*Article* **Non-PGM Electrocatalysts for PEM Fuel Cells: A DFT Study on the Effects of Fluorination of FeNx-Doped and N-Doped Carbon Catalysts**

**Mohamed Cherif, Jean-Pol Dodelet, Gaixia Zhang , Vassili P. Glibin, Shuhui Sun and François Vidal \***

Centre Énergie, Matériaux, Télécommunications, Institut National de la Recherche Scientifique, 1650 Bd. Lionel-Boulet, Varennes, QC J3X 1S2, Canada; mohamed.cherif@inrs.ca (M.C.); jean-pol.dodelet@inrs.ca (J.-P.D.); gaixia.zhang@inrs.ca (G.Z.); vassili.glibin@gmail.com (V.P.G.); shuhui.sun@inrs.ca (S.S.)

**\*** Correspondence: francois.vidal@inrs.ca

**Abstract:** Fluorination is considered as a means of reducing the degradation of Fe/N/C, a highly active FeNx-doped disorganized carbon catalyst for the oxygen reduction reaction (ORR) in PEM fuel cells. Our recent experiments have, however, revealed that fluorination poisons the FeNx moiety of the Fe/N/C catalytic site, considerably reducing the activity of the resulting catalyst to that of carbon only doped with nitrogen. Using the density functional theory (DFT), we clarify in this work the mechanisms by which fluorine interacts with the catalyst. We studied 10 possible FeN<sup>x</sup> site configurations as well as 2 metal-free sites in the absence or presence of fluorine molecules and atoms. When the FeNx moiety is located on a single graphene layer accessible on both sides, we found that fluorine binds strongly to Fe but that two F atoms, one on each side of the FeNx plane, are necessary to completely inhibit the catalytic activity of the FeNx sites. When considering the more realistic model of a stack of graphene layers, only one F atom is needed to poison the FeNx moiety on the top layer since ORR hardly takes place between carbon layers. We also found that metal-free catalytic N-sites are immune to poisoning by fluorination, in accordance with our experiments. Finally, we explain how most of the catalytic activity can be recovered by heating to 900 ◦C after fluorination. This research helps to clarify the role of metallic sites compared to non-metallic ones upon the fluorination of FeNx-doped disorganized carbon catalysts.

**Keywords:** oxygen reduction reaction; proton exchange membrane fuel cell; fluorination; density functional theory; non-noble metal catalyst; N-doped carbon catalyst

#### **1. Introduction**

While promising non-noble metal catalysts for the oxygen reduction reaction (ORR) in proton-exchange membrane (PEM) hydrogen fuel cells have been synthesized over the years [1–5], their stability in fuel cells remains the main obstacle to their widespread use [6,7]. One of the most promising non-noble metal catalysts synthesized to date is FeNx-doped disorganized carbon [8–11]. There are several experimental and theoretical pieces of evidence that the Fe atom is the site where the ORR takes place [11–19]. It has been observed that this type of catalyst suffers from a decrease of almost half of its activity in a few hours of operation in fuel cells, followed by a much slower decrease thereafter. The current delivered by the fuel cell versus time can be fitted by a double exponential decay [20]. Few hypotheses have been put forward to explain the first rapid decay of catalytic activity. These include demetallation of the metal catalytic sites [20–23] and chemical reactions with H2O<sup>2</sup> [24–27]. The slower decay has not attracted as much interest as the fast one to date. Recent simulation work suggests that planar M3(C6O6)<sup>2</sup> [28] and M3(C6S3O3)<sup>2</sup> [29] structures, where M is a transition metal, may also be promising candidates but these have not yet passed the test of experiment.

**Citation:** Cherif, M.; Dodelet, J.-P.; Zhang, G.; Glibin, V.P.; Sun, S.; Vidal, F. Non-PGM Electrocatalysts for PEM Fuel Cells: A DFT Study on the Effects of Fluorination of FeNx-Doped and N-Doped Carbon Catalysts. *Molecules* **2021**, *26*, 7370. https://doi.org/10.3390/ molecules26237370

Academic Editors: Jingqi Guan and Yin Wang

Received: 26 October 2021 Accepted: 1 December 2021 Published: 4 December 2021

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There are indications that the fluorination of materials in acidic media improves their oxidative stability. Examples of such systems include Nafion ionomer, Pt/C, and platinum group metal-free catalysts for PEM fuel cells [30–32]. Recently, we also fluorinated a highly active FeNx-doped carbon catalyst in the hopes that fluorine would increase its stability in PEM fuel cells [33]. However, even after a short (2 min) exposure to a room-temperature F2:N<sup>2</sup> (1:1; *vol.*) gas stream, fluorination considerably inhibited the catalyst performance in H2/O<sup>2</sup> PEM fuel cells. Even if these experiments did not yield the expected results, they enabled several important observations to be made regarding the properties of the catalytic material under study:


The observations made in the context of these fluorination experiments have in fact provided a unique opportunity to improve our understanding of the nature of our FeNxdoped carbon catalysts and of the decay mechanisms of their catalytic activity in PEM fuel cells. In order to support and deepen the conclusions of our experimental study, in this paper we report density functional theory (DFT) calculations, based on the current understanding of the atomic structure of the catalytic sites and processes, and study the catalytic properties of these sites in the absence/presence of adsorbed fluorine.

Several theoretical studies have already focused on MNx-doped carbon catalysts, most often Fe [34–38], Co [35], Mn [35,39], and Ni [35]. Per the indications of several experimental studies [34,39–41], they conclude that the catalytic site is, specifically, the M atom within a functional group MN<sup>x</sup> embedded in a planar carbon layer. It is generally thought that the ORR catalyzed on these sites follows the four-electron exchange process [42–45]

$$\begin{array}{ccccc} \ast + \text{O}\_{2} + 4\text{(H}^{+} + e^{-}) & \rightarrow & \ast \text{O}\_{2} + 4\text{(H}^{+} + e^{-}) & \rightarrow & \text{O}\_{2} \\ & \text{I} & \text{II} & & \\ \ast \text{COOH} + 3\text{(H}^{+} + e^{-}) & \rightarrow & \ast \text{O} + \text{H}\_{2}\text{O} + 2\text{(H}^{+} + e^{-}) & \rightarrow & \text{O}\_{2} \\ & \text{III} & & \text{IV} & \\ \ast \text{OH} + \text{H}\_{2}\text{O} + (\text{H}^{+} + e^{-}) & \rightarrow & \ast + 2\text{H}\_{2}\text{O} \\ & \text{V} & & \text{VI} & \\ \end{array} \tag{1}$$

where \* denotes the adsorption site and the labels I to VI refer to the six reaction steps.

For several MNx-doped carbon structures, it has been found that, at low enough potentials, the free energy of each step of the reaction sequence (1) decreases uniformly from the first to the last step, indicating that the reaction sequence (1) is thermodynamically viable at these potentials. Other possible pathways, such as those involving spontaneous O<sup>2</sup> dissociation or H2O<sup>2</sup> formation, are less likely due to the increase in free energy at some stage of the process [24]. Several DFT studies have also been carried out for catalysts without metal [45–51]. These generally consider N-doped carbon structures and assume that the reaction sequence (1) still takes place at low enough potentials. These catalysts appear to be thermodynamically viable for some carbon sites near a nitrogen atom. However, O<sup>2</sup> adsorbs weakly or not at all on the catalytic sites (step II in the sequence (1)). This characteristic likely explains, at least in part, the much lower activity of these sites compared to the higher activity obtained with metal sites.

In a recent work, we theoretically studied the fluorination of two single-layer porous FeN4-doped carbon structures, one with pyrrolic nitrogen atoms and the other with pyridinic nitrogen atoms at the FeN<sup>4</sup> sites, and we assumed that the catalytic reaction took place through the sequence (1) [52]. Subsequent work has investigated ORR for various adsorbates bound to several transition metals on an MN<sup>4</sup> site [53,54]. However, actual catalysts most likely contain many embodiments of MN<sup>x</sup> moieties as well as M-free N

atoms in carbon layers, and involve more than one carbon layer. In order to provide a more complete picture of the fluorination process and of its influence on ORR, we performed DFT calculations for nine additional atomic structures of FeNx-doped carbon sites with x between 1 and 4 as well as for two N-doped metal-free carbon structures. We also examined the possibility of whether the ORR can be catalyzed on an Fe site between two parallel carbon layers in the presence of a F atom bound on Fe on the opposite side.

