graphite // CoMoO₄/graphite electrodes demonstrated the specific capacitance equal to ca. 210 F/g (total electrode mass) or the energy density equal to 74.7 W.h/kg at $P = 1.4$ kW/kg.

As it has been mentioned, the presence of structural defects, in particular, oxygen vacancies, in α-Fe₂O₃ results in an increase of its electronic conductivity due
to the modification of its electronic states. The doping of α-Fe₂O₃ by the elec-	ron acceptor cations, in particular, Ti⁴⁺, leads to the formation of the stabilized
oxygen vacancies and reduced cation Fe²⁺ sites [140]. These modifications provide
a significant improvement of the charge storage performance of α-Fe₂O₃ with the
capacitance retention above 90% after 3 · 10⁴ cycles.

**Fe₃O₄**, or FeO·Fe₂O₃, is a spinel oxide of iron, formed by the high temper-
ature reduction of Fe₂O₃. Its maximal charge storage capacitance, according to
(3.2) with *x* = 3 and *M* = 231.5 g/mol, *Q*<sub>C</sub> = 1250.4 C/g. This value is even
higher than the value of *Q*<sub>C</sub> for Fe₂O₃. Contrary to the hematite oxide, Fe₃O₄
has a high conductivity at the ambient conditions: 10 – 10² S.cm⁻¹ [141]. Thus,
Fe₃O₄ was proposed as a promising component for supercapacitors electrode ma-
terial [136]. Despite a relatively high conductivity, the stand-alone films of Fe₃O₄
nanoparticles (diameter is ca. 20 nm) demonstrated a moderate specific capac-
itance *C*<sub>g</sub> = 118.2 F/g and a weak capacitance retention of 89% after just 500
cycles in 1M Na₂SO₃ [142]. The moderate performance of Fe₃O₄ can be attributed
to the high interparticles resistance in the absence of the conductive component
of the electrode layer [143]. The other problem of utilisation of Fe₃O₄ as an elec-
trode material is a narrow (ca. 200 mV) potential window of its thermodynamic
stability, see the Pourbaix diagram (Fig.3.6(E)). At more positive potentials one
can expect the oxidation of Fe<sup>II</sup> sites into Fe<sup>III</sup> resulting in the performance loss.
The electrodes of Fe₃O₄ are often tested in the Na₂SO₃ electrolytes, because the
sulfite is a reductant and helps to stabilize the Fe<sup>II</sup> sites. Also, the fine dispersion
of the Fe₃O₄ nanoparticles on the carbon substrates allows to avoid the interpar-
ticles resistance complications and stabilizes the performance. The high specific
capacitance of the Fe₃O₄ nanoparticles with *d* = 5 nm, which are dispersed on the
reduced graphene oxide sheets, is *C*<sub>g</sub>=361 F/g at 1 A/g in 1M KOH. This high
specific capacitance is stable up to 10³ cycles at 5 A/g [144].

Case study: hierachial deposits of Fe₃O₄@graphene / PEDOT:PSS

In this subsection the hierarchial electrode material with the Fe₃O₄ nanoparticles
deposited on the graphene and arranged into the layered structure is described.
This material was developed in the common project between IPCMS, ICS and
ICPEES with the participation of the author [12].

The uniform deposits of the Fe₃O₄ nanoparticles on the few-layer graphene
<figure><img src="image_45.png" /><figcaption>Figure 3.7: The Fe₃O₄/FLG // PEDOT:PSS electrodes, prepared by the layer-by-layer deposition [12].</figcaption></figure>

(FLG) support were utilized to construct thick electrode layers up to 72 µm [12]. The regular structure of the deposit was preserved by the layer-by-layer deposition of the Fe₃O₄/FLG layers alternating with the PEDOT/PSS layers, which played the role of structurizing element (Fig.3.7). Preserving the structure of the layer resulted in the linear increase in the surface specific capacitance with the number of deposited layers. However, after more than 20 bilayers were deposited, the slope of the increase decreased due to the partial disturbance of the structure (Fig.3.8). The mass specific capacitance had the maximal value C_g=116 F/g in 0.5 M Na₂SO₃ for the deposit with 15 bilayers (Fig. 3.9). The stability of the performance was demonstrated by the retention of 114% for the initial capacitance after 1.5 · 10³ cycles at 1 A/g in 0.5M Na₂SO₃. The increase of the capacitance was attributed to the electrochemically induced recovery of the Fe₃O₄ chemical state. The chemical state of Fe₃O₄ was assumed to be altered due to the partial oxidation during the materials and final electrode preparations [12].

The oxyhydroxide of Fe<sup>III</sup>, **FeOOH**, occurs mainly in 2 allotropies: the goethite α-FeOOH and the lepidocrocite γ-FeOOH. These 2 oxides are formed spontaneously on the surface of iron as main components of the rust layer (together with hematite and magnetite) [145]. They can also be grown on the surface of the Fe electrode under the controlled electrochemical conditions [146]. The ratio of the components depends on the conditions of formation. The goethite is claimed to be a protective component, while the lepidocrocite – a dissolving component.

The conductivity of these oxides is low, ca. 10<sup>-9</sup> S.cm<sup>-1</sup> [147], which strongly limits their application in supercapacitors. The oxohydroxides of Fe are used as the supercapacitors electrodes in the form of nanoparticles supported on a conductive substrate. For example, the composites of the goethite nanoparticles on rGO
<figure><img src="image_46.png" /><figcaption>Figure 3.8: The typical SEM images of 15 layers of the Fe₃O₄/FLG // PEDOT:PSS deposit: (a) atop and (b) cross-section views; dependencies of the deposit thickness (c) and mass of Fe₃O₄ (d) on the number of deposited layers [12].</figcaption></figure>
<figure><img src="image_47.png" /><figcaption>Figure 3.9: (a) The typical CV curves of the 15-layered (Fe₃O₄/FLG // PEDOT:PSS) electrode in 0.5M Na₂SO₃ at 1 mV/s. The dependence of the specific capacitance of the electrode layer on the number of deposited layers: the mass specific capacitance (b) and the capacitance referred to the geometric surface area (c) [12].</figcaption></figure>
demonstrated the capacitance equal to 132 F/g at 0.5 A/g and the retention equal to 87% after 200 cycles in 0.1M Na₂SO₃ [148]. The nanorods of the goethite uniformly dispersed on rGO sheets demonstrated the capacitance above 450 F/g (at 1 A/g) in 1M KCl [149], and the retention of 87% of the initial capacitance after 10³ cycles at 2 A/g. It is claimed that the uniform distribution and fine contact of oxide material with the flexible rGO substrates is responsible for this impressive performance.

