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Article

Influence of the Thermal Treatment on the Structure and Cycle Life of Copper Hexacyanoferrate for Aqueous Zinc-Ion Batteries

by
Mohsen Baghodrat
1,
Giorgia Zampardi
1,*,
Jens Glenneberg
2 and
Fabio La Mantia
1,2,*
1
Energiespeicher- und Energiewandlersysteme, Universität Bremen, Bibliothekstraße 1, 28359 Bremen, Germany
2
Fraunhofer Institute for Manufacturing Technology and Advanced Materials—IFAM, Wiener Straße 12, 28359 Bremen, Germany
*
Authors to whom correspondence should be addressed.
Batteries 2023, 9(3), 170; https://doi.org/10.3390/batteries9030170
Submission received: 1 February 2023 / Revised: 1 March 2023 / Accepted: 4 March 2023 / Published: 15 March 2023
(This article belongs to the Section Battery Materials and Interfaces: Anode, Cathode, Separators and Electrolytes or Others)
Browse Figures
Versions Notes

Abstract

:
Copper hexacyanoferrate (CuHCF) has become an attractive Zn2+ insertion material as a positive electrode in aqueous zinc-ion batteries thanks to its high reversibility towards Zn2+ (de-)insertion, its simple, inexpensive and easily scalable synthesis route, its low toxicity, and its high working potential. It is known that the physiochemical properties of CuHCF can be modified by manipulating its synthesis parameters. However, the effect of these parameters on the material’s electrochemical performance and cycle life needs further investigation. Here, the structure and composition of CuHCF treated at different temperatures are studied through crystallographic, compositional, and thermogravimetric analyses. The resulting CuHCF powders were galvanostatically cycled to assess their electrochemical performance in relation to their annealing temperature. The results showed that the annealed CuHCF electrodes exhibited longer cycle life while maintaining a coulombic efficiency ≥ 99.5%. The longest cycle life was achieved by annealing the CuHCF electrodes at 100 °C.