It is generally believed that the catalytic sites for the ORR of the FeNx-doped carbon materials consist of planar FeN<sup>x</sup> moieties located in a carbon structure which is generally approximated by a single carbon layer. Figure 1 shows some of the possible embodiments of such structures. Of course, the set of structures shown in Figure 1 is not exhaustive. Many variants of each structure are possible, such as, for example, pores in the carbon layer (as in Figure 1j), the FeN<sup>x</sup> moiety being near the edge of the carbon layer (as in Figure 1g,i), and N atoms being randomly distributed in the carbon layer (as in Figure 1e). Also, the N atoms surrounding the Fe ion may be either of a pyridinic type (Figure 1a–i) or of a pyrrolic type (Figure 1j). We will consider the 10 structures shown in Figure 1, expecting that they will be representative of the main effects of fluorination on the properties of the catalysts. For the purpose of the following discussion, special attention will be paid to the structure in Figure 1j. For this structure, the F–N bond length is 2.00 Å and the N–Fe–N angles are 197.19◦ and 82.81◦ .

**Figure 1.** DFT optimized configurations of FeNx-doped carbon catalysts investigated in this work. Panels (**a**–**j**) refer to the 10 configurations considered in this work. Color code: grey is carbon, blue is nitrogen, orange is iron.

It can be found in the literature that the free energy at zero potential of the steps of the reaction sequence (1) for the pristine structures of Figure 1c [55], 1h [56], and 1j [52] is uniformly descending at each reaction step (see Section 2.1). The structure of Figure 1b has also been investigated and was found to be not thermodynamically viable [55] (because the free energy presents a minimum at step V of the reaction sequence (1)—see Section 2.1). The other sites shown in Figure 1 have, to the best of our knowledge, never been investigated so far.

Figure 2 shows the two metal-free nitrogen-doped carbon catalytic sites considered in this work. As reported above, the activity of the fluorinated Fe/N/C catalysts became similar to that of Fe-free nitrogen-doped carbon catalysts. It is assumed that the reaction sequence is given by (1) where \* now denotes an active carbon site. Other nitrogendoped carbon structures have been investigated and shown to be potential catalysts for ORR [46–49]. The structures shown in Figure 2 were selected for this work because they turn out to have the lowest formation energies [45] and are, therefore, most likely to be found in actual catalysts. Only two of these structures are considered in this work because we show that their catalytic properties are immune to fluorination and this result is sufficient for our purpose.

**Figure 2.** DFT optimized configurations of metal-free catalysts investigated in this work. (**a**) Armchair configuration of N-doped carbon and (**b**) zigzag configuration of N-doped carbon. The meaning of the numbers on the carbon atoms is discussed in Section 2.3.

#### **2. Results**

#### *2.1. Fluorination of the FeNx Sites—Single Carbon Layer*

We first performed DFT optimizations of all the non-fluorinated carbon-based catalysts considered in this work. The resulting structures are shown in Figures 1 and 2. We then introduced F atoms and F<sup>2</sup> molecules at different locations on these atomic structures and performed new DFT optimizations.

We first considered the adsorption of F<sup>2</sup> on the Fe atom of the FeN<sup>x</sup> sites. We found that F<sup>2</sup> adsorbs on Fe for all the FeN<sup>x</sup> configurations of Figure 1. When adsorbed on Fe, the F<sup>2</sup> molecule is strongly stretched relative to the free F<sup>2</sup> molecule. In the example of Figure 3a, where we use the basic structure of Figure 1j, the spacing between the two F atoms of the adsorbed F<sup>2</sup> is 2.32 Å while the spacing of the two F atoms in the free F<sup>2</sup> molecule is 1.11 Å. In fact, the adsorbed F<sup>2</sup> molecule is subject to dissociation because, for example, the binding energy of the dissociated F<sup>2</sup> molecule with one F atom adsorbed on Fe and the other F atom adsorbed on a near C atom either on the same side (Figure 3b) or on the opposite side (Figure 3c) of the carbon layer is, respectively, 1.36 eV and 1.40 eV lower. The fluorinated structure is considerably more stable when two F atoms are adsorbed on the Fe atom on opposite sides of the carbon layer (Figure 3d) because the energy of the system is 6.12 eV lower than that of F<sup>2</sup> adsorbed on Fe (Figure 1a). The binding energies of the fluorine adsorbates shown in Figure 3 are given in Table 1. All F–Fe and F–C bond lengths appearing in Figure 3 are shown in Table A1 of the Appendix A.

In the case where two F atoms are adsorbed on Fe on both sides of the carbon layer (Figure 3d), the Fe site becomes unavailable to catalyze ORR. Indeed, the binding energy of the O<sup>2</sup> molecule on the Fe atom of the pristine structure of Figure 3 is only −0.71 eV, meaning that the fluorinated iron sites of Figure 3d,e are very stable and the adsorbed F is unlikely to be spontaneously replaced by O2. The same conclusion holds for all structures of Figure 1, except for Figure 1d, for which the binding energy of Fe–O<sup>2</sup> is stronger than that of Fe–F, as shown in Table 2. However, this structure appears to be inactive in the absence of fluorine, as can be seen in Figure 4d.

**Figure 3.** Adsorption of F<sup>2</sup> and F on the Fe site for the basic structure of Figure 1j. (**a**) Adsorption of F<sup>2</sup> ; (**b**,**c**) adsorption of one F on Fe and one F on a nearby carbon site; (**d**) adsorption of two F on the Fe site on both sides of the carbon layer; and (**e**) adsorption of a single F on Fe. Color code: grey is carbon, blue is nitrogen, orange is iron, and green is fluorine.


− − − − −

−

**Table 1.** Binding energies of the fluorine adsorbates for the structures shown in Figure 3.

**Table 2.** Binding energies in eV for F adsorbed on Fe (Fe–F), for two F adsorbed on Fe on both sides of the carbon layer (F–Fe–F), and for O<sup>2</sup> adsorbed on Fe (Fe–O<sup>2</sup> ) for the structures shown in Figure 1.


**Figure 4.** Relative free energy at zero potential for the six steps (I–VI) of the ORR sequence (1) for the atomic structures shown in Figure 1 with (red segments) and without (blue segments) a F atom adsorbed on Fe. The panels (**a**–**j**) correspond to the structures shown in Figure 1a to j.

For a fluorinated FeN<sup>x</sup> site located in a single carbon layer, the adsorption of a single F atom on the Fe site can turn a poor catalyst into an effective one due to the weakening of the binding energy of the adsorbates of the reaction sequence (1). This can be seen in Figure 4, which shows the relative free energies (with respect to the initial state) of each step of the catalytic reaction sequence (1) in the cases where the Fe sites are free of fluorine (blue segments) and where a F atom is adsorbed on Fe (red segments). One can see for the structures corresponding to Figure 4a,b,d,f, that the ORR catalytic process, which seems unlikely without the adsorption of F because of the very low energy level of step V, ∗OH + H2O + H<sup>+</sup> + *e* <sup>−</sup>), becomes thermodynamically viable with the adsorption of F on Fe. However, the structure corresponding to Figure 4i remains a poor catalyst with and without the adsorption of F on Fe. On the other hand, the structure corresponding to Figure 4h, which seems to be a possible good catalyst without fluorine, becomes less effective with fluorine because O<sup>2</sup> can no longer adsorb on Fe (step II). Table A2 shows the F–Fe, O–Fe–O, and O–O bond lengths in O2, Fe–O2, and F–Fe–O2. It can be seen that the F–Fe bond length in F–Fe–O<sup>2</sup> is almost the same as in Figure 3e (Table A1). However, the Fe– O and O–O bond lengths in F–Fe–O<sup>2</sup> are slightly smaller than in Fe–O2, indicating that the Fe–O<sup>2</sup> bond is weakened due to the presence of F, in agreement with the discussion above.

From this, we conclude that carbon-based catalysts with sufficient separation between carbon layers (which is equivalent to considering the catalytic sites located on a single carbon layer) could theoretically benefit from partial fluorination by the weakening of the adsorbate free energy. This effect leads to a more favorable free energy distribution for most of the fluorinated ORR active sites as compared to that of the sites without any F–Fe bond. The conclusion that ORR is generally promoted when an adsorbate is bound to the metal atom on the opposite side of the carbon plane was also obtained via the DFT calculations reported in [53,54].