The doping of the goethite crystals with Co results in a significant improvement of the capacitance and its retention [150]: the deposit of 3.6% Co-doped α-FeOOH demonstrated the capacitance equal to 463.18 F/g at 0.1 A/g in 1M KOH and the excellent retention equal to 96.3% after 10³ cycles. The improvement of the performance was explained by tje influence of the dopant Co³⁺ cations on the energy level of the oxygen p-electrons and the decrease in the band gap by 0.24 eV.

The lepidocrocite γ-FeOOH is less stable than the goethite. However, the ad-
vantage of the lepidocrocite is its 2D-layered structure with the interlayer channels
facilitating the ion transport [136]. For example, the thick (30-50 nm) nanosheets
of γ-FeOOH with the length ca. 1.4 µm were electrochemically deposited on the
carbon cloth. The high value of the mass specific capacitance C_g = 310.3 F/g was
detected at 0.2 A/g and more importantly, only a mediocre drop of the capacitance
to 219.5 F/g was observed at very high charging rates of 12.6 A/g (both in 3M
KCl) [146]. This indicated the efficient ion transport within the electrode layer.

Oxides of Mn

The manganese oxides form the second most studied group of oxides (after RuO₂)
as electrode materials for supercapacitors. Numerous reviews were published ex-
clusively devoted to the manganese oxides electrode materials [151, 82, 152]. The
manganese oxides are relatively abundant and cheap, and environmentally neutral.
In contrast to the most of transition metals, Mn has a wide range of stable oxidation
states, from Mn⁰ as the metal to Mn⁷⁺ as the permanganate anion MnO₄⁻. Mn
forms several oxides stable in aqueous solutions: MnO (easily oxidizes to Mn₂O₃),
Mn₃O₄, Mn₂O₃, MnO₂. Another oxide Mn₂O₇ is a strong oxidant, reacting with
water to form MnO₂ and O₂.

The oxide of Mn⁴⁺, namely, **MnO₂**, is by far the most intensively utilized
oxide of manganese in the field of electrochemical energy storage. The cathode redox transition $MnO_2/Mn_2O_3$ is commonly used in the non-rechargeable alkaline batteries with the Zn anodes:

$$Zn + 2 MnO_2 = ZnO + Mn_2O_3$$

The reversibility of both Zn anodes and $MnO_2$ cathodes can be improved by the application of nanostructured materials, in particular, the $MnO_2$ nanofibers [153]. The retention of 92% of the initial capacitance of such batteries after 5000 cycles at 5C rate was reported [153]. The $MnO_2$-based cathodes are also commonly utilized as cathodes of the Li-ion rechargeable batteries [95, 154].

Since the work [91], where the capacitive behavior of the $MnO_2$ oxide in neutral electrolyte (2M KCl) was first demonstrated and $C_g = 199$ F/g and its 100% retention after 100 cycles were claimed, there were intensive studies of application of $MnO_2$ as electrode materials for supercapacitors. The maximal charge storage, assuming $1e^-$ per Mn atom was $Q_c=1102$ C/g or 918 F/g for the potential range of 1.2V, which was utilized in [91]. In [155] an even higher value of the capacitance equal to 1250 C/g was obtained for the thin deposit of amorphous particles of $MnO_2$. The amorphous hydrated $\alpha$-$MnO_2$ is typically formed by the low-temperature chemical synthesis or the electrodeposition from $Mn^{2+}$ salts. High observed values of $Q_c$ indicate the high efficiency of the charge storage in $\alpha$-$MnO_2$, related to the following suggested redox transition [155]:

$$MnO_x + \delta H^+ + \delta e^- \rightleftharpoons MnO_{x-\delta}OH_{\delta} \quad (3.11)$$

However, such high value $Q_c=1250$ C/g or the capacitance $C_g=1380$ F/g were reported for the thin films of $MnO_2$ only. For example, the low loading (13.3 $\mu g/cm^2$) $MnO_2$ deposit on the Pt substrate [155] showed the high capacitance, while the increase of the loading to 66.7 $\mu g/cm^2$ resulted in the film of the thickness equal to ca. 5 $\mu m$ with the capacitance decreased to 930 F/g. In the most of experimental studies, the films with the thicknesses around 100 $\mu m$ were used, and the capacitances $C_g \approx 130-150$ F/g were generally reported [104, 156]. Such significant decrease in the electrochemical capacitance with the deposit thickness was attributed to a limited conductivity of $MnO_2$, both electronic and ionic.

The conductivity of $MnO_2$ varies with its crystal structure. The rutile-type $\beta$-$MnO_2$ (pyrolusite) is an n-type semiconductor [157] with a low electronic con-
ductivity $\sigma \approx 10^{-6} \text{ Ohm}^{-1}\text{.cm}^{-1}$ [103]. $\beta$-MnO$_2$ consists of the MnO$_6$ tetrahedra forming the compact structure with (1x1) tunnels, i.e. a tunnel with dimensions close to the size of the MnO$_6$ tetrahedra, 2.3x2.3 Å. These relatively narrow tunnels provide not very high ionic conductivities. If MnO$_2$ is formed relatively fast in aqueous solutions, then water molecules can stabilize larger tunnels. The examples are the (2x2) tunnels in $\alpha$-MnO$_2$ (cryptomelane), the (2x1) tunnels in ramsdellite structure, or the alternating ramsdellite (2x1) and rutile (1x1) domains in $\gamma$-MnO$_2$ [95]. These polymorphs of MnO$_2$ are reported to have higher electronic and ionic conductivities [103] comparing to $\beta$-MnO$_2$. Upon the reversible cation insertion, the $\alpha$-MnO$_2$ structure with large (2x2) tunnels can become destabilized [95], resulting in the loss of performance. The ramsdellite and $\gamma$-MnO$_2$ show a better stability [158, 159].

The controlled annealing of the amorphous MnO$_2$, which is obtained by the reduction of the aqueous solution of KMnO$_4$ and contains a significant amount of $K^+$ inclusions, results in the formation of the $\delta$-MnO$_2$ oxide (birnessite) [160]. The birnessite has the layered structure of $\beta$-MnO$_2$ sheets separated by water and the ions insertion. In particular, the decrease of $K^+$ content in the birnessite, caused by its washing in acidic or neutral aqueous solutions, results in the increase of the inter-layer separation and the SSA of the material. $\delta$-MnO$_2$ shows the electrical conductivity comparable to that of $\beta$-MnO$_2$. However, $\delta$-MnO$_2$ has a significantly higher ionic conductivity [103].

The spinel oxide $\lambda$-MnO$_2$ can be obtained by the de-lithiation of the spinel LiMn$_2$O$_4$ by the hydrolysis in HCl. Despite the absence of large regular tunnels in the spinel structure, the high ionic conductivity of $\lambda$-MnO$_2$ was reported [103].