Graphical Abstract

1. Introduction

To meet the increasing energy demands of current society, together with the need for independence from fossil fuels, it has become of critical importance to develop sustainable and low-cost methods to store electrical energy harvested from renewable sources [1,2,3]. Developing bulk energy-storage systems that meet the strict standards required by the stationary market in terms of high-rate capability, high safety, and low environmental impact is therefore essential for the cost-effective and sustainable integration of renewable energy into the power grid [4,5,6,7,8]. Aqueous metal-ion batteries are ideal candidates for such applications, as they are intrinsically cheap and safe energy-storage devices. However, they require further advancements to reach higher efficiency and longer cycle life, in order to compete with the more mature organic-based Li-ion technology [9,10,11,12].
In the current research on aqueous metal-ion batteries, there is a rising interest in developing aqueous Zn-ion batteries (A-ZIBs). The high specific power and high reversibility in aqueous solutions of these batteries, combined with their low cost, environmental friendliness, and earth-abundance of metallic zinc, have made them appealing choices for large-scale grid-energy-storage applications [5,6,13,14,15,16,17].
Despite the promising advantages of A-ZIBs, their commercial application remains limited due to the low efficiency of the zinc electrodeposition reaction occurring at the negative electrode [16,18] and the lack of suitable Zn-insertion materials for the positive electrode. To resolve the latter problem, it is crucial to consider different aspects, such as economics, safety, and ecology. Ideally, the cathode material needs to be non-toxic and cost-effective, and, at the same time, highly durable when cycled in aqueous electrolytes without compromising its reversibility [4,9,19,20].
Among the insertion materials for the positive electrode of an A-ZIB, the hexacyanometallates of transition metals, generally known as Prussian blue analogues (PBAs), have generated significant interest due to their high reversibility towards the (de-)insertion of many cations, low volume changes during the ion-insertion reaction, low toxicity, and low costs [4,21,22,23,24]. As shown in Figure 1, the compounds belonging to the PBA family are characterized by an open-framework structure, with large cavities and channels created by two transition metals that are octahedrally linked together by CN ligands [25,26,27]. The face-centered cubic crystal lattices of PBAs allows the rapid movement of ions, enabling high-rate performance due to the fast (de-)insertion of a wide variety of ions, such as Li+, Na+, K+, Zn2+, Mg2+, and Al3+ [4,20,28,29,30,31,32,33].
Among the materials belonging to the PBA family, copper hexacyanoferrate (CuHCF) has attracted significant attention as an active material for the positive electrode in aqueous Zn-ion batteries since the first time it was used for this purpose, in 2015 [31].
Copper hexacyanoferrate shows high reversibility towards the (de-)insertion reaction of various monovalent, divalent, and trivalent ions in aqueous electrolytes [6,19,34]. It can be synthesized using abundant and non-toxic elements through a simple and inexpensive synthesis route that can be easily scaled up to an industrial level. Moreover, this material is well suited to power-grid applications because of its excellent power-rate capability and high cell-working potential, of about 1.7 V, compared with Zn2+/Zn. The latter enables CuHCF-based A-ZIBs to utilize the electrochemical-stability window of water almost fully [30,34,35,36,37].
Despite its many advantages, CuHCF suffers from a relatively short cycle life when operated in rechargeable aqueous-based systems, compared to the commercial insertion materials employed in organic lithium-ion batteries. Interestingly, it has been observed that the stability and the electrochemical properties of CuHCF can be tuned by changing various parameters, such as the reaction time, the temperature, and the reactant concentration during the material-synthesis procedure. According to previous studies, these strategies mainly affect the positioning of the elements, the potassium or iron content, and the amount of coordinated (interstitial) water in the CuHCF lattice [6,20,35,38,39].
As part of our previous efforts to optimize CuHCF’s properties [6], we adjusted the reactants’ (i.e., Cu(NO3)2 and K3Fe(CN)6) ratio and concentration during the synthesis of CuHCF. Our physicochemical analyses of our pristine CuHCF samples synthesized with different reactant concentrations and ratios indicated different potassium contents in the lattice. In particular, an increasing amount of potassium within the initial CuHCF lattice led to smaller bond distances, namely smaller crystal dimensions, resulting in the higher long-term stability of the material and, therefore, an improved cycle life. The initial potassium content within the lattice of CuHCF was also considered by Ojwang et al. [35,38,40]. In their work, K2S2O3 was incorporated into the synthesis according to Equation (1).
Cu II [ Fe III ( CN ) 6 ] 2 3 . n H 2 O + 2 x / 3 . K + + 2 x / 3 . S 2 O 3 2 − ↔ K 2 x 3 Cu II [ Fe x II Fe 1 − x III ( CN ) 6 ] 2 3 . n H 2 O + 1 x / 3 . S 4 O 6 2 −
Infra-red (IR) spectroscopy measurements associated with synchrotron-based experiments showed that a lower FeIII/FeII ratio in the lattice and smaller cell dimensions were associated with a higher initial potassium content. Variations in the FeIII/FeII ratio in the lattices of hexacyanometallates were also observed by Gerber et al. [41] while employing different precursors, such as Fe, Co, Ni, and Cu, in the coprecipitation reaction of their hexacyanoferrates. Their IR measurements revealed that the utilized metal precursor strongly influenced the peak position and the intensity of the FeIII-CN and of the FeII-CN bands in the IR spectrum. In another study on the magnetic properties of CuHCF, by Ng et al. [42], variations in the FeIII/FeII ratio in the lattice were also observed while investigating the effects of thermal treatment on CuHCF’s crystal structure. According to their findings, annealing the CuHCF powder led to the conversion of the material’s microstructure from an initial composition of FeIII-CN-CuII to FeII-CN-CuIII and CuII-CN- FeIII, with a progressive decrease in the lattice parameter.
In light of the studies mentioned above, both the crystal structure and the chemical composition of CuHCF can be modified by adjusting the synthesis parameters. Moreover, it appears that the potassium and the iron (either FeII or FeIII) contents in the CuHCF lattice play a pivotal role in determining its unit cell’s parameters. There are, however, only a few studies on how these modifications influence the electrochemical performance of the material. The tuning of the crystal structure of PBAs in general, and of CuHCF in particular, is primarily studied in non-battery research areas, and, therefore, assessments of the electrochemical performance are not routinely provided.
In our previous study, the highest initial potassium content related to the smallest unit cells was associated with the longest cycling lifetime of our CuHCF-based electrodes [6]. In this work, we addressed the question of how thermally induced structural changes influence the electrochemical performance of CuHCF. Here, CuHCF powder was synthesized via the routinely used coprecipitation method and thermally treated at various temperatures before its electrochemical testing in three-electrode flooded cells. Structural and chemical analyses were performed on the thermally treated CuHCF powders to correlate the effect of the annealing temperature with the changes in the material’s crystal structure. Moreover, galvanostatic cycling of the CuHCF-based electrodes was performed to assess the material’s cycle life as a function of its thermal treatment.