#### *2.2. Fluorination of the FeNx Sites—Double Carbon Layer*

The useful FeN<sup>x</sup> active sites (those able to produce ORR) are actually thought to be mostly embedded at the surface of continuous graphene layer stacks or located between two discontinuous graphene layers of micro- or mesopores [20]. Therefore, the question naturally arises as to whether the ORR can take place between the first two carbon layers when a F atom is adsorbed on Fe on the more accessible free side of the FeN<sup>x</sup> site. To answer this question, we considered a model using two carbon layers: on the top, a carbon layer containing the FeN<sup>x</sup> moiety and, under this first layer, another one composed of a single graphene layer parallel to the first one and located at 3.6 Å from the first layer. We selected this interplanar distance, which is a little larger than that of d = 3.35 Å in graphite, because carbon is disorganized after the pyrolysis stage. As a matter of fact, a fairly broad distribution of d-spacings (between 3.5 and 4.1 Å) was measured for furnace turbostratic carbon black grades, regardless of their particle size and structure. The average TEM measured d-spacings range between 3.83 and 3.92 Å and are significantly larger than the X-ray measured d-spacings ranging from 3.52 to 3.56 Å [57]. Thus, we used the intermediate value of 3.6 Å.

The first step in ORR is the adsorption of O<sup>2</sup> on Fe. Then, according to the reaction sequence (1), each adsorbate combines with an H<sup>+</sup> ion and an electron. Therefore, O<sup>2</sup> and H<sup>+</sup> have to migrate between the carbon layers to reach the Fe atom of the FeN<sup>x</sup> site. This migration is certainly easier for porous carbon structures such as the one of Figure 1j or when the Fe atom is close to the edge of the carbon layer, as in Figure 1g,i. Figure 5 shows the basic structure of Figure 3e to which a parallel graphene plane was added under the graphene plane containing the F–FeN<sup>x</sup> site. One notes that there is apparently enough room between the two carbon layers to accommodate an oxygen atom or molecule because the spacing between the carbon layers of our amorphous carbon catalyst is assumed here to be 3.6 Å, while the theoretical radii of Fe, C, and O are 1.56 Å, 0.67 Å, and 0.48 Å [58], respectively, so that the sum of the radius of Fe, the diameter of O2, and the radius of C is 3.19 Å, which is smaller than the assumed spacing between the carbon layers.

**Figure 5.** The first catalytic reaction steps between two carbon layers with a F atom adsorbed on Fe on the free side for the basic structure of Figure 1j. (**a**) Adsorption of O<sup>2</sup> on Fe; (**b**) result of the spontaneous dissociation of OOH into O on Fe and OH on the opposite carbon layer; and (**c**) formation of OH adsorbed on Fe between the two layers. Color code: grey is carbon, blue is nitrogen, orange is iron, green is fluorine, and red is oxygen.

(Fe O) + (G OH) + 3൫ ା + ݁ି൯ III

(Fe OH) + G + H2O + ൫ ା + ݁ି൯ V

Fe + (G OH) + H2O + ൫ ା + ݁ି൯ V

 (Fe OH) + G + H2O + ൫ ା + ݁ି൯ V

 → (Fe Oଶ) + G + 4൫ ା + ݁ି൯ → II

> H ା + ݁ି

Path 1: (Fe O) + G + H2O + 2൫ ା + ݁ି൯ → IV

Path 2: (Fe OH) + (G OH) + 2൫ ା + ݁ି൯ → IV

Path 3: (Fe OH) + (G OH) + 2൫ ା + ݁ି൯ → IV

Fe + G + <sup>ଶ</sup> + 4൫ ା + ݁ି൯ I

Fe + G + 2H2O

H ା + ݁ି

The free energy diagram of the catalytic reaction is shown in Figure 6. O<sup>2</sup> can adsorb on the Fe atom between the planes (Figure 5a and step II in Figure 6) but it needs around 1 eV to get there. This is a consequence of the stress induced on the surrounding structure by the insertion of the O<sup>2</sup> molecule. Our transition state calculation indicates that the activation energy is around 1.5 eV between the state of O<sup>2</sup> in the pore and its adsorbed state on the Fe atom. However, from the latter adsorbed state (step II in Figure 6), a continuously decreasing free energy sequence can be found, but with modifications with respect to the sequence (1). When adding H<sup>+</sup> + *e* <sup>−</sup> to the adsorbed O2, OOH spontaneously dissociates into an O adsorbed on Fe and an OH which adsorbs on a carbon atom of the opposite layer (Figure 5b and step III in Figure 6). The reactions up to step III in Figure 6 are as follows:

$$\begin{array}{ccccc} \text{Fe} + \text{G} + \text{O}\_2 + 4\text{(H}^+ + \text{e}^-) & \rightarrow & (\text{Fe} - \text{O}\_2) + \text{G} + 4\text{(H}^+ + \text{e}^-) & \rightarrow & (\text{Fe} - \text{O}) + (\text{G} - \text{OH}) + 3\text{(H}^+ + \text{e}^-) \\\text{I} & \text{II} & \text{III} & \text{III} \end{array} \tag{2}$$

where G is the bottom graphene plane. Then three distinct paths are possible, depending on how the successive H<sup>+</sup> + *e* <sup>−</sup> are added to the adsorbates. These paths are illustrated in Figure 6 by black, blue, and red segments, respectively. The last step (VI) is Fe + G + 2H2O, i.e., the formation of two water molecules after the exchange of four electrons. The steps IV and V for the three paths are, respectively,

$$\begin{array}{ccccc}\text{Path 1:} & (\text{Fe}-\text{O}) + \text{G} + \text{H}\_{2}\text{O} + 2\text{(H}^{+} + \text{e}^{-}) & \rightarrow & (\text{Fe}-\text{OH}) + \text{G} + \text{H}\_{2}\text{O} + \text{(H}^{+} + \text{e}^{-})\\ & \text{IV} & \text{V} \end{array} \tag{3}$$

$$\begin{aligned} \text{Path 2: } \text{(Fe-OH)} + \text{(G-OH)} + 2\begin{pmatrix} \text{H}^{+} + \text{e}^{-} \end{pmatrix} & \rightarrow & \text{Fe} + \begin{pmatrix} \text{G} - \text{OH} \end{pmatrix} + \begin{pmatrix} \text{H}^{+} + \text{e}^{-} \end{pmatrix} \\ \text{IV} & \text{V} \end{aligned} \tag{4}$$

$$\begin{aligned} \text{Path 3: } \text{(Fe-OH)} + \text{(G-OH)} + 2\begin{pmatrix} \text{H}^{+} + \text{e}^{-} \end{pmatrix} &\rightarrow & \begin{pmatrix} \text{Fe}-\text{OH} \end{pmatrix} + \text{G} + \begin{pmatrix} \text{H}^{+} + \text{e}^{-} \end{pmatrix} \\ \text{IV} &\text{V} \end{aligned} \tag{5}$$

The most complex intermediate state, where one OH is adsorbed on Fe and the other OH is adsorbed on the opposite graphene layer G, is shown in Figure 5c and corresponds to step IV of paths 2 and 3.

**Figure 6.** Relative free energy at zero potential for the six steps of the ORR sequences between two carbon layers for the structure of Figure 3e to which a parallel graphene plane was added under the graphene plane containing the F–FeNx site. The three possible paths are explained in the text. The energy levels have been shifted slightly to facilitate path identification.

− − Since the theoretical radius of F (0.42 Å) is smaller than that of O (0.48 Å) [58], F and F<sup>2</sup> can also be accommodated between the two layers. We also performed DFT calculations for that case. The results are illustrated in Figure 7. The binding energies of F<sup>2</sup> and F in the cases of Figure 7a,b are −1.84 and −4.48 eV, respectively. Here the reference structure is composed of the two planes with the external adsorbed F atom on Fe. Thus, as in the

−

−

−

monolayer case of Figure 3, the dissociation of F<sup>2</sup> is favored between the two layers of carbon. In addition, contrary to O2, the adsorption of F<sup>2</sup> (or its dissociated form) between the plane is exothermic. For the sake of comparison with the single plane case of Figure 3d, the binding energy of the two F on both sides of the plane is −6.24 eV (Figure 7c) vs. −7.88 eV in the case of Figure 3d. Again, the difference in binding energy is due to the stress induced on the surrounding structure by the insertion of the F atom. The existence of F–Fe–F bonds in the catalyst could correspond to the specific peak at ~685.4 eV assigned to the adsorption of two F atoms on Fe in the F1s XPS spectrum of the catalyst [33]. −

**Figure 7.** Adsorption of F<sup>2</sup> and F on the Fe site between two carbon layers for the basic structure of Figure 1j. (**a**) Adsorption F<sup>2</sup> on Fe; (**b**) adsorption of F on Fe and F on a nearby carbon site (dissociated form of F<sup>2</sup> ); and (**c**) adsorption of F on Fe between two carbon layers.