The electrochemical properties of MnO$_2$ depend significantly on its crystal structure. This dependence was studied in details in [160] and [103]. It is important to emphasize that it is difficult to isolate the effect of crystal structure, because the material capacitance depends also on SSA and porosity. Since the conditions of the synthesis of different polymorphs of MnO$_2$ are substantially different, the SSA and porosity of final materials varied too, which can explain certain disagreement between the studies [160] and [103]. Nevertheless, these studies present several common observations. In particular, the lowest specific capacitance was observed for the pyrolusite $\beta$-MnO$_2$ and explained by its low conductivity. $\beta$-MnO$_2$ is the most compact phase among the above mentioned oxides. Significantly higher capacitances were detected for more open phases, such as the cryptomelane $\alpha$-
MnO₂. Even a higher capacitance was reported for the birnessite δ-MnO₂. It is important to notice that the DFT modeling demonstrates that despite a low bulk conductivity, the surface conductivity of β-MnO₂ is much higher, because of the presence of the surface states [161]. Thus, the nanostructured MnO₂ mixed with a conductive component (carbon) makes a promising class of electrode materials for supercapacitors.

Most of the studies of the MnO₂ electrodes for supercapacitors were performed in neutral solutions, namely, chlorides or sulfates of K or Na. There are only few studies of the capacitance of the manganese oxides performed in alkaline electrolytes. It was observed [162] that the capacitance of the same material (MnO₂ deposited on the carbon nanotubes) was ca. 25% higher in 1M KOH than in 0.5M Na₂SO₄, while the same mechanism of charge storage was assumed. The electrodeposited manganese oxide also demonstrated a substantially higher specific capacitance in 1 M NaOH [163] comparing to 0.5M Na₂SO₄ [164]. The intriguing aspect of the studies in alkaline electrolytes is the existence of a relatively wide potential range of the stability for the reduced manganese oxides: Mn₂O₃ (at pH > 7) and Mn₃O₄ (at pH > 10.5) (Fig.3.6(D)). It was suggested that the redox transitions of these oxides could allow a more significant change of the oxidation state of Mn in the potential range of 1V [165], thus, increasing the pseudocapacitance of the oxide.

The detailed XPS analysis of the changes in intensities of the O 1s and Mn 2p XPS signals between the fully charged (oxidized) and discharged (reduced) α-MnO₂ in 0.1M Na₂SO₄ confirmed that the reaction (3.11) was the main mechanism of charge storage. This reaction results in the change of average oxidation state of Mn from 3.1 for the reduced state to 4 for the oxidized state. The high values of the detected *Q*<sub>C</sub> indicate that the reaction (3.11) can propagate into the bulk of MnO₂ particles. On the other hand, despite numerous claims in literature of electrolyte cation insertion into the bulk of MnO₂, the amount of Na and S detected by the XPS is attributed to the adsorption of the electrolyte cations Na⁺ and anions SO₄²⁻ exclusively on the surface of oxide [9].

The oxides of manganese with lower oxidation states, namely, the birnessite Mn₂O₃, MnOOH, and the spinel Mn₃O₄ are studied significantly less than MnO₂. The main reason is that the intrinsic electronic conductivity of these oxides is much lower [157]. In particular, the conductivity of Mn₃O₄ is σ = 10⁻⁷ Ohm⁻¹·cm⁻¹. The other problem is the instability of lower oxidation states of Mn²⁺/Mn³⁺. This
instability is caused by the irreversible oxidation to Mn⁴⁺ upon potential cycling, which was proved by the XANES measurements [166]. This effect was observed for both Mn₂O₃ and Mn₃O₄ oxides and resulted in the unexpected decrease of the mass specific capacitance (measured in 0.1 M Na₂SO₄) by 21% in 300 cycles: from 230 F/g to 180 F/g. It is interesting that a significantly higher capacitance of Mn lower oxidation state oxides was observed in alkaline electrolytes, comparing to neutral electrolytes. For example, Mn₂O₃ on the mesoporous carbon in 6M NaOH provided 651 F/g with 85% retention after 800 cycles [152]. A composite of the spinel Mn₃O₄ and the porous carbon demonstrated the specific capacitance equal to 600 F/g in 1M KOH, but at a very low charging rate (0.1 mV/s) [167]. The capacitance dropped to 340 F/g at 1 mV/s rate, indicating the limitations related to the low conductivity of Mn₃O₄. It is important to emphasize that in the studies of electrochemical properties of Mn₂O₃ and Mn₃O₄ in alkaline electrolytes, few pairs of CV redox peaks were clearly observed, contrary to MnO₂, which demonstrated relatively flat CV curves. The presence of the CV redox peaks clearly indicates the contribution of pseudocapacitance into the charge storage.

## Case study: Mn₃O₄/C composites as supercapacitors electrode material

In this subsection the synthesis and the properties of the spinel (hausmannite) Mn₃O₄ oxide in the composite with the Vulcan carbon are discussed. This composite was developed in the ICPEES and studied by the author in details [13].

The electrochemical properties of the Mn₃O₄ spinel oxide were studied thoroughly for the composite of the oxide nanoparticles and the XC-72 Vulcan carbon [13]. The composites were prepared by the solid state combustion synthesis (SSC), which is commonly used for the preparation of the oxide nanosized powders [168, 169, 170]. The SSC method has been adapted in the ICPEES by Dr. Kérangueven for the preparation of the composites of metal oxide nanoparticles with carbon, which are used as electrocatalysts in the oxygen reduction reaction [171]. The modified SSC method was named as the in situ autocombustion or the ISAC method.

In the original SSC method, a precursor salt of a transition metal (mostly nitrate or carbonate) is mixed with a fuel component. A fuel component is a substance capable to decompose at moderate temperature with a strong exothermic effect. Most commonly used fuel components are glycine (C₂H₅NO₂), citric acid
<table><caption>Table 3.2: The parameters for the calculation of reaction temperature (3.12) for the synthesis of Mn₃O₄ from the Mn nitrate with glycine as a fuel.</caption><thead><tr><th>Substance</th><th>Δ<sub>r</sub>H, kJ/mol</th><th>C<sub>p</sub>, J·mol<sup>-1</sup>·K<sup>-1</sup></th></tr></thead><tbody><tr><td colspan="3"><strong>Reagents</strong></td></tr><tr><td>Mn(NO<sub>3</sub>)<sub>2</sub>·4H<sub>2</sub>O</td><td>500</td><td>-</td></tr><tr><td>Glycine</td><td>-973</td><td>-</td></tr><tr><td colspan="3"><strong>Products</strong></td></tr><tr><td>Mn<sub>3</sub>O<sub>4</sub></td><td>-</td><td>44</td></tr><tr><td>C, graphite</td><td>-</td><td>8.5</td></tr><tr><td>N<sub>2</sub></td><td>-</td><td>31</td></tr><tr><td>CO<sub>2</sub></td><td>-</td><td>54</td></tr><tr><td>H<sub>2</sub>O</td><td>-</td><td>45</td></tr></tbody></table>