2. Experimental Methods

2.1. Material Synthesis

Copper hexayanoferrate (CuHCF) was synthesized through the standard coprecipitation method reported in [43]. Briefly, under vigorous stirring at room temperature, two solutions, of 50 mM Cu(NO3)2•3H2O (Sigma Aldrich, Munich, Germany) and 100 mM K3Fe(CN)6 (Sigma Aldrich), were added simultaneously and dropwise to 60 mL of deionized water. A brown suspension was formed immediately, which was subsequently bath-sonicated for 30 min and then allowed to settle overnight. The formed precipitate was centrifuged and subsequently washed with a solution containing 1 M KNO3 (Sigma Aldrich) and 10 mM HNO3 (Sigma Aldrich), followed by rinsing with deionized water, in order to eliminate any remaining impurities or unreacted precursors from the synthesis of the CuHCF. Subsequently, the drying of the CuHCF powder at 60 °C was carried out overnight. The resulting material was then ground using mortar and pestle.

2.2. Electrochemical Characterization

The electrochemical measurements were carried out using a BioLogic VMP3 instrument in a flooded three-electrode cell consisting of a CuHCF-based working electrode, two zinc foils (99.99%, Good Fellow) as counter and reference electrodes, and a 100-millimolar ZnSO4 (Heptahydrate, 99.95%, Sigma Aldrich) aqueous solution as electrolyte. The working electrodes were prepared by hand brushing the CuHCF-based slurry on carbon cloth (Fuel Cell Earth) current collector with a mass loading of approximately 10 mg cm−2. To make the slurry, CuHCF powder, amorphous carbon (Super C65, Timcal), polyvinylidene fluoride (PVdF) (Solef S5130, Solvay, Brussels, Belgium), and graphite (SFG6, Timcal) were dispersed in N-methyl-2-pyrrolidone (NMP) (Sigma Aldrich) with a weight ratio of 80:9:9:2. The dispersion was then mixed thoroughly for 30 min at 4000 rpm using an Ultra-Turrax disperser (T10, IKA). Prior to the assembly of the electrochemical cells, the CuHCF-based working electrodes were annealed at 60°, 80°, 100°, 120°, and 150 °C for 6 h under vacuum. Before each electrochemical test, the open-circuit potential was measured for one minute.

2.3. Material Characterization

The thermo-gravimetric properties of the CuHCF were studied using a NETZSCH STA 449 F3 thermogravimetric analyzer under an inert argon atmosphere at 35–250 °C (5 K min−1 rate) in Al2O3 crucibles.
Prior to the scanning electron microscopy (SEM), IR spectroscopy analysis and X-ray powder diffraction (XRPD), pristine CuHCF powder samples were annealed under the same thermal conditions as the electrodes (namely, at 60°, 80°, 100°, 120°, and 150 °C for 6 h under vacuum). The attenuated total reflectance infrared Fourier transform (ATR-FTIR) spectra were generated and acquired using a Bruker ALPHA II compact spectrometer configured in the mid-IR range. The spectra were acquired with a resolution of 2 cm−1.
To record the XRPD patterns, a Miniflex Rigaku® diffractometer was utilized, with CuKα radiation at room temperature in the 2θ range of 10–60° at a scan speed of 5 s per step and a step width of 0.03°. A quartz holder was used without the help of any solvents to hold the powder sample. The diffractograms were normalized based on their highest peak intensity.
The SEM images were acquired using a FEI Helios NanoLan 600 DualBeam® apparatus with an acceleration voltage of 10 kV. To overcome the inadequate electronic conductivity of the particles, platinum/palladium coatings were employed.