From this study of the fluorinated bilayer configuration, it appears that ORR catalysis between two carbon layers is inefficient, primarily because O<sup>2</sup> requires about 1 eV to occupy the catalytic site, although the subsequent reaction steps are thermodynamically viable. Furthermore, upon fluorination of the catalyst, F<sup>2</sup> and F can occupy the catalytic site at a lower energy cost, thereby poisoning the FeN<sup>x</sup> sites.

#### *2.3. Fluorination of Metal-Free Sites*

− −

We now turn to the metal-free catalysts shown in Figure 2, which are also known to contribute to the ORR, but to a much lesser extent than the FeN<sup>x</sup> metal sites free of fluorine [59]. It has been demonstrated that these catalysts can produce uniformly descending free energy steps for the reaction sequence (1), as for some of the FeN<sup>x</sup> sites (see Figure 4). However, O<sup>2</sup> hardly adsorbs on these structures [45–49]. Our DFT calculations are in good agreement with those previous works, as shown by the blue segments in Figure 8a,b, which were obtained by considering the ORR active sites 5 and 1 in Figure 2a,b, respectively.

**Figure 8.** Relative free energy at zero potential for the six steps (I–VI) of the ORR sequence (1) for the sites shown in Figure 2 with (red segments) and without (blue segments) a F atom adsorbed on a carbon atom. (**a**) Armchair configuration of N-doped carbon; (**b**) zigzag configuration of N-doped carbon. The blue segments in (a) were obtained by considering the ORR activity of site 5 in Figure 2a. The blue segments in (b) were obtained by considering the ORR activity of site 1 in Figure 2b. The red segments in (a) were obtained by considering a F atom adsorbed on site 1 in Figure 2a, while keeping site 5 as ORR active. The red segments in (b) were obtained by considering a F atom adsorbed on site 1 in Figure 2b, while considering the ORR activity of site 6.

We then verified whether the carbon catalytic sites can be poisoned by F or F2. For the armchair configuration of Figure 2a, adsorption of F on the sites numbered from 1 to 5 were tested. Table 3 shows that F can adsorb on the five sites with different binding energies.

−

− − − − − − − − − −

−

Site 1 has the strongest binding energy of −2.57 eV. When a F atom is adsorbed on site 1, Figure 8a (red segments) shows that the catalytic site 5 remains active.

**Table 3.** Binding energies in eV for a F atom adsorbed on the 5 carbon sites of Figure 2a (armchair) and Figure 2b (zigzag).


For the zigzag structure we found that the catalytic site 1 in Figure 2b has the strongest binding energy of −3.78 eV for F. In this case, we carried out a calculation of the catalytic sequence on site 6 in Figure 2b. Even if the active site 1 is occupied by a fluorine atom, the ORR can take place on site 6, as can be seen in Figure 8b (red segments). Additional F–C bond length data for the five sites are given in Table A3. We note the anti-correlation between F–C bond length and binding energy as well as the smaller values compared to the F–Fe bond lengths given in Table A1.

We were unable to find a site where F<sup>2</sup> could be adsorbed on both the armchair and zigzag structures, so that dissociation of F<sup>2</sup> is unlikely on these structures. Therefore, F<sup>2</sup> has to be dissociated elsewhere, such as on the Fe sites, for instance, or at any oxygenated functionality like COH, COOH, or C–O–C, known (by XPS) to be present at the surface of our (and many other) non-PGM catalysts [33]. These calculations tend to confirm that fluorination does not affect much the catalytic activity of the metal-free catalysts considered here, thus providing an explanation for the residual ORR catalytic activity found after fluorination up to a value of F/C = 0.27 (measured by NMR) of the fluorinated Fe/N/C catalyst observed in the experiments reported in [33].

#### *2.4. Defluorination of the FeNx Sites*

Here we consider a question that has puzzled us for some time: the possibility to thermally de-fluorinate at 900 ◦C previously fluorinated FeN<sup>x</sup> catalytic sites such as the ones illustrated, for instance, in Figure 3d,e. Let us label these configurations F–FeN4–F and F–FeN4, respectively. From Table 1, the binding energies of the fluorine adsorbates on the Fe atom are −4.56 eV for F–FeN<sup>4</sup> and −7.88 eV for F–FeN4–F. Despite the high binding energies of these bonds, it was experimentally found that the latter are broken after a 30 min heat treatment at 900 ◦C in Ar of the fluorinated catalysts (no F1s XPS signal anymore, as seen in Figure 6A of [33]). How is this possible when the thermal energy at 900 ◦C is only around 0.1 eV?

It is certainly not because of the special nature of the catalytic sites such as those illustrated in Figure 3e,d, as Table 2 confirms that the FeN<sup>x</sup> sites illustrated in Figure 1 are all characterized by Fe–F and F–Fe–F binding energies of several eV. De-fluorination is only possible if the fluorinated FeN<sup>x</sup> sites are involved in reactions also involving either radicals or small molecules released from the catalyst surface under heat treatment. Reactions (6–9) below are examples of possible de-fluorination reactions. Our thermodynamic calculations show that all these reactions are characterized by a negative free energy change ∆G at 900 ◦C, meaning that they are spontaneous at that temperature.

$$(\text{F} - \text{FeN}\_4 - \text{F})\_{\text{solid}} + (\text{CH} \bullet)\_{\text{gas}} \leftrightarrow (\text{F} - \text{FeN}\_4)\_{\text{solid}} + \text{C}\_{\text{solid}} \text{ (graphene)} + (\text{HF})\_{\text{gas}} \tag{6}$$

$$(\text{F} - \text{FeN}\_4)\_{\text{solid}} + (\text{CH}\bullet)\_{\text{gas}} \leftrightarrow (\text{FeN}\_4)\_{\text{solid}} + \text{C}\_{\text{solid}} \text{ (graphene)} + (\text{HF})\_{\text{gas}} \tag{7}$$

$$(\text{F} - \text{FeN}\_4 \text{--F})\_{\text{solid}} + (\text{CF} \bullet)\_{\text{gas}} \leftrightarrow (\text{F} - \text{FeN}\_4)\_{\text{solid}} + (\text{CF}\_2 \bullet)\_{\text{gas}} \tag{8}$$

$$2(\text{F} - \text{FeN}\_4 \text{-F})\_{\text{solid}} + (\text{C}\_2\text{F}\_6)\_{\text{gas}} \leftrightarrow 2(\text{F} - \text{FeN}\_4)\_{\text{solid}} + 2(\text{CF}\_4)\_{\text{gas}} \tag{9}$$

In these examples, (CH•) gas, (CF•) gas, and (C<sup>2</sup> F<sup>6</sup> ) gas are decomposition products [60] generated at 900 ◦C from the carbonaceous or from the fluorinated carbonaceous

supports of the catalysts. The evidence for the release of such gases is documented by the TGA curves already reported in several figures of [33] for these fluorinated catalysts.

#### **3. Computational Methods**

All DFT calculations reported here were done using the Vienna ab initio software package (VASP) [61–64]. The calculations were performed using the generalized gradient approximation (GGA) with the Perdew–Burke–Ernzerhof (PBE) functional [65]. The convergence criterion on the relative energy was set to 10−<sup>5</sup> and the plane wave energy cut-off was set to 500 eV for all calculations. The Brillouin zone was sampled on regular 4 × 4 × 4 gamma grids. A graphene sheet with cell dimensions of a = 20.22 Å and b = 14.88 Å was used as a model for the carbon support. A void of 15 Å was included in the normal direction to avoid interactions between the periodic FeNx-doped carbon layers. The doped carbon structures were created by substituting carbon atoms of the graphene sheet by FeN<sup>x</sup> groups or by N atoms in the case of metal-free catalysts. The positions of all atoms were fully relaxed, except in the case of the two carbon layers, where the positions of the carbon atoms were fixed to prevent the planes from moving relative to each other. However, for the calculation of the activation energy of O<sup>2</sup> transiting between the two planes, in relation to Figure 5a, we used constraints where the edges of the planes were fixed along the *x*-axis while keeping the edges along the *y*-axis fixed, and vice versa. The activation energy was almost the same (1.5 eV) in both cases. The binding energy of an adsorbate on a given site was calculated using

$$E\_{\text{catalyst}+\text{adsorbat}} - \left(E\_{\text{catalyst}} + E\_{\text{adsorbat}}\right) \tag{10}$$