(C₆H₈O₇), urea ((NH₂)₂CO). The heat generated by the decomposition dissipates fast, providing enough heat to decompose the metal salt into well crystallized oxide particles and, at the same time, avoiding the sintering of the particles. The size of the particles depends on the weight ratio of salt to fuel [170]. The temperature T<sub>r</sub>, generated within the reaction media, can be estimated by the equation

$$T_r = T_0 + \frac{\sum_i (-\Delta_r H_i \cdot n_i)}{\sum_j C_{p,j} \cdot m_j}, \quad (3.12)$$

where T₀ is the externally applied temperature, i is the index of decomposing reagents, j is the index of the products of the decomposition reactions. The iᵗʰ reagent generates the heat –ΔₜHᵢ (in J/mol) and is present in the nᵢ molar quantity. The jᵗʰ product has the specific heat capacitance Cₚ,ⱼ (in J·mol⁻¹·K⁻¹) and the mass mⱼ. The parameters of the equation (3.12) for the synthesis of Mn₃O₄ by the decomposition of Mn(NO₃)₂·4H₂O with glycine as a fuel are given in Table 3.2. It is important to note that the equation (3.12) gives an estimate of the temperature and not the exact value, because of the assumption about a weak dependence of the specific heat capacitance on the temperature.

In comparison to the original SSC synthesis, the presence of carbon in the ISAC synthesis decreases the temperature of the media due to a high heat conductivity of the carbon. Using (3.12), the effect of the presence of the carbon in the reaction media is evaluated. The dependence of the reaction temperature on the weight percentage of the oxide in the Mn(NO₃)₂ - glycine - carbon mixture is shown in Fig.3.10. It appears that the loading of carbon up to ca. 70% decreases the
<figure><img src="image_48.png" /><figcaption>Figure 3.10: The dependence of the reaction temperature in the ISAC synthesis for the Mn(NO₃)₂·4H₂O - glycine - carbon mixture on the weight ratio of the oxide. The molar ratio of nitrate and glycine is fixed to provide the optimal conditions.</figcaption></figure>

temperature of the reaction media only by ca. 20%. It can be explained by a relatively low heat capacitance of the carbon, comparing to the other products of the ISAC reaction (see Table 3.2).

Indeed, the composites of Mn₃O₄/C can be prepared by the ISAC method in a relatively wide composition range: at least from 7%wt to 87%wt of the oxide [13]. The composite consists of the Mn₃O₄ nanoparticles with the size range 20-30 nm, which are distributed uniformly over the carbon substrate. As it was discussed in Chapter 3.2, the mixing of the components in the ISAC synthesis is much more efficient comparing to the mechanical mixing. Such effective mixing results in a high specific capacitance and its weaker dependence on the charging rate (Fig. 3.4).

The specific capacitance of these composites shows strong dependence on their compositions (Fig.3.11). In particular, the presence of a relatively small (18%wt) amount of the oxide in the composite results in the 2-fold increase of the specific capacitance. The maximum capacitance $C_g = 270$ F/g was obtained for the com-
<figure><img src="image_49.png" /></figure>

<figure><img src="image_50.png" /><figcaption>Figure 3.11: (A) The CV curves of the Mn₃O₄/C composites with various compositions measured in 1M NaOH, 20 mV/s. (B) The dependence of the specific capacitance of the Mn₃O₄/C composites in 1M NaOH on their composition [13].</figcaption></figure>
posite containing 34%$_{wt}$ of the oxide in Mn$_{3}$O$_{4}$/C (at 1 mV/s in 1M KOH ). Further increase of the oxide content resulted in the rapid decrease of the total capacitance due to the agglomeration of the oxide particles. The cyclic voltammetry of the Mn$_{3}$O$_{4}$/C composites clearly showed several pairs of the quasi-reversible redox peaks, indicating the contribution of pseudocapacitance into the charge storage ( Fig.3.11(A)).

By taking into account the specific capacitance of the carbon, measured sepa-
rately, the specific capacitance of oxide particles can be calculated from the total
capacitance of the composite. For the optimized composition the specific capac-
itance of the Mn₃O₄ nanoparticles was found to be equal to C₉=754 F/gₐₓide at
1 mV/s and 439 F/gₐₓide at 500 mV/s. The surface-related charge storage capac-
itance (including the EDL and the pseudocapacitance) for MnO₂ was estimated
to be equal to 100 µF/cm² [155]. Using the SSA value for 34%$_{wt}$ Mn₃O₄/C, the
contribution of the surface related capacitance to the total capacitance was esti-
mated to be equal to 114 F/g. Thus, it can be concluded that the bulk of the
Mn₃O₄ nanoparticles strongly participates in the charge storage in the composite.
Indeed, the maximal charge storage capacitance for Mn₃O₄ calculated by (3.2) is
Qc = 1264 F/g. The experimental value of C₉=754 F/gₐₓide indicates that more
than a half of the bulk Mn atoms are involved in the charge storage at least at
slow potential sweep rates.

The dependence of the charge stor-
age potential profile on the charging
rate can be analyzed by comparison
of the CV curves measured at differ-
ent sweep rates. For better compari-
son, the currents are recalculated to the
specific differential capacitance c<sub>g</sub> using
the following equation:

<figure><img src="image_51.png" /><figcaption>Figure 3.12: The dependence of the differential capacitance of the 34% Mn₃O₄/C composite on the potential sweep rate [13].</figcaption></figure>

$v$ is a sweep rate (in V/s) (Fig.3.12).

According to the Pourbaix diagram (Fig.3.6(D)), the peaks below -0.4V vs

$$
c_g = \frac{i_m}{v}, \qquad (3.13)
$$

where *i*<sub>*m*</sub> is the mass-specific current
density measured by the CV (in A/g),
Hg/HgO (0.52V vs. SHE) can be attributed to the reduction of $Mn^{3+}$ to $Mn^{2+}$. The peaks in the range -0.20..0.00V Hg/HgO (0.72 .. 0.92V SHE) can be attributed to the oxidation of $Mn^{3+}$ to $Mn^{4+}$. As Fig.3.12 shows, the pseudocapacitance related to the $Mn^{3+}/Mn^{2+}$ transition is slower, and its contribution is decreasing faster with the potential sweep rate increase, in comparison to the decrease for the $Mn^{4+}/Mn^{3+}$ transition. The latter transition provides a significant contribution even at 500 mV/s.