3. Results and Discussion

Initially, the thermal stability of our synthesized CuHCF powder was analyzed with a thermogravimetric analyzer between 35 °C and 250 °C. The thermogravimetric analysis (TGA) was carried out under an inert argon atmosphere. As observed in the thermogravimetric curve reported in Figure 2, upon increasing the temperature, the CuHCF powder showed a mass loss, which slowly started at around 80 °C and was continuous until a temperature of ca. 170 °C was reached. When the CuHCF powder was exposed to a temperature of 148 °C, a sharp, exothermic peak appeared in the corresponding differential scanning calorimetry (DSC) curve. This exothermic peak was attributed to the rapid decomposition of the cyanide groups present in the CuHCF lattice [42,44], which led to the significant degradation of its crystal structure. It is worth noticing that an initial partial decomposition of cyanide groups might start even at temperatures lower than 148 °C. However, the initial decomposition starting point may vary between CuHCF samples synthesized through different synthesis routes and reactant concentrations [42,44]. Based on the peak location on the DSC curve, it appears that our pristine CuHCF powder was stable up to nearly 120 °C, after which it slowly started to partially degrade until completely losing its stability at temperatures ≥150 °C.
Successively, the synthesized CuHCF powders were annealed at different temperatures, namely 60°, 80°, 100°, 120°, and 150 °C. A structural analysis was performed on all the thermally treated CuHCF powders and the untreated powder. The X-ray powder diffractograms of all the analyzed powders (Figure 3) showed a pattern corresponding to an F-centered cubic unit cell (fm3m space group), which was in agreement with the primary reflections corresponding to K2x/3Cu[Fe(CN)6]2/3•nH2O [38].
Despite the similarities in the XRPD patterns of all the analyzed CuHCF powders, a slight 2θ shift and a change in the peak intensity ratio of the 220/200 planes (I220/I200) were observed when increasing the annealing temperature. The shift of the 220 and 200 reflections towards higher 2θ values appeared more prominent when the CuHCF powder was annealed at 120 °C and 150 °C compared to the other thermally treated samples. The changes in the crystal parameters are shown more clearly in Figure 4, which represents the variations in the I220/I200 peak-intensity ratio and the lattice constant in the dependency on the temperature employed during the annealing of the CuHCF powders. These results are in agreement with those of previous works [6,38,42,45]; accordingly, the increased I220/I200 ratio and the shift of the peaks to higher 2θ values indicate a rearrangement of the elements in the structure towards a smaller crystal unit [6,35,38,40].
The SEM analysis showed that there were no significant variations in the average particle sizes or the morphologies of all the thermally treated CuHCF powders (Figure S1), which resembled those of the untreated CuHCF shown in Figure 5.
The ATR-FTIR spectra of all the investigated CuHCF samples are shown in Figure 6. The cyanide complex can be easily identified as sharp bands stretching between 2000 cm−1 and 2200 cm−1 [38]. Interestingly, with the Prussian blue analogues, it is possible to differentiate between the ferrocyanide (FeII-CN) and ferricyanide (FeIII-CN) groups in the IR spectrum because of the C-N band position. Due to the higher oxidation number of the Fe and the stronger σ-bond, the positions of FeIII-CN bond peaks are expected to appear at higher wavenumbers than the peaks of FeII-CN bonds [38]. The two peaks at 2100 cm−1 and 2170 cm−1 in Figure 6 indeed correspond to the FeII-CN-CuII and the FeIII-CN-CuII groups of the CuHCF lattice, respectively. Upon increasing the temperature employed during the thermal treatment of the material, a significant change in the relative peak intensities of the FeIII-CN and the FeII-CN peaks was clearly observed. In particular, as the temperature increased, the intensity of the 2170 cm−1 band (FeIII-CN-CuII) diminished, while the intensity of the 2100 cm−1 band (FeII-CN-CuII) increased. When a temperature of 150 °C was chosen for the annealing, the FeIII-CN-CuII band almost disappeared and, at the same time, the FeII-CN-CuII band displayed a broad shoulder extending to wavenumbers as low as 2020 cm−1. According to other studies [38,42], this shoulder can be attributed to CuII-CN-FeII, and it is expected to be observed at slightly lower wavenumbers than FeII-CN-CuII due to the higher electronegativity of CuII compared to FeII and the lower σ-donation.