where *Ecatalyst*+*adsorbate* is the energy of the carbon-doped catalyst with the adsorbate, *Ecatalyst* is the energy of the catalyst alone, and *Eadsorbate* is the energy of the adsorbate far from the catalyst. The energy of H<sup>+</sup> + *e* <sup>−</sup> is taken as half the energy of the H<sup>2</sup> molecule, since H<sup>2</sup> is at an equilibrium with its dissociated form 2 H<sup>+</sup> + *e* <sup>−</sup>) at the anode [43]. The molecules of the gas phases considered in this work, namely O2, H2, and F2, are assumed to be non-interacting with each other, which implies that only single molecules have been considered. Each step of the catalytic sequence (1) corresponds to a free energy given by

$$\mathbf{G} = \mathbf{G}\_0 + ZPE + TdS + \mathbf{G}\_{\text{sol}} \tag{11}$$

where *G*<sup>0</sup> is the energy of the structure per cell, *ZPE* is the zero point energy, *TdS* is the entropy term, and *G*sol is the solvation energy arising by the aqueous medium. As was done in some of our previous works [52,54], for simplicity we assumed that the sum of the last three contributions nearly cancels, in agreement with [43,66]. However, corrections were brought to the intermediate state ∗O + H2O + 2 H<sup>+</sup> + *e* <sup>−</sup>) and ∗ + 2H2O of the reaction sequence (1), which were inferred to be +0.4 and −0.6 eV, respectively [43].

#### **4. Conclusions**

We used DFT to examine the consequences of fluorination of the FeNx-doped and N-doped carbon catalysts used for ORR at the cathode of H2/O<sup>2</sup> fuel cells. The main objectives of these calculations were to rationalize some of the experimental observations and to verify our conceptual representation of the catalytic sites and processes. We have considered several moieties of catalytic sites of FeNx-doped carbon with x ranging from 1 to 4. Most of them seem to be suitable catalysts for ORR because the free energy of the supposed catalytic sequence decreases regularly at zero potential. When the FeN<sup>x</sup> sites are located on a single graphene layer, it turns out that F<sup>2</sup> binds to Fe at FeN<sup>x</sup> sites, with a binding energy of approximately −2 eV, but is subject to dissociation, leaving a single F on Fe with a binding energy of approximately −4 eV, which is stronger than the typical binding energy of O<sup>2</sup> on Fe. In these conditions, ORR cannot happen on the F-poisoned FeN<sup>x</sup> side, but is still possible on the other side of the F–FeN<sup>x</sup> site, even transforming some otherwise poor un-poisoned FeN<sup>x</sup> catalytic configurations into better F–FeN<sup>x</sup> active

ones. In addition, two F atoms can also bind to Fe on both sides of the carbon layer with almost twice the binding energy of a single F. When this happens, the Fe site is completely poisoned on both sides and is no longer able to catalyze ORR.

The occurrence of single graphene layers in actual catalysts is probably quite exceptional. Those are certainly better represented by several stacks of disorganized graphene layers forming a network of connected micropores and mesopores. Therefore, we have also examined the double graphene layer case where there is a second parallel carbon layer at a distance of 3.6 Å from the upper carbon layer carrying the FeN<sup>x</sup> sites. We found that O<sup>2</sup> adsorption on Fe between the two carbon layers is stable and that OOH dissociates spontaneously into O adsorbed on Fe and OH adsorbed on the opposite carbon layer. Because O<sup>2</sup> adsorption increases the free energy by about 1 eV (and needs an activation energy of around 1.5 eV) relative to free O2, the catalytic process is unlikely in this case, even though the free energy of subsequent steps decreases monotonically. On the other hand, we found that F<sup>2</sup> can adsorb on Fe between the two carbon layers without energy expenditure, making this process more likely than for O2. These results suggest complete poisoning of the FeN<sup>x</sup> sites through extensive fluorination of the catalyst, in agreement with the experimental observations.

We then focused on the residual catalytic activity after fluorination by considering Fefree N-doped carbon armchair and zigzag structures for which previous DFT calculations suggested a viable catalytic process although O<sup>2</sup> hardly adsorbs on these structures. For both structures, the active catalytic site is a carbon atom near a N atom. We found that these catalytic structures are not poisoned by F or F2, thus justifying a residual ORR catalytic activity similar to that of the Fe-free catalysts observed for fluorinated Fe/N/C catalysts.

Finally, we provided an explanation for the recovery of ORR upon heating to 900 ◦C after fluorination. This explanation is based on the presence of radicals or small molecules released from the catalyst surface upon heat treatment. Most of the calculations presented in this work are based on free energy levels that only indicate whether a catalytic process is thermodynamically viable or not. A more thorough study would include the determination of activation energies. These calculations are very computationally demanding and will be the subject of future work.

**Author Contributions:** Conceptualization, M.C., J.-P.D. and F.V.; funding acquisition, S.S. and F.V.; investigation, M.C. and V.P.G.; methodology, M.C., J.-P.D. and F.V.; supervision, F.V.; writing—original draft, M.C.; writing—review and editing, M.C., J.-P.D., G.Z., V.P.G., S.S. and F.V. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Natural Sciences and Engineering Research Council of Canada under grant STP 521582-18.

**Data Availability Statement:** Additional data are available upon request from the corresponding author. The data set generated by the simulations performed in this work has not been made publicly available due to its large quantity and diversity.

**Conflicts of Interest:** The authors declare no conflict of interest.

**Sample Availability:** No Samples are available.

#### **Appendix A**

**Table A1.** Lengths in Å of the bonds involving the F atoms in Figure 3a–e.



**Table A2.** Bond lengths in Å for free O<sup>2</sup> and for adsorbates on the structure of Figure 1j.

**Table A3.** F–C bond lengths in Å for sites 1 to 5 shown in the structures of Figure 2.


#### **References**


## *Article* **A Nanosheet-Assembled SnO2-Integrated Anode**

**Xiaoli Wang <sup>1</sup> , Xinyu Zhao 2, \* and Yin Wang 2, \***


**Abstract:** There is an ever-increasing trend toward bendable and high-energy-density electrochemical storage devices with high strength to fulfil the rapid development of flexible electronics, but they remain a great challenge to be realised by the traditional slurry-casting fabrication processes. To overcome these issues, herein, a facile strategy was proposed to design integrating an electrode with flexible, high capacity, and high tensile strength nanosheets with interconnected copper microfibre as a collector, loaded with a novel hierarchical SnO<sup>2</sup> nanoarchitecture, which were assembled into core–shell architecture, with a 1D micro-fibre core and 2D nanosheets shell. When applied as anode materials for LIBs, the resultant novel electrode delivers a large reversible specific capacity of 637.2 mAh g <sup>−</sup><sup>1</sup> at a high rate of 1C. Such superior capacity may benefit from rational design based on structural engineering to boost synergistic effects of the integrated electrode. The outer shell with the ultrathin 2D nanoarchitecture blocks can provide favourable Li + lateral intercalation lengths and more beneficial transport routes for electrolyte ions, with sufficient void space among the nanosheets to buffer the volume expansion. Furthermore, the interconnected 1D micro-fibre core with outstanding metallic conductivity can offer an efficient electron transport pathway along axial orientation to shorten electron transport. More importantly, the metal's remarkable flexibility and high tensile strength provide the hybrid integrated electrode with strong bending and stretchability relative to sintered carbon or graphene hosts. The presented strategy demonstrates that this rational nanoarchitecture design based on integrated engineering is an effective route to maintain the structural stability of electrodes in flexible LIBs.

**Keywords:** anode; flexible electronics; nanosheets; SnO<sup>2</sup>

#### **1. Introduction**

Nowadays, an urgent and key task for energy conversion storage systems, in particular lithium-ion batteries (LIBs), is to develop advanced electrode materials with mechanical durability and superior Li-storage performance for booming flexible energy storage applications in foldable smartphones, wearable electronic systems, and implantable device [1–5]. However, the design and fabrication of such a high-performance flexible electrode is still a major challenge via a facile method because of the lack of optimal materials with the feature of high special capacity and robust mechanical flexibility in electrochemical environments. The traditional slurry-coating fabrication technology is not suitable for flexible LIBs because the active materials often suffer from exfoliation or cracking in the process of frequent bending. Furthermore, these additional binders in slurry would hinder electronic transport and reduce the specific energy density of the battery, and conductive carbon black should be added into the slurry to improve electrical performance.