A more detailed insight on the nature of the pseudocapacitance for the $Mn_3O_4$ particles was obtained from the measurements performed in the set of $Na_2SO_4$ + NaOH electrolytes with various pH (Fig.3.13). It is important to note that the composition of electrolytes in [13] was adjusted to change the value of pH and keep a constant concentration of the $Na^+$ cation. The inset in Fig.3.13 shows the dependence of the open circuit potential $E_{OCP}$ of the $Mn_3O_4$/C electrode on pH. The systematic shift of $E_{OCP}$ equal to $80 \pm 20$ mV/pH was detected. For all values of pH, the value of $E_{OCP}$ was observed in the range of the $Mn_2O_3$ stability according to the Pourbaix diagram. This indicates that the $Mn_3O_4$ particles are partially oxidized, at least on the surface, during the preparation of the electrode.

The systematic decrease of the pseudocapacitance-related peaks of the CV with the decrease of pH (Fig.3.13(A)) was detected. This is in agreement with the above mentioned comparative studies of the $MnO_2$ capacitance in NaOH and $Na_2SO_4$ electrolytes, although for the $Mn_3O_4$/C composites the effect is much stronger. The observed shift of the CV peaks with pH $dE/dpH \approx 70 \pm 10$ V/pH confirms that the pseudocapacitance in the manganese oxides can be attributed to the process described by (3.11). It can be concluded that $OH^-$ plays an important role in this process. One may also detect that the capacitance decrease with the potential sweep rate increase gradually becomes more significant at lower pH (Fig.3.13). The separation between the anodic and cathodic peaks around 0.00V Hg/HgO also increases from 70 mV at pH=14 to 140 mV at pH=11. It allows to conclude that the pseudocapacitance processes become gradually slower with the pH decrease.

## Spinel oxides electrode materials

It is interesting to compare the capacitances reported in the literature for 3 spinel oxides: $Mn_3O_4$, $Co_3O_4$ [172], $Fe_3O_4$ [173] (see Table 3.3). $Mn_3O_4$ has the smallest
<figure><img src="image_52.png" /></figure>

<figure><img src="image_53.png" /><figcaption>Figure 3.13: (A) The CV curves of the 34% Mn₃O₄/C composite at 20 mV/s measured at various pH. (B) The decrease in the total specific capacitance with the potential sweep rate increase at various pH.</figcaption></figure>
<table><caption>Table 3.3: The comparison of the structure and properties of the spinel oxides</caption><thead><tr><th>oxide</th><th>spinel</th><th>lattice parameter, Å</th><th>density, g/cm³</th><th>specific capacitance, F/g</th><th>Ref.</th></tr></thead><tbody><tr><td>Mn₃O₄</td><td>normal</td><td>5.8</td><td>4.86</td><td>750</td><td>[13]</td></tr><tr><td>Co₃O₄</td><td>normal</td><td>8.1</td><td>6.07</td><td>400</td><td>[172]</td></tr><tr><td>Fe₃O₄</td><td>inverse</td><td>8.4</td><td>5.17</td><td>120</td><td>[173]</td></tr></tbody></table>

lattice parameter and the most compact crystal structure (as evidenced by its lowest density). Mn₃O₄ has also the highest reported mass specific capacitance. However, it is hard to explain a higher capacitance of the Mn oxide by its more compact structure. In fact, as mentioned above, an increase in the interatomic spacing for the MnO₂ oxides resulted in a higher specific capacitance of α-MnO₂ comparing to β-MnO₂.

The ionic conductivity is claimed to be an important parameter, determining the efficiency of the charge storage. However, among the 3 spinel oxides, the conductivity decreases in the following order: Fe₃O₄ > Co₃O₄ > Mn₃O₄ [174], inversely to the increase in capacitance. Thus, the factors other than ionic conductivity play a more important role.

As discussed above, the redox transformation, which is responsible for the pseudocapacitance in Mn₃O₄, is a pH-dependent reaction (3.11). This reaction involves OH⁻ as a reagent. Thus, one can expect that the rate of reaction (3.11) depends on OH⁻ interaction energy with the oxide surface. It is interesting to note that Co₃O₄ and Mn₃O₄ are considered among the most promising catalysts for the oxygen electroreduction in the alkaline solutions [175]. The modeling of the interaction of water with the Mn₃O₄ cluster demonstrates, that the water molecules are significantly stronger and dissociatively adsorbed on the Mn²⁺ sites of Mn₃O₄ than on the Mn³⁺ sites [176]. This can explain a higher activity of the Mn₃O₄ oxide in the oxygen electroreduction in comparison to the MnO₂ and Mn₂O₃ oxides, and, more importantly for the present work, its higher pseudocapacitance in alkaline solutions. Moreover, a higher activity of the substituted spinel MnFe₂O₄ comparing to Fe₃O₄ in the oxygen electroreduction was reported [177], confirming that the reaction is facilitated in the presence of the Mn²⁺ sites, most probably, due to a stronger water adsorption. Similarly, a higher activity of the MnCo₂O₄ spinel comparing to Co₃O₄ was demonstrated [178].
It allows to suggest that the dissociative adsorption of water influences the rate of the process of pseudocapacitance in the TMO nanoparticle.

In Section 3.3.4 the impedance of Fe₃O₄/C and Mn₃O₄/C electrodes is discussed. These systems are considered as the examples of systems with significant contribution of pseudocapacitance to the total charge storage.

### 3.3.4 Impedance of porous electrode/electrolyte interface with pseudocapacitive behavior

The generalized staircase model of the impedance of porous electrodes is described in Section 1.3. This model predicts that the influence of the porous structure on the measured impedance can be detected at the low frequencies under the condition that the redox pseudocapacitance reaction is sufficiently fast. However, it is difficult to fit the measured EIS spectra quantitatively. As for the carbon electrodes, the macro-structure of the deposit of the material must be taken into account 2. Thus, only qualitative interpretation of the experimental Nyquist plots is given below.

The dependence of the impedance on the rate of the interfacial reaction is shown in Fig.1.22(A) in Section 1.3.

Fig.3.14 shows the impedance spectra for 2 various electrodes with the strong contribution of pseudocapacitance: the Fe₃O₄/FLG 15-layered deposit with the PEDOT as a structurizing agent (Fig.3.14(A)) and the Mn₃O₄/C deposit (Fig.3.14(B)). The spectra were measured at the potential of the open circuit, i.e. close to the equilibrium potential of the redox transition of the oxide materials (in the absence of other redox active components in the system).