From the analysis of the ATR-FTIR spectra, it is possible to extract the relative amount of FeIII with respect to the total iron content of the lattice based on the ratio between the 2100 cm−1 and the 2170 cm−1 bands’ intensities. As shown in Figure 7, in the case of the untreated CuHCF powder, the oxidation state of the Fe was primarily (+3). By contrast, when the CuHCF powders were treated with increasing temperatures, the relative amount of FeIII decreased, and the amount of FeII increased. This trend followed a similar pattern as the lattice constant (Figure 4a), in which increasing the treatment temperature caused the unit cell to shrink. Therefore, it appears that the thermal treatment might have caused a rearrangement of the CuHCF crystal structure, which may have led to a change in the amount and/or type of defects present in the lattice, ultimately resulting in the reduction of a portion of the FeIII present in the crystal structure. However, further investigations are needed to clarify the role of annealing in this change in the oxidation state of the iron atoms in CuHCF lattices.
It is worth noting that within the FTIR spectra, a broad peak was observed for all the investigated samples at around 1600 cm−1, which was attributed to the interstitial water content within the material’s lattice [44,46]. In agreement with the literature [39,42], the FTIR spectra demonstrated that the annealing of the CuHCF up to 140 °C had little effect on the amount of interstitial water, but it primarily eliminated the water adsorbed on the material surface. The intensity of the peak at 1600 cm−1 displayed negligible fluctuations with increasing temperatures up to 100 °C, particularly in comparison with the variations observed in the intensity of the peaks at 2100 and 2170 cm−1. This suggests that the thermal treatment had a more significant effect on the relative FeIII-to-FeII content than on the interstitial water.
The electrochemical performance of the annealed CuHCF electrodes was assessed through galvanostatic cycling in the 100 mM ZnSO4 aqueous solutions at a current rate (C-rate) of 1C. This C-rate was chosen considering that the usual operational currents required by a storage device for the power grid range from 0.5C to 2C [5,47]. According to Table 1, the initial specific discharge capacity (Q0,discharge) decreased slightly when increasing the annealing temperature. In particular, the CuHCF electrode annealed at 150 °C showed an extremely low initial discharge capacity, of 15.2 mAh g−1. This was a direct consequence of the irreversible degradation of the material’s crystal structure at this temperature, as previously shown through the thermogravimetric analysis (Figure 2).
Interestingly, the open-circuit potential (OCP) of the annealed CuHCF electrodes decreased when increasing the annealing temperature. Indeed, before cycling, as the annealing temperature increased, the amount of FeII became predominant within the lattice, while the amount of FeIII decreased. Consequently, the OCP value shifted towards lower potentials due to the lower electronegativity of the FeII.
Figure S2 shows further proof of the increase in the FeII content in the CuHCF lattice upon increasing the annealing temperature. Indeed, the initial oxidation charge in the annealed CuHCF electrodes that was needed to reach the fully charged state during the first cycle (Q0,charge) increased upon increasing the annealing temperatures, due to the higher initial amount of FeII that could be oxidized to FeIII. Both the decrease in the OCP and the increase in the initial Q0,charge represent electrochemical proofs ascribed to the increase in the initial FeII content in the CuHCF lattice, which is in agreement with the ATR-FTIR analysis (Figure 7).
The long-term galvanostatic cycling of the CuHCF electrodes also changed depending on the annealing temperature (Figure 8). In particular, the cycle life of the CuHCF increased with the increase in the annealing temperature up to 100 °C. However, it gradually decreased when annealed at temperatures higher than 100 °C. As shown in Figure 8a, the 100 °C sample retained 80% of its initial discharge capacity after 340 cycles, whereas the untreated electrode achieved the same capacity retention only after 170 cycles.
The CuHCF electrode annealed at 120 °C exhibited a shorter lifespan, of around 290 cycles. Interestingly, the standard deviation in the capacity retention along the cycles was more significant for the 120 °C sample than that the other annealed electrodes. This evident variation in the electrochemical behavior of the CuHCF annealed at 120 °C can be explained by taking into consideration the TG curve (Figure 2): at this temperature, it appears that an initial partial degradation of the CuHCF lattice started to occur, and this may have caused a lower stability in the CuHCF crystal structure when cycled electrochemically. Thus, we believe that both the shorter lifespan and the larger standard deviation were probably due to the start of this partial degradation of the material lattice at 120 °C.