In order to keep pace with the development of flexible energy storage systems and solve these problems of excessive consumption of adhesives and carbon black (~30 wt%), one of the most effective strategies is to construct an integrated electrode instead of slurrycoating fabrication technology [6–11]. In this regard, searching for an appropriate flexible

**Citation:** Wang, X.; Zhao, X.; Wang, Y. A Nanosheet-Assembled SnO2-Integrated Anode. *Molecules* **2021**, *26*, 6108. https://doi.org/ 10.3390/molecules26206108

Academic Editors: Gregorio F. Ortiz and Munkhbayar Batmunkh

Received: 1 August 2021 Accepted: 7 October 2021 Published: 10 October 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

current collector and high-energy-density active material is crucial to achieving these goals. To date, noteworthy progress in novel flexible electrodes has been achieved through chemical synthesis routes, including hydrothermal [12], sol–gel techniques [13], and CVD [14]. Nevertheless, most reported flexible energy storage devices are performed by using biodriven carbon or free-standing film, which are not strong enough to withstand frequent mechanical deformations during practical use. Hence, an integrated electrode with robust tensile strength and enhanced electrochemical performance is still plagued and needs to be further improved.

Moreover, group IV element Sn in the form of metal oxide (SnO2) has particularly attracted extensive attention as a promising anode material to replace conventional graphitic carbon in current LIBs because of its uniqueness in terms of low cost, safe working potential, high theoretical capacity, and environmental friendliness [15–19]. Nevertheless, the simple structure, relatively low intrinsic conductivity, and vast structural variation during the reversible insertion/deinsertion processes of the bulk SnO<sup>2</sup> powder keep it from achieving its full capacitance potential. One of the effective strategies to mitigate these problems is to develop structural engineering, including morphology control and hybrid construction, which could alleviate the mechanical stress induced by large volume change and prevent aggregation of the active domains [20–25]. However, these materials still need to be mixed with a binder and carbon black and pressed onto metal substrates or, alternatively, by being deposited onto a conductive substrate before they are assembled into batteries, which makes them less flexible and have a low energy density. Integrated electrodes, in which electrochemically active nanostructures are conformably coated on conductive collectors, have been demonstrated with ultrafast power rate and long lifespan. The successful integration of the sturdy conductive matrix support with elegant nanostructures improves the electrode performance and endows it with robust mechanical flexibility. All these merits render this type of electrode very attractive for flexible power sources.

To combine the aforementioned merits and promote the development of flexible LIBs, herein, we devised a novel route to fabricate SnO2-integrated electrode assembled by core–shell architecture, with a 1D micro-fibre core and 2D nanosheets shell, targeting high capacity and tensile strength LIBs. The copper micro-fibre clothes function as a superior conductive pathway facilitating fast electrons transfer along the axial orientation but also provide a high mechanical substrate assisting independent growth of SnO<sup>2</sup> nanosheets. The nanosheets shells are vertically distributed on copper micro-fibre, forming core–shell structure, which can provide a large electrode/electrolyte interfacial contact area. As expected, when measured as anode materials of LIBs, we obtained a reversible lithium storage capacity of 637.2 mAh g−<sup>1</sup> at a current density of 1C. The presented synthetic strategy is effective and with low cost, which provides a novel route to design advanced electrode materials for flexible LIBs.

#### **2. Materials and Methods**

#### *2.1. Sample Synthesis*

Prior to the synthesis, a piece of copper micro-fibre (CMF) textile (approximately 5 × 5 cm<sup>2</sup> ) (Liaoning Copper Group, Liaoyang, China) was treated by ultrasonication with 1 M hydrochloric acid solution (Aladdin, Shanghai, China), in order to remove the CuO layer. In a typical hydrothermal synthesis of the SnO<sup>2</sup> nanosheets array, 12 mmol SnCl2•2H2O and 24 mmol NH4F were first mixed in 70 mL deionised water under magnetic stirring (Aladdin, China). The clear and transparent solution was continuously stirred for 0.5 h in the air. Then, the mixed solution was transferred into a 100 mL Teflon-lined stainless steel autoclave. Afterwards, the CMF substrate was immersed in the solution. The sealed autoclave was heated to 180 ◦C for 24 h. Subsequently, the grey cloth was cleaned repeatedly with deionised water and ethanol under ultrasonic treatment and then dried under N<sup>2</sup> gas flow. To investigate the formation of the nanosheet SnO2, a series of parallel experiments were carried out by adjusting the molar ratio of fluoride/Sn.

#### *2.2. Materials Characterisation*

The crystallographic information of the as-prepared materials was recorded with powder X-ray diffractometry (XRD, Bruker D8 Advance, Bremen, Germany). The morphological features and microstructure of the sample were observed by using a field-emission scanning electron microscopy (FE-SEM, S-4800, Hitachi, Tokyo, Japan) and transmission electron microscopy (TEM, JEOL JEM-2100F, Japan). X-ray photoelectron spectroscopic (XPS, ESCALAB 250, Thermo Scientific, Waltham, MA, USA) was used to characterise the surface composition of the sample.

#### *2.3. Electrochemical Measurements*

Electrochemical properties were measured using coin cells (CR2025), which were assembled in an argon-filled glovebox (Mbraun, Unilab, Germany). The copper micro-fibretextile-supported ultrathin SnO<sup>2</sup> nanosheets were used directly as the working electrode for the subsequent electrochemical tests without binders and conductivity carbon black. The coated cloth was cut into disk electrodes (12 mm in diameter). To fabricate the SnO<sup>2</sup> power working electrode, the active materials, super-P, and PVDF with a mass ratio of 80:10:10 were ground in NMP solvent to form a homogeneous slurry, which was then coated onto the Cu foil by the doctor blade method and dried by heating in a vacuum oven (Yiheng, Shanghai, China). The lithium foil was used as the cathode electrode. The commercial electrolyte in the present measurements was a mixture of LiPF<sup>6</sup> in ethylene carbonate, dimethyl carbonate, and diethyl carbonate (EC-DEC-EMC, 1:1:1 in *v*/*v*). Galvanostatic cycling performances of the as-prepared coin cell were operated at room temperature on a multi-channel battery testing system (Land CT2001A, Wuhan, China) with a cut-off voltage of 1.2–0.01 V versus Li/Li + . Cyclic voltammetry (CV) curves were carried out by applying a CHI-760E electrochemical workstation at a scanning rate of 0.1 mV s −1 . – −

#### **3. Results and Discussion**

The flexible integrated electrode samples were fabricated via a facile in situ hydrothermal depositions of active material on the CMF without any post-treatment and use of conventional carbon black additive and binder. The overall fabrication process of the SnO2 integrated electrode is shown schematically in Figure 1. As detailed in the Experimental Section, the acid-treated copper micro-fibre textile was immersed in a mixed solution of stannous chloride dihydrate and ammonium fluoride, followed by a hydrothermal reaction at 180 ◦C for 24 h, yielding the large-size free-standing SnO<sup>2</sup> nanosheets on the CMF. The SnO2-integrated electrode materials display superior flexibility and robust mechanical tensile, which can be cut directly into electrode pieces with a diameter of 12 mm.

**Figure 1.** A schematic diagram for the fabrication procedure of SnO<sup>2</sup> nanosheets on the CMF.

Firstly, the powder X-ray diffraction (XRD) technique was provided to determine the composition characterisation of the as-prepared SnO<sup>2</sup> samples. Figure 2a displays the crystallinity and phase purity of the SnO<sup>2</sup> nanosheets assembled integration materials. The main three peaks at around 43.3, 50.4, and 74.1 ◦ correspond to the metal copper substrate. Removing the Cu element signals from the substrates, all the diffraction peaks can be well

indexed to the tetragonal rutile structure of SnO<sup>2</sup> (JCPDS Card no.41-1445). No impurities, such as SnO or Sn, were detected, indicating the formation of pure SnO<sup>2</sup> nanosheets. In addition, a piece of SnO2/CMF-integrated composite sample was put into a concentrated nitric acid solution to remove the Cu substrates, and the resulting powers were purified with deionised water and ethanol for subsequent testing. As shown in Figure 2b, the XRD pattern indicated that the as-prepared powers after concentrated nitric acid treatment were in a pure tetragonal rutile crystalline phase. Supplementary Figure S1 presents the XRD pattern of the SnO<sup>2</sup> nanoflower assembled nanosheets synthesised by hydrothermal method without CMF, which is in accordance with the above-mentioned sample.

**Figure 2.** XRD pattern of SnO<sup>2</sup> : (**a**) nanosheets on the CMF; (**b**) nanosheets after HNO<sup>3</sup> treatment.