In both cases the low frequency segments of the Nyquist plots show the straight inclined line, indicating the capacitive behavior. A slight inclination corresponds to the CPE constant $a = 0.90$ for Fe₃O₄/FLG and $a = 0.80$ for Mn₃O₄/C electrodes. Higher values of $a$ for Fe₃O₄/FLG are expected, because the Fe₃O₄/FLG deposit has the ordered layered structure with the PEDOT interlayers, thus, having the open regular porous structure. In comparison to this material, the structure of Mn₃O₄/C is more disordered, resulting in a lower $a$ value.

At high frequencies a semi-circle segment of the Nyquist plot for Fe₃O₄/C electrode is observed (Fig.3.14). This behavior is predicted by the generalized staircase impedance model (Fig.1.22(A)). The transition from the semi-circle to the
<figure><img src="image_54.png" /><figcaption>Figure 3.14: (A) The EIS spectra of the (Fe₃O₄@FLG/PEDOT)₁₅ electrodes in 0.5M Na₂SO₃ [12]. (B) The EIS spectra of the Mn₃O₄/C electrode in 1M NaOH [13]. The insets show the high frequency range of the EIS spectra.</figcaption></figure>
straight inclined segment for Mn₃O₄/C is observed at the higher frequencies than for Fe₃O₄/C. This difference indicates a faster kinetics of the pseudocapacitance reaction (i.e. a lower Rct) for Mn₃O₄/C. This observation correlates well with the significantly higher values of Cg for Mn₃O₄/C comparing to Fe₃O₄/C.

The application of 10000 charging cycles results in the strong loss of capacitance for the Mn₃O₄/C electrodes due to low retention. This is in contrast with the high retention of the Fe₃O₄/C electrodes under the similar conditions. This difference is reflected in the impedance spectra. For Mn₃O₄, the strong increase of the diameter of semi-circle is detected, reflecting the increase of Rct, i.e. the significant decrease in the rate the pseudocapacitance reaction. The semi-circle part of the EIS spectra for Fe₃O₄/C barely changes, except the higher values of Cdl result in the transition to the capacitive behavior at the lower frequencies.

## 3.4 Conclusions

Due to the pseudocapacitance phenomena, the transition metal oxides are the promising materials for the application as electrodes in supercapacitors. The pseudocapacitance involves the redox reaction of the transition metal cation and the intercalation of the electrolyte cation (for example H⁺, or Na⁺, K⁺). The tuning of the TMO structure allows to achieve a high utilisation of the charge storage capacity of the material. However, the following important points have to be taken into account.

* *Well crystallized oxides.* One of the factors limiting the efficiency of the TMO in supercapacitors is their low electronic conductivity. Electronic conductivity depends on the oxide structure and crystallinity: amorphous oxides are generally less conductive due to a higher resistance at the crystal grain boundaries. Thus, well crystallized oxides are preferable as the electrode materials for supercapacitors.

* *Oxide/carbon composites.* Even for relatively high intrinsic electronic conductivities, the interparticles resistance, related to the small area of contact between particles, results in the high serial resistance. Thus, the efficient strategy is to mix the TMO particles with a conductive material, in particular, with carbon. The TMO/C composites are among the most promising materials for supercapacitor electrodes. It is important that the TMO
nanoparticles are distributed uniformly over the carbon, in order to avoid the agglomeration and interparticle contacts.

* *A weak influence of the ionic conductivity.* Contrary to the frequent claims in the literature, the ionic conductivity of TMO seems to have little or even no influence on the TMO performance in supercapacitors. This fact is derived from the evaluations of the conditions for the ionic transport in the metal oxide particles under the uniform potential distribution. For this evaluation the solid state diffusion coefficient values reported in the literature were used.

* *The fast kinetics of interfacial charge transfer.* The important factor of the TMO performance as supercapacitor materials is the kinetics of the pseudo-capacitance reaction, involving the interfacial charge transfer and preceding the ion intercalation. The impedance modeling (Section 1.3) demonstrates that the fast kinetics is necessary to observe a significant impact of the pseudo-capacitance at the potential modulation rates typical for the supercapacitors utilization. This conclusion is supported by the frequent observation of the correlation between the capacitance of oxide materials and their electro-catalytic activities in the reactions with electrolyte components, such as the oxygen electroreduction.

* Electrochemical stability. Reversibility of the charge/discharge process is another important characteristic of the supercapacitors performance. Thus, the retention of initial capacitance at least 80% after $10^5$ charging cycles is one of the crucial requirements. However, a significant decrease in the capacitance of the TMO electrodes with the cycling is often observed. This decrease is caused by the degradation of the TMO. The degradation can be related to the irreversible formation of an oxide phase with a lower conductivity or the formation of a soluble phase. The analysis of the stability of the oxide phase can be done using the Pourbaix diagrams (Fig.3.6). The most of TMO are stable in neutral or weak alkaline electrolytes, thus, limiting the choice of aqueous electrolytes. The oxides of Mn, Fe, Co and Ni are among the oxides, which show a sufficient stability in neutral/alkaline electrolytes and the reversible redox transition. From the viewpoint of sustainable development, the first two oxides in this group are of a special interest due to their high abundance and low toxicity.
* *Spinel oxides as electrode materials.* The spinel oxides $Mn_3O_4$ and $Fe_3O_4$ are the most interesting candidates among the binary oxides of Mn and Fe for using as supercapacitors electrode materials. This is due to their relatively high ionic and electronic conductivity. In particular, the $Mn_3O_4$/C composite demonstrates the specific charge capacitance of the oxide particles above 750 F/g at 1 mV/s in 1M NaOH. Therefore, the $Mn_3O_4$ nanoparticles demonstrate the high efficiency of charge storage. The capacitance of the magnetite $Fe_3O_4$ is not very high, ca. 120 F/g. However, the magnetite is outstandingly stable: the retention of 114% of the initial capacitance was detected after 1500 cycles at 1A/g charging rate. The development of the composites with the substituted spinel oxides, based on the Mn and Fe spinels, such as $MnFe_2O_4$, is a promising approach towards new electrode materials.
# Chapter 4

## Future projects

The transition towards the energy networks based on the renewable primary sources, such as solar energy and wind, is one of the most pressing challenges of our time. Considering the intermittent character of these sources, the efficient and reversible energy storage devices are necessary and important components of such network. The electrochemical energy storage devices, in particular batteries, supercapacitors and hybrid cells are known to provide an outstanding energy storage reversibility and efficiency, compact structure and zero operation emission. In particular, supercapacitors can provide the energy density above 10 W.h/kg at high power output (higher than 1 kW/kg) and keep at least 80% of their performance after 10⁵ charging cycles.