The initial specific discharge energy (E0,discharge) of all the CuHCF electrodes ranged around 80–90 mWh g−1 (Table 1), except the electrode annealed at 150 °C. In the case of the initial discharge capacity, the initial specific energy also slightly decreased with the increasing annealing temperature. In the case of the electrode annealed at 150 °C, the low initial specific energy was due to the degradation of the material lattice that occurred at this temperature, as discussed above.
As with the discharge-capacity retention, the specific energy retention was strongly influenced by the annealing temperature of the CuHCF-based electrodes (Figure 8b). Similarly, the CuHCF electrode annealed at 100 °C reached 80% energy retention after 400 cycles, compared to the 190 cycles retained by the untreated CuHCF.
It is generally known that the cycle life of CuHCF in terms of delivered energy density is longer than that estimated in terms of the electrode’s specific capacity. This behavior can be explained by examining the potential profiles derived from galvanostatic measurements. Figure 9 and Figure S3 illustrate the potential profiles of a selection of cycles for both the untreated CuHCF-based electrode and the annealed electrodes. In all cases, a plateau appeared at 1.6 V vs. Zn2+/Zn at the beginning of the charge curve, which was due to the FeIII/FeII redox associated with the (de-)insertion of the Zn2+. However, the (de-)insertion potential was partially shifted towards higher values, at approximately 1.7 V vs. Zn2+/Zn, with the material cycling. This higher insertion plateau enabled a two-phase Zn2+ insertion in the CuHCF lattice. This can be more clearly visualized with the aid of the differential charge plots (Figure 10 and Figure S4). Here, two pairs of peaks can be observed at 1.6 V vs. Zn2+/Zn and 1.7 V vs. Zn2+/Zn. Notably, the first pair of peaks, at 1.6 V vs. Zn2+/Zn, was present from the beginning of the cycling of both electrodes, whereas the second pair of peaks, at 1.7 V vs. Zn2+/Zn, appeared after approximately 200 cycles. These peaks were correlated with the developing plateau at around the same potential in the galvanostatic profiles. The first pair of peaks at around 1.6 V vs. Zn2+/Zn, corresponded to the FeIII/FeII redox associated with the (de-)insertion of the Zn2+, and the development of the second pair of peaks around 1.7 V vs. Zn2+/Zn was correlated with the development of a two-phase insertion. In the case of the annealed electrodes, the second peaks at 1.7 V vs. Zn2+/Zn were considerably sharper than that of the untreated electrode, particularly after approximately 300 cycles. We previously argued that the occurrence of such a two-phase insertion mechanism at higher potential compensates, at least partially, for the faded capacity of the material. This is the reason for the increased cycle life when the state of health is defined with respect to the specific energy (E) rather than the specific capacity (Q; refer to Table 1) [6,19].
There has been speculation that such phase transitions may affect the aging of CuHCF. In our previous work [6], we suggested that such phase transitions and changes in the potential profile may also be related to the formation of a different insertion site, with Zn2+ occupying vacancies in the CuHCF’s host structure, followed by the nucleation of ZnHCF after long-term cycling.
All the annealed CuHCF electrodes showed very high coulombic efficiency (≥99.5%), regardless of the temperature employed during the thermal treatment, as reported in Figure 11.
It is worth noting that the electrode with the longest cycle life, namely the 100 °C-annealed CuHCF, had an initial iron content of about 45% FeIII and 55% FeII in its lattice (according to the ATR-FTIR analysis), while maintaining its structural stability (as demonstrated by the TGA). The better electrochemical performance (i.e., longer cycle life) of the material could have been due to the stabilizing effect on the crystal structure caused by the thermal treatment at 100 °C, which is likely to have provoked a rearrangement of the lattice. This lattice rearrangement might have been a consequence of a change in the amount/type of the defects present in the CuHCF’s crystal structure, ultimately resulting in a change in the initial FeIII/FeII content ratio.
Since the tuning of CuHCF’s properties through thermal treatment is limited by the possibility of the decomposition and structural deterioration of the lattice, other strategies should be developed in order to utilize an optimal FeIII/FeII content ratio in CuHCF crystal structures to achieve the best possible electrochemical performance.