To further verify the near-surface chemical composition and the oxidation state of element Sn in the as-prepared SnO<sup>2</sup> products, XPS measurement was then performed, as shown in Figure 3. A survey XPS spectrum of SnO<sup>2</sup> nanosheet is clearly observed in Figure 3a, indicating the existence of Sn and O elements. For the high resolution of element Sn, two strong peaks centred at 486.8 and 495.2 eV in the XPS spectrum could match with Sn 3d5/2 and 3d3/2 states, respectively. This clearly indicated that tin was in the Sn (IV) state in the nanosheets sample, which is in good accordance with previously reported SnO<sup>2</sup> [26].

**Figure 3.** XPS spectrum of SnO<sup>2</sup> nanosheets: (**a**) survey XPS pattern; (**b**) high-resolution XPS of Sn 3d.

FE-SEM and TEM images with different magnifications could provide information about the surface morphology and crystallographic properties of the as-received products. Figure 4a shows a typical SEM image of the uncovered SnO<sup>2</sup> nanosheets micro-fibre cloth composed of perpendicular and smooth copper micro-fibre, which was interconnected into a textile structure (20 × 20 µm<sup>2</sup> squared pore) and served as the backbone for the growth of SnO<sup>2</sup> nanosheets. As shown from Figure 4b, the hydrothermal treatment resulted in a significant morphology change of the copper micro-fibre cloth from a relatively smooth surface to a very rough surface, which indicated that the conductive substrate had been

covered with the numerous SnO<sup>2</sup> nanosheets and formed a CMF@SnO<sup>2</sup> hybrid material. Based on recorded top-view high magnification SEM images (Figure 4c–e), the outer sheath in the presented 1D architecture consisted of a uniform sheet structure with a smooth surface. Apparent open space between adjacent SnO<sup>2</sup> nanosheets presented sponge-like porous architecture from the magnified SEM images (Figure 4d,e), which provided a large interfacial area for electrolyte ion diffusion and ensured a short solid-state diffusion length for fast Li-ion insertion/extraction. The direct growth of SnO<sup>2</sup> nanosheets on the copper micro-fibre network collector enabled good contact and strong binding between SnO<sup>2</sup> and copper micro-fibre without any binder and carbon black. Due to the high surface energy effect of nanostructured materials, a large number of nanosheets were attached to the outer of the integrated electrode (Supplementary Figure S2), which maybe bring more negative effects on the subsequent electrochemical testing. Hence, the SnO2-integrated electrode was purified under ultrasonic treatment. As a result, SnO<sup>2</sup> nanosheets in situ grown on the CMF were not peeled off even after repeated bending or ultrasonic treatment due to the present interconnected nanoarchitecture. Figure 4f shows the photograph of a copper textile coated with a layer of grey SnO<sup>2</sup> nanosheets. Notably, similar to micro-fibre cloth, the SnO2-integrated electrode with the nanosheets coating could be easily bent without damage to the nanosheets, making them interesting for flexible batteries. –

20 μm

**Figure 4.** SEM images of (**a**) pristine copper microfibre cloth, where sub-millimetre pores can be clearly observed; (**b**) and (**c**) low magnification images of the copper textile surfaces after SnO<sup>2</sup> nanosheets layers are grown; (**d**) and (**e**) highmagnification SnO<sup>2</sup> nanosheets; (**f**) shows a photo of a piece of folded copper textile coated SnO<sup>2</sup> nanosheets.

Further microstructure information about SnO<sup>2</sup> nanosheets building blocks was obtained from transmission electron microscopy (TEM) (Figure 5). As shown in Figure 5a–d, the nanosheets structure could be clearly seen at low magnification. TEM further revealed that the hierarchical SnO<sup>2</sup> nanoarchitecture was built up of highly porous interconnected nanosheets. The curled geometrical morphology further showed that the presented SnO<sup>2</sup> electrode materials exhibited significantly improved microstructural flexible performance through nanostructured engineering. The average lateral size and thickness of the nanosheets were found to be approximately 300 and 20 nm, respectively. The morphology and size of the sample obtained from TEM were in good agreement with those observed in the SEM images. Furthermore, Figure 5d shows the HR-TEM image and measured lattice

–

fringes as 3.35 Å, which is consistent with the (111) interplanar distance of the SnO<sup>2</sup> phase. The inset of Figure 5d is the SAED pattern of a piece of SnO<sup>2</sup> nanosheets film, which clearly demonstrates the polycrystalline nature of the nanosheets films.

– **Figure 5.** (**a**–**c**) Typical TEM images of the nanosheets SnO<sup>2</sup> ; (**d**) HRTEM images of the nanosheets SnO<sup>2</sup> . The inset of d shows that these nanosheets are polycrystalline.

To gain insight into the effect of ammonium fluorides (NH4F) on the morphology of SnO<sup>2</sup> nanostructure, fluoride-dependent experiments were carried out. As shown in Supplementary Figure S3a, SnO<sup>2</sup> nanoparticles with an irregular shape and size of several nanometres were obtained in the absence of NH4F, which confirmed that the NH4F was key to the formation of the 2D SnO<sup>2</sup> nanoarchitecture. When the F/Sn molar ratio was 1, hollow-nanosphere-shaped samples with coarse surfaces could be detected (Supplementary Figure S3b). Once the F/Sn molar ratio was increased to 2, the well-defined 3D flower-like SnO<sup>2</sup> hollow nanoarchitecture (or nanoflower) constructed by many 2D nanosheets were generated, and their detailed characteristics are described in Supplementary Figure S3c,d. However, the SnO<sup>2</sup> products could not be found in the Teflon-lined stainless-steel autoclave if the molar ratio was increased to 4.

To prove the efficacy of the as-synthesis unique CMF@SnO<sup>2</sup> nanosheets architecture as a potential anode material suitable for practical application, we next performed an electrochemical evaluation of this integration electrode without adding any binders (PVDF) and conductive black additives (CB) in half-cells. It is well known that the electrochemical reaction process in the anode can be divided into two steps, as follows:

$$\text{SnO}\_2 + 4\text{Li}^+ + 4\text{e}^- \rightarrow \text{Sn} + 2\text{Li}\_2\text{O} \tag{1}$$

$$\text{Sn} + \text{xLi}^+ + \text{xe}^- \leftrightarrow \text{Li}\_\text{x}\text{Sn} \ (0 \le x \le 4.4) \tag{2}$$

$$\text{Li}^+ + \text{e}^- + \text{electrolyte} \rightarrow \text{SEI layer} \tag{3}$$

To gain better insight into the electrochemical performance of the novel integration anode, cyclic voltammetry (CV) characteristics of the initial four cyclic were first performed between 0.01 and 3 V, at a scan rate of 0.1 mV s −1 . As can be seen from Figure 6a, there was a substantial difference between the first and the subsequent cycles. In the first cathodic process, there were significant reduction peaks at 0.8 to 1.0 V, which is related to the conversion of tin dioxide to metallic tin and lithium oxide, and the formation of SEI layer on the electrode surface, as illustrated in Equations (1) and (3) [23,24]. Note that there was a strong reduction peak located at around 0.05 V, which is assigned to the formation of the LixSn alloys process (Equation (2)). In the subsequent anodic scan, a strong peak at 0.6 V corresponded to the phase transition from LixSn alloy to metallic Sn, which increased in intensity as the cycle number was increased. This phenomenon could be explained by the activation of the reversible reaction that occurred in the electrode materials [27]. Beyond that, oxidation peaks around 1.3 and 1.9 V were also observed for the SnO2-integrated anode materials [23,24,28]. This observation might suggest partial reversibility of the reaction by Equation (1), which is also observed for the nano-sized SnO<sup>2</sup> and SnO2/carbon composites [29,30]. In the following process of CV testing, the peak potentials of anodic and cathodic curves were almost coincident, and the peak intensity changed very slightly, indicating robust cycling stability of SnO2-integrated anode. − − metal's remarkable

−

−

– – **Figure 6.** (**a**) Cyclic voltammograms (CVs) of nanosheet-assembled SnO<sup>2</sup> -integrated electrode; (**b**) discharge–charge voltage profile of nanosheet-assembled SnO<sup>2</sup> -integrated electrode between 0.01 and 1.2 V at a rate of 1 C; (**c**) cycling performance and coulombic efficiency (CE) of nanostructured SnO<sup>2</sup> electrode; (**d**) cycling performance at different rates (1–4 C) of nanosheet-assembled SnO<sup>2</sup> -integrated electrode.