The performance of supercapacitors is determined by the properties of electrode materials, such as electronic and ionic conductivity, specific surface area, accessibility for electrolyte, electrochemical stability and structural integrity during charging/discharging [1]. Thus, the development of supercapacitors with high performance requires the optimization of these properties of electrode materials.

Two classes of electrode materials play a key role in the development of supercapacitors: the carbon materials and the transition metal oxides. In the current section the author presents the planned projects which are the continuation of the current activities in the development of new materials for supercapacitors electrodes.
## 4.1 Bio-sourced activated carbon

The carbon materials have a high electronic conductivity and electrochemical stability. They are used as electrode materials in the electrochemical double layer capacitors (EDLC). Their performance is limited by a moderate charge storage capacitance, which is less than 15 µF/cm². Thus, a high specific surface area (SSA), which is readily accessible for an electrolyte, is crucial for the high energy density.

The activated carbon (AC) materials provide outstanding SSA values (above 2000 m²/g), resulting in high specific charge capacitances. The AC carbons are prepared by the physical (thermal) or chemical activation (see Section 2.4). The chemical activation provides a better control on the porosity of the final material and requires a lower energy input. Thus, the chemical activation is generally preferred for the AC synthesis. An additional advantage of the chemical activation method is that relatively cheap and abundant sources of carbon can be used for the mass-scale AC synthesis. For example, the AC electrodes with high charge capacitances (close to 300 F/g), excellent retention of capacitances with fast charging rates up to 20 A/g and with charging cycles up to 10⁴ were prepared from a fig fruit [10]. This approach is very useful from the viewpoint of the valorisation of bio-mass wastes: agricultural, agro-industrial, urban and wood production wastes [179, 180, 181].

In the recent years, the particular interest has been attracted to the microwave-assisted chemical carbon activation due to its high yield, short synthesis time, and low energy input [180, 181]. When the KOH/NaOH-etching approach is assisted by the microwave treatment, the AC with SSA large than 2000 m²/g and the high contribution of micropores are synthesized.

The tuning of the conditions of the microwave-assisted synthesis of the AC from biomass wastes requires careful study, considering the variability, complex structure and composition of biomass sources. This study is proposed as a part of the future project devoted to the development of new carbon electrode materials for supercapacitors.

In particular, the food waste treatment is very beneficial from the viewpoint of sustainable development [89]. The food waste is a renewable source, highly rich in cellulose and sugar. Contrary to agricultural wastes, the food waste is rich in the lignocellulose. The valorisation of food wastes decreases the landfill pressure. The
thermal pretreatment (in the air with $t^\circ < 300^\circ C$) is needed for standardization of waste sources and decrease of water content in the precursor.

The activation conditions must be tuned in order to obtain a desired AC structure with high SSA, controlled porosity, structural integrity and electrochemical stability. For example, the addition of NaCl in a certain proportion to the carbon source increased the graphitization of the final carbon product and its charge capacitance and electrochemical stability [11]. The control of the material porosity is important to provide the possibility of applying the final material as a carbon electrode in supercapacitors or using it as a component for the TMO/C composite materials. In the first case, the mircoporosity with the pore size $d < 1$ nm is favored in order to provide a high SSA. For the synthesis of the TMO/C composites, larger pores are required in order to provide the substrate for TMO nanoparticles growth.

Another important aspect of tuning of conditions for the AC synthesis is the optimization of its costs and energy input. Such optimization provides the potential for mass scale production of bio-sourced ACs for electrodes materials for supercapacitors.

## 4.2 TMO/C composite materials

The TMO materials are another most important class of materials for supercapacitor electrodes due to the TMO high energy density related to the phenomenon of pseudocapacitance. The pseudocapacitance involves a reversible redox transformation of the transition metal cation and/or intercalation of cation from an electrolyte into the sub-surface and the bulk of oxide phase. High rate and reversibility of charge storage reaction are crucial for the performance of supercapacitors. These requirements impose the necessity of shortening the ion diffusion path in the TMO bulk. Thus, the nanostructured TMO (nanoparticles, nanotubes, nanowhiskers, ...) are commonly used as electrode materials. For the nanoparticles, the optimal size is found in the range of 10-20 nm, providing a sufficiently large volume for the charge storage (high energy density) and a relatively short ion diffusion path (for high power output). The main factors, which limit the performance of the TMO as supercapacitor materials, are electronic conductivity (many TMO are semiconductors) and electrochemical stability.

The best performance in supercapacitors is reported for $RuO_2$ and related
oxides [115, 114, 110]: the charge capacitance larger than 900 F/g and the retention more than 80% after 10⁵ cycles. However, these oxides are too expensive and toxic for wide commercialization. The Mn oxides (mostly MnO₂, while lower oxidation degree oxides, such as Mn₂O₃, Mn₃O₄, MnOOH, have also been studied) also demonstrates the high charge capacitance larger than 600 F/g [9]. The high abundance and low toxicity of Mn make its oxides even more desirable from the viewpoint of sustainable development. In this context, the oxides of Fe are also considered as a promising candidate for supercapacitor electrodes materials. The tow electrical conductivity of Fe₂O₃ prevents its utilisation in supercapacitors. However, a more conductive β-FeOOH was studied as an electrode material for supercapacitors [82, 150].

The spinel oxides, namely the oxides with the formula AB₂O₄, where B is, most commonly, a trivalent transition metal, and A is a divalent metal cation, are of particular interest. These oxides often demonstrate a relatively high electronic conductivity and electrochemical stability. Namely, the spinels with B=Mn³⁺, Fe³⁺, and A=Mn²⁺, Cu²⁺, Mg²⁺ showed the high conductivity [174]. For example, the spinel Mn₃O₄ was studied as electrodes materials for supercapacitors, and the high charge capacitance was reported [182, 183].

High loading of electrode materials (above 100 mg/cm²) requires relatively thick deposits of the TMO nanoparticles. Even for relatively conductive oxides, the interparticles resistance results in high serial resistance of such layers and requires the utilisation of a conductive additive, for example, the carbon. The carbon substrates with the well developed surface provide a possibility to immobilize the TMO particles and avoid their agglomeration. Thus, the development of the TMO/C composite materials is a promising strategy to obtain the electrodes materials with high performance in supercapacitors.

As mentioned above, the TMO nanoparticles with the size in the range of 10-20 nm are preferred. The well crystallized particles are expected to have a better performance, as the conductivity of the amorphous oxides is generally lower. Unfortunately, the preparation of the well crystallized TMO particles often require the thermal treatment at the elevated temperature, resulting in ripening, sintering, and agglomeration of the particles.