4. Conclusions

Here, structural changes in a CuHCF lattice were induced by thermally treating the material, with an apparent effect on its cycle life. In particular, it was found that our synthesized CuHCF is thermally stable up to nearly 120 °C. A structural analysis showed that the CuHCF’s crystal lattice shrank upon increasing the annealing temperature. This change in crystal dimensions was attributed to a decrease in the lattice’s FeIII/FeII content ratio, as suggested by our ATR-FTIR analyses, which was also in agreement with the values of the open-circuit potential of the annealed samples and with the initial oxidation charge during the first galvanostatic cycle of the samples. The galvanostatic cycling showed that the CuHCF-based electrode annealed at 100 °C, with a FeIII/(FeIII + FeII) content ratio in the range of 45%, exhibited a longer cycle life of ca. 400 cycles at 1C, compared to the 190 cycles reached by the untreated CuHCF.
Based on our experiments, it appears that the arrangement of elements in the crystal structure, particularly the FeIII/FeII content ratio in the crystal structure, may affect the stability of CuHCF lattices and, therefore, the electrochemical performance of CuHCF, when cycled in mild acidic aqueous electrolytes containing Zn2+. Considering the challenges that are yet to be addressed to increase the stability of CuHCF when cycled in the presence of Zn ions, it is clear that thermal treatment is a suitable strategy to increase its cycle life and, therefore, accelerate the commercialization of aqueous Zn-ion batteries.

Supplementary Materials

The following supporting information can be downloaded at: https://www.mdpi.com/article/10.3390/batteries9030170/s1, Figure S1: SEM images of untreated and treated CuHCF powder; Figure S2: The first galvanostatic cycle of CuHCF-based electrodes; Figure S3: Galvanostatic cycles of all CuHCF-based electrodes; Figure S4: Differential charge plots of all CuHCF-based electrodes.

Author Contributions

Conceptualization, M.B., G.Z. and F.L.M.; formal analysis, M.B.; investigation, M.B. and J.G.; data curation, M.B. and G.Z.; writing—original draft preparation, M.B.; writing—review and editing, G.Z.; supervision, G.Z. and F.L.M.; project administration, F.L.M.; funding acquisition, F.L.M. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the German Federal Ministry of Education and Research (BMBF) grant number FKZ 03XP0204A.