Furthermore, the galvanostatic charge and discharge measurements of the nanosheetassembled SnO2-integrated electrode were carried out at a high rate of 1 C for 1st, 2nd, 3rd, 50th, and 70th cycles in order to study Li-storage performance. It can be clearly seen that the nanosheet-assembled SnO2-integrated electrode showed a similar curve to those of SnO<sup>2</sup> anode materials and delivered a large initial discharge and charge capacity of 1518.1 and 818.6 mAh g −1 , respectively. The capacity loss, compared with the first

cycle, may be mainly attributed to irreversible side reactions, such as the trapping of some lithium in the lattice, formation of solid–electrolyte interface (SEI) layer, and electrolyte decomposition [23,24,28–30]. Nonetheless, perfect reversibility of the capacity was still obtained, and the charge and discharge capacities gradually stabilised in the following 70 cycles.

Cycle stability is another important parameter of the nanosheet-assembled SnO2 integrated electrode. Figure 6c displays the cycling performance of the nanosheet-assembled SnO2-integrated electrode with a voltage window of 0.01–1.2 V at a current rate of 1 C. It can be seen that the cycling of the nanosheet-assembled SnO2-integrated electrode was quite stable, which delivered a high reversible capacity of 637.2 mAh g−<sup>1</sup> after 70 cycles. It is should be noted that an average Coulombic efficiency of higher than 99% could be obtained up to 70 cycles after the second cycle. In Table S1, we compared the specific capacity of the proposed SnO2-integrated electrode with the relevant free-standing anode, including SnO2/CNTs, SnO2/graphene, and SnO<sup>2</sup> nanoarchitecture on the metal conductive substrate. Given these results, it is expected that synthesising SnO2-integrated anode assembled by nanosheets could promote the electrochemical performance of SnO2. Moreover, we also performed the as-synthesised nanoflower, nanosphere, and nanoparticle test as an anode material by the traditional slurry-casting fabrication processes. In comparison, the charge capacity of nanoflower remains 438.5 mAh g−<sup>1</sup> , outperforming that of the theoretical capacity of graphite (372 mAh g−<sup>1</sup> ). The nanosphere assembled by nanoparticles had an initial capacity of 1383.3 mAh g-1, quickly decreasing to 287.5 mAh g−<sup>1</sup> after 70 cycles. The nanoparticles without fluorine-doped SnO<sup>2</sup> showed a relatively low capacity of 123.08 mAh g−<sup>1</sup> after 70 cycles; in particular, the first 10 cyclings deteriorated sharply with the increasing cycle. It is suggested that such nanosheet-assembled SnO2-integrated electrodes could provide more interconnection between the building blocks and a more stable porous structure due to the effective prevention of dense aggregation of the nanosheets. Such a superior Li-storage performance of nanosheet-assembled SnO2-integrated electrode should be ascribed to the reasonable structural design on structural engineering to boost synergistic effects. The ultrathin 2D nanoarchitecture blocks provide favourable Li<sup>+</sup> lateral intercalation lengths and more beneficial transport routes for electrolyte ions, and the interconnected 1D micro-fibre core with outstanding metallic conductivity can offer an efficient electron transport pathway along axial orientation to shorten electron transport. More importantly, the metal's remarkable flexibility and high tensile strength endow the hybrid integrated electrode with strong bending and stretchability relative to sintered carbon or graphene hosts. Further, for the electrode realised by the traditional slurry-casting fabrication processes, the better cyclability of nanoflower and nanosphere than that of nanoparticle also supports the key role of fluorine-doped strategy [31,32].

To further evaluate the improved electrochemical performance of the SnO2-integrated anode, the rate capability at diverse larger current densities was also investigated, which is important for practical applications of LIBs. Benefiting from its unique structure, the nanosheet-assembled SnO2-integrated electrode revealed an exceptional cycling response to a continuously varying current rate. As displayed in Figure 6d, the representative specific capacities were about 731.17, 677.55, 567.30, and 423.71 mAh g−<sup>1</sup> at current rates of 1 C, 2 C, 3 C, and 4 C, respectively. The specific capacity slightly decreased as the current density increased, and it could still be maintained a very stable cycling capacity of above 423.71 mAh g−<sup>1</sup> , corresponding to nearly 100% Coulombic efficiency when the current density was up to 4 C, which was still higher than the theoretical capacity of graphite (372 mAh g−<sup>1</sup> ). More importantly, when the current density was adjusted back to the original current density again, the specific capacity of the SnO2-integrated electrode still regained the initial reversible value after the high-rate test for the 40 cycles, implying superior stability of the present SnO2-integrated electrode.

Overall, using microscopic 2D self-assembled materials combined with a macroscopic 2D metal conductor to prepare integrated electrodes is one of the most promising strategies to increase ion and electron transport kinetics toward present LIBs. According to the above results, the presented SnO2-integrated anode showed significantly improved capacity and rate performance. Clearly, the excellent improvement of electrochemical performance of the nanosheets SnO<sup>2</sup> arrays on the copper micro-fibre electrode can be attributed to two reasons. Firstly, self-supported nanosheets arrays growing directly on a current-collecting substrate represents an attractive nanoarchitecture for LIBs. Such structural feature of thin sheets combined with enriched pores built from the stacking of nanosheets is beneficial to rapid Li<sup>+</sup> intercalation and diffusion of electrolyte into the inner region of the electrode, high electrode–electrolyte contact area, and good stain accommodation. Moreover, each nanosheet has its own contact with the substrate at the bottom, which can ensure every nanosheet participates in the electrochemical reaction and effectively prevent the aggregation of the SnO<sup>2</sup> nanosheet. Secondly, the copper grid was selected as an effective substrate because of its high conductivity, 2D planar structure, and larger mechanical strength, compared with carbon materials. Furthermore, the proposed technique also saves the tedious process of mixing active materials with ancillary materials such as carbon black and polymer binders.

#### **4. Conclusions**

In this study, we presented a cost-effective, scalable, and effective approach to fabricate SnO<sup>2</sup> nanosheets cluster arrays directly grown on a 2D-interconnected conductive network with robust mechanical flexibility via a facile hydrothermal route. Such integrated electrodes possess a network configuration, which offers more beneficial transport routes for electrolyte ions and guarantees an intimate contact between active and current collectors. As a result, the SnO2-integrated electrode with novel nanoarchitecture shows an excellent electrochemical Li-storage performance with a high capacity up to 637.2 mAh g−<sup>1</sup> at 1 C rate and excellent rate capability of 423.71 mAh g−<sup>1</sup> at 4 C rate. Such superior cyclic stability and capacity may benefit from the well-designed electrode to boost synergistic effects, which include shortened Li<sup>+</sup> diffusion distance in the 2D nanoarchitecture blocks, sufficient void space among the nanosheets to reduce volume expansion, and a substrate with superior flexibility and robust tensile strength. The presented strategy provides a new synthetic idea for engineering tin-based energy storage systems with high electrochemical performance and robust mechanical flexibility.

**Supplementary Materials:** The following are available online, Figure S1: XRD pattern of nanoflower SnO<sup>2</sup> assembled by nanosheets without CMF, Figure S2: SEM image of the integrated electrode before ultrasonic treatment, Figure S3: SEM image of SnO<sup>2</sup> samples at 180 ◦C for 24 h: (a) without NH4F; (b) F/Sn = 1; (c,d) F/Sn = 2, Table S1: Comparison of electrochemical properties of nanosheetassembled SnO<sup>2</sup> -integrated electrode with other reported Sn-based free-standing anode materials.

**Author Contributions:** Conceptualisation, X.W. and X.Z.; methodology, X.Z. and Y.W.; investigation, X.W. and X.Z.; writing—original draft preparation, X.Z. and Y.W.; writing—review and editing, X.W. and X.Z. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Natural Science Foundation of Liaoning Province (Project No: 20180550749), the Foundation of Liaoning Key Laboratory of Chemical Additive Synthesis and Separation (Project No: ZJNK2008), the Doctoral Scientific Research Foundation of Inner Mongolia University for Nationalities (Project No: BS422), and the Opening Foundation from Inner Mongolia Key Lab of Carbon Nanomaterials (Project No: MDK2018050).

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

**Acknowledgments:** X.Z. deeply thanks the West Light Foundation's Visiting Scholar Research Program.

**Conflicts of Interest:** The authors declare no conflict of interest.

**Sample Availability:** Samples of the compounds are available from the authors.

#### **References**