The solid state combustion (SSC) synthesis is a preparation method providing the possibility to form the crystallized TMO nanoparticles while avoiding the risk of the particles sintering. The SSC method was upgraded in the ICPEES for the
preparation of the TMO/C composites. In the upgraded method, the precursors of metal (nitrate or carbonate salts) are dissolved with a combustion additive (glycine, urea, or citric acid). After that, the carbon is ultrasonically dispersed in this solution. Then, the solvent is evaporated until the dry uniform mixture is obtained. At the next step, the fuel agent is ignited at a relatively low temperature (100° < t° < 300°C), providing a local heat, which is necessary to decompose metal salts into a corresponding oxide. A fast evacuation of the heat allows to avoid the sintering of the TMO particles.

The upgraded SSC method was applied to synthesize the Mn₃O₄/C composite with controlled structure and composition. This composite demonstrated a high charge capacitance (larger than 750 F/g<sub>oxide</sub>) in 1M KOH, but a low capacitance retention (ca. 60% after 1000 cycles) [13]. The upgraded SSC method was also used for the preparation of the LaMnO₃/C composites for the oxygen electroreduction, proving that the composites with ternary oxides can be synthesized by the SSC.

In the planned project, the SSC method will further extended to the preparation of the spinel TMO/C composites with the general composition MₓFe₂ₓO₄, where M = Mn²⁺, Cu²⁺, Zn²⁺, Ca²⁺. The ferrite spinels are characterized by a high electrochemical stability (for example, the Fe₃O₄ oxide is used as a supercapacitor electrode [12]) and relatively high electronic conductivity. The capacitance of the MnFe₂O₄ oxide is reported to be higher than that of Fe₃O₄ [184, 185, 186] due to the stronger interaction of Mn²⁺ with water molecules and more facile cation insertion.

The extended SSC method will be applied to synthesize the MₓFe₃₋ₓO₄ oxides. The tuning of the ratio of metal precursors and combustion additives will be performed to ensure the control of the composition, including the substitution degree *x* and the crystal phase of the final oxide. After that, the method will be utilized to prepare the composites of MₓFe₃₋ₓO₄/C with the controlled composition and structure of TMO particles. Both commercial carbon (for example, the XC-72 Vulcan) and the bio-sourced AC, prepared in the framework of the planned project, will be used for the composites.

## 4.3 Electrochemical study of supercapacitors performance

Firstly, the supercapacitors tests will include the utilisation of the model 3-electrode configuration with the electrode loadings below 1 mg/cm². This will allow to avoid
the problems related to a possible non-uniformity of the thicker deposits. Also, the
measurements in the 3-electrode cell allow to address more precisely the properties
of the electrode material. This configuration also allows to avoid the complications
related to the presence of a separator with low ionic conductivity comparing to
a liquid electrolyte. In particular, the impedance spectroscopy in the optimal 3-
electrode configuration allows to minimize the contributions of impedances of the
counter electrode and the electrolyte.

The 3-electrode configuration will allow to determine the initial specific charge capacitance (from the CV measurements) and its dependence on charging rate and on the number of applied charging cycles (from the galvanostatic cycling). Using the EIS method, the mechanism of charge storage will be evaluated (EDL or pseudocapacitance). Moreover, the factors, which limit the contribution of pseudocapacitance for the TMO/C materials will be determined.

The 2-electrode configuration will allow to test the electrode materials in the conditions of z real supercapacitor, including high (above 1 mg/cm²) loadings of the electrodes. It will also allow to employ the physicochemical methods (XRD, XPS, SEM/TEM) to perform the post-mortem analysis of the electrode materials. This will allow to detect the changes in the materials composition and structure and provide the valuable information about degradation mechanisms.

The results of the EIS measurements in the 3-electrode configuration will be an-
alyzed using the generalized staircase impedance model, developed in the current
work. This will allow to have a deeper insight on the dependence of the elec-
trochemical properties of the complex electrode materials (carbons and TMO/C
composites) on the parameters of their porosity. On the other hand, the model
will be also developed further in order to take into account the macrostructure of
the thick deposits of the electrode materials. Ultimately, the generalized staircase
model will be applied for the quantitative description of the impedance of thick
electrode materials in both 3- and 2-electrode configurations.
Acknowledgments

First of all, I would like to express my deepest gratitude to the head of Electrocatal-
ysis and Energy Conversion group of the ICPEES, **Prof. Elena R. Savinova**.
During more than 15 years of our work together, she has been always giving me
all the support and autonomy necessary to develop myself as an independent re-
searcher and to work out my own scientific personality. I am both honored and
privileged to work under the Elena's supervision.

I would like to thank most sincerely **Dr. Phamm-Huu Cuong**, who accepted the burden of being my garant for this work. I am also very grateful to Cuong for sparkling my interest to the topic of electrode materials for EESD applications, for inviting me to participate in the projects related to this field, and for being always ready to share his experience and knowledge in the field of carbon materials.

I would like to express my appreciation to **Prof. Aziz Dinia** (IPCMS, Strasbourg), **Dr. Frederick Favier** (ICGM, Montpellier), **Dr. Camélia Matei Ghimbeu** (IS2M, Mulhouse) and **Prof. Philippe Serp** (ENCI-ACET, Toulouse) for their acceptance of being the members of my jury, for their attention to my work.

This work would not be possible without active participation of other members
of ICPEES. I am grateful to **Dr. Gwénaëlle Kérangueven** for helping me to
learn ISAC synthesis and working together on the further development of this
synthetic route, pioneered out by her the ICPEES. I also have a privilege to work
with very skillful and helpful colleagues of the group of Dr. Pham-Huu, namely
**Housseinou Ba** and **Lai Truong Phuoc**, who are always ready to help, to share
the ideas and their knowledge of carbon materials.

I am grateful to **Thierry Romero** for all the sessions of SEM/TEM analysis
of our complicated samples. The analysis of our samples by XRD would not be
possible without guidance of **Dr. Thierry Dintzer**, who is also always very
helpful and creative in solving all the problems related to informatics and fitness.

By this work I am strongly indebted to **Dr. Nina Shokina**, my wife. She has supported and helped me throughout the whole long period of the preparation of this work. She has also significantly contributed to this work. An excellent mathematician, she worked out and programmed the algorithm of the calculations by the generalized staircase model. This important part of my work would not be possible without her impact. She has also gladly volunteered to undertake a huge work on testing and improving the comprehensibility of the text of my thesis.

Finally, I would like to express my everlasting gratitude to my first tutor after
my PhD, [Prof. Dr. Thomas Wandlowski], who died too soon in the summer
of 2015 at the age of 57. Besides having strongest impact on my development as
a chemist and a researcher, Thomas was also a very devoted and honest friend,
which will never be forgotten.
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