Data Availability Statement

The data are available upon request from the corresponding authors.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Schematic representation of the open-framework structure of a general PBA, where: A: alkaline cation, M,M’: transition metals, C: carbon, N: nitrogen. The structural water molecules have been omitted for the sake of clarity.
Figure 1. Schematic representation of the open-framework structure of a general PBA, where: A: alkaline cation, M,M’: transition metals, C: carbon, N: nitrogen. The structural water molecules have been omitted for the sake of clarity.
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Figure 2. Thermogravimetric curve (black) and differential-scanning-calorimetry curve (red) of the synthesized CuHCF powder.
Figure 2. Thermogravimetric curve (black) and differential-scanning-calorimetry curve (red) of the synthesized CuHCF powder.
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Figure 3. The XRPD patterns of the CuHCF powders thermally treated at different temperatures.
Figure 3. The XRPD patterns of the CuHCF powders thermally treated at different temperatures.
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Figure 4. (a) Average lattice constant of the CuHCF powder as a function of the temperature employed during the thermal treatment, (b) average variation in X-ray-diffraction-intensity ratio between 220 and 200 planes (I220/I200) of the CuHCF powder as a function of the temperature employed during the thermal treatment. The mean values and standard deviations were estimated by comparing at least two different samples resulting from two different thermally treated CuHCF powders.
Figure 4. (a) Average lattice constant of the CuHCF powder as a function of the temperature employed during the thermal treatment, (b) average variation in X-ray-diffraction-intensity ratio between 220 and 200 planes (I220/I200) of the CuHCF powder as a function of the temperature employed during the thermal treatment. The mean values and standard deviations were estimated by comparing at least two different samples resulting from two different thermally treated CuHCF powders.
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Figure 5. SEM image of untreated CuHCF powder.
Figure 5. SEM image of untreated CuHCF powder.
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Figure 6. The ATR-FTIR spectra of the synthesized CuHCF powders thermally treated at different temperatures.
Figure 6. The ATR-FTIR spectra of the synthesized CuHCF powders thermally treated at different temperatures.
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Figure 7. Fraction of FeIII with respect to the total iron content of the lattice as a function of the temperature employed during the treatment of the CuHCF powder.
Figure 7. Fraction of FeIII with respect to the total iron content of the lattice as a function of the temperature employed during the treatment of the CuHCF powder.
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Figure 8. Mean value and standard deviation of (a) specific discharge-capacity retention and (b) specific discharge-energy retention of the synthesized CuHCF-based electrodes galvanostatically cycled at 1C annealed at different temperatures. The mean value and the standard deviations were calculated according to at least two different measurements of two different synthesis batches for each annealing temperature.
Figure 8. Mean value and standard deviation of (a) specific discharge-capacity retention and (b) specific discharge-energy retention of the synthesized CuHCF-based electrodes galvanostatically cycled at 1C annealed at different temperatures. The mean value and the standard deviations were calculated according to at least two different measurements of two different synthesis batches for each annealing temperature.
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Figure 9. Galvanostatic cycles of (a) untreated and (b) 100 °C-annealed CuHCF electrodes, recorded at a C-rate of 1C.
Figure 9. Galvanostatic cycles of (a) untreated and (b) 100 °C-annealed CuHCF electrodes, recorded at a C-rate of 1C.
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Figure 10. Differential charge plots for (a) untreated and (b) 100 °C-annealed CuHCF electrodes.
Figure 10. Differential charge plots for (a) untreated and (b) 100 °C-annealed CuHCF electrodes.
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Figure 11. The mean value and standard deviation of coulombic efficiency of the synthesized CuHCF-based electrodes, annealed at different temperatures and galvanostatically cycled at 1C. The mean value and the standard deviations were calculated according to at least two different measurements of two different synthesis batches for each annealing temperature.
Figure 11. The mean value and standard deviation of coulombic efficiency of the synthesized CuHCF-based electrodes, annealed at different temperatures and galvanostatically cycled at 1C. The mean value and the standard deviations were calculated according to at least two different measurements of two different synthesis batches for each annealing temperature.
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Table
Table 1. Open-circuit potential, average amount of charge exchanged at full state of charge in the first cycle, average initial discharge capacity, average initial discharge energy, and average cycle life based on discharge capacity and discharge energy of all the synthesized CuHCF-based electrodes. The errors in OCP measurements are within ±5 mV. All the values reported within this table were averaged by comparing at least two different samples resulting from two different thermally treated CuHCF powders.
Table 1. Open-circuit potential, average amount of charge exchanged at full state of charge in the first cycle, average initial discharge capacity, average initial discharge energy, and average cycle life based on discharge capacity and discharge energy of all the synthesized CuHCF-based electrodes. The errors in OCP measurements are within ±5 mV. All the values reported within this table were averaged by comparing at least two different samples resulting from two different thermally treated CuHCF powders.
Treatment Temperature (°C)OCP
(V)		Q0,charge
(mAh g−1)		Q0,discharge
(mAh g−1)		E0,discharge
(mWh g−1)		Cycle Life Based on QdischargeCycle Life Based on
Edischarge
Untreated1.6913.6 ± 1.358.1 ± 1.392.5 ± 2.1170190
601.6719.6 ± 0.256.4 ± 0.289.6 ± 0.2210320
801.6521.8 ± 0.254.6 ± 0.287.0 ± 0.4280350
1001.6229.4 ± 0.254.5 ± 0.287.4 ± 0.4340400
1201.4846.6 ± 1.949.8 ± 1.983.3 ± 0.8290340
1501.32-15.2 ± 1.226.3 ± 2.1--
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Baghodrat, M.; Zampardi, G.; Glenneberg, J.; La Mantia, F. Influence of the Thermal Treatment on the Structure and Cycle Life of Copper Hexacyanoferrate for Aqueous Zinc-Ion Batteries. Batteries 2023, 9, 170. https://doi.org/10.3390/batteries9030170

AMA Style

Baghodrat M, Zampardi G, Glenneberg J, La Mantia F. Influence of the Thermal Treatment on the Structure and Cycle Life of Copper Hexacyanoferrate for Aqueous Zinc-Ion Batteries. Batteries. 2023; 9(3):170. https://doi.org/10.3390/batteries9030170

Chicago/Turabian Style

Baghodrat, Mohsen, Giorgia Zampardi, Jens Glenneberg, and Fabio La Mantia. 2023. "Influence of the Thermal Treatment on the Structure and Cycle Life of Copper Hexacyanoferrate for Aqueous Zinc-Ion Batteries" Batteries 9, no. 3: 170. https://doi.org/10.3390/batteries9030170

APA Style

Baghodrat, M., Zampardi, G., Glenneberg, J., & La Mantia, F. (2023). Influence of the Thermal Treatment on the Structure and Cycle Life of Copper Hexacyanoferrate for Aqueous Zinc-Ion Batteries. Batteries, 9(3), 170. https://doi.org/10.3390/batteries9030170

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