IV. Multiscale Hyperporous Silicon Flake Anodes for high Initial Coulombic Efficiency
4.3 Results and discussion
A crystal structure of talc clay minerals is depicted in Figure S1a (see Supporting Information, (SI)), which consist of 2:1 ratio of two different layers, silicate tetrahedral and magnesium oxide layers. It belongs to the so-called smectite group among the various classes of clay minerals.32 Talc clay has tightly stacked flake structure with 1 nm-thick single layer of phyllosilicates, where those regular laminating structures are sustained by interfacial repulsive forces between adjacent layers (see SI, Figure S1b-d). Because it is natural sources, broad distribution in their size (0.5–5 μm) and thickness (10–50 nm) are unavoidable.
One of the scalable production methods for silicon is magnesiothermic reduction process, generally following reaction equation: SiO2 (s) + 2Mg (g) → Si (s) + 2MgO (s) (weight ratio of SiO2 : Mg = 1 : 0.8). Likewise, the clay minerals are subjected to magnesiothermic reduction after considering stoichiometric ratio of reactants. The mixture of talc and Mg was loaded to a stainless steel reactor filled with inert Ar gas (see SI, Figure S2) and gradually heated up to 650 C for 3hours. Above the melting point of Mg, it vaporizes and is trapped inside the reactor on the purpose of preventing escape of Mg vapors, delivering high pressure on overall system. As a result, Mg vapors can easily react with clay minerals beginning from the surface dual phase of silicates and silica.
Subsequently, all of the byproducts are completely etched out after reduction for sufficient time. As schematically shown in Figure 4-1a, the resultant HPSF has uniform distribution of macropores and meso-/micro-pores on the framework. Scanning electron microscopy (SEM) characterization of HPSF shows multi-stacked hyperporous flake structure retaining original morphology of clay minerals (Figure 4-1b and 4-1c). Furthermore, sequential phase transition was investigated by X-ray diffraction pattern, as shown in Figure 1d. Talc has typical XRD patterns for monoclinic clay minerals (JCPDS 13-0558, bottom in Figure 4-1d), which are further corroborated by a high-powder XRD pattern analysis (see SI, Figure S3). Crude products from reduction reaction, cotaines only silicon and MgO particles originated from as-existing MgO octahedral layers and consequence of reaction (JCPDS 89-7746, middle in Figure 4-1d). From SEM images, newly formed particles identified as MgO by energ-dispersive X-ray (EDX) analysis are located in between of layers before etching (see SI, Figure S4). In this intermediate state, as-reduced HPSF samples become thicker up to 200 nm, compared to bare clay minerals (~50 nm), but still preserved flake-type silicon structure are shown. By removing all the MgO-
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relavant byproducts, pure silicon phase of HPSF is obtained (JCPDS 27-1402, top in Figure 4-1d) along with calculated crystallite size of 72 nm.33
What differentiates our system from conventional magnesiothermic reduction is no need to introduce any external sources as a heat scavenger or sink (for example, molten salts) for reducing the excessive amounts of exothermic heat generated from reduction reaction (ΔHrxn = -586.7 kJ/molSilica).34-36 Instead, we utilized the MgO layers sandwiched by two adjacent silicate layers as a sort of heat scavenger, due to its ability to absorb the heat from the reduction reaction by serving as a negative catalyst.37 Because negative catalysts (also known as inhibitors of catalysis) have high specific heat capacity as much as typical salts and accordingly regulate the reaction rate, structural collapse can be greatly prevented.
Apart from MgO layers, any kind of clays which contains inactive metal oxide layers towards metal reductant can reproduce the same phenomena.
Structural superiority of developed 3D HPSF on the strength of negative catalyst layers was examined by transmission electron microscopy (TEM). As presented in Figure 4-2, HPSF has multi-stacked hierarchical porous flake structures with highly pure silicon phase developed (see SI, Figure S5). First, about 100 nm of macropores are regularly developed over the few micrometer-sized flake frame through selective chemical reduction process which is analogous to inverse-opal morphology (Figure 4-2a). By observing the degree of overlapped macropores, we can surmise that more than two flakes are irregularly stacked each other, of which those thickness is still what can be called nanoflakes (<50nm) and substantiated by the scanning transmission electron microscopy (STEM) dark-field image (Figure 4-2b). Interestingly, macropores have distinctive morphologies of bowl-like cavities with end macropores (Figure 4-2c and 4-2d). The width of those cavities is around 150 nm and 100 nm macropores. Arched-sides are rough-textured surface containing plenty of meso-/micro-pores (Figure 4-2e and 4-2f) showing typical polycrystalline phase of silicon based on d-spacing values (0.32 nm) and corresponding Fast-Fourier-Transform diffraction pattern (inset of Figure 4-2b).
As control experiment, we figured out whether adding NaCl salts act as heat scavenger and/or competitively participate in the magnesiothermic reduction process of talc. As- synthesized products with NaCl involved (denoted as NPSF, non-porous silicon flake) do not contain any sort of pores on the framework. Flake-like structure is well-retained after magnesiothermic reduction, while multi-scale porosity was not observed. The NPSF has even poor purity confirmed by EDX analysis. In addition, small circular plate or particles (100-200 nm) are uniformly distributed over the structure, considered as intermediate form of macropores. This result implies that two kinds of heat scavengers are competitively
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involved in the reduction system. Then, most of thermal energy is taken away from clays to NaCl salt which have much affinity by forming molten salt than MgO layers. As a result, MgO layers cannot be activated enough to generate regular macropores, which left behind just trace of it (see SI, Figure S6).
Nitrogen adsorption analysis undergirds the hyperporosity on framework of silicon flakes (Figure 4-2g). HPSF has Brunauer-Emmet-Teller (BET) specific surface area of 188 m3 g-1 and broad pore size distribution ranging from 1 to 100 nm in the Barret-Joyner- Halenda (BJH) curve (Figure 4-2h). More specifically, major portion of pores originate from macropores on the frame occupying over 80% (0.306 cm3 g-1) of total pore volume (0.368 cm3 g-1). The rest part of total pores is composed of mesopores (~17%, 0.061 cm3 g-1) and micropores (<0.2%, 0.001 cm3 g-1), respectively, as summarized in Figure 4-2i.
Unlike actually existing 5-nm-sized mesopores, protruding peak at around 50nm have contributed to the irregularly stacked silicon flake rather than the meso-porosity on the framework itself. As the 100-nm-sized macropores overlap each other, almost half of macropores could be veiled in the nitrogen adsorption measurement. Also, tiny micropores less than 2 nm were observed high-resolution TEM analysis (see SI, Figure S7).
Formation of hyperporosity was investigated in a systematic manner. The external MgO particles are formed in two different state during the magnesiothermic reduction. At initial stage, only outermost oxygen atoms meet the vaporized Mg to generate MgO particles on the surface generating a little expanded channels, where the Mg vapors penetrate the interface of the silicate and existing-MgO layers shared by the oxygen atoms. Unlike the surface reduction, the internal formation of MgO can make a continuous phase with the existing-MgO layers. The removal of the former brings many oxygen vacancies at a few nanoscale. Then, the latter leads to the formation of macropores and cause structural instability resulting in exfoliation of nanoflake from tightly stacked talc clay minerals.
Direct evidence for mechanism of pore formation is corroborated by time-dependent TEM analysis shown in Figure 4-3. All of the samples are characterized after acid-etching process. At t=10min, there are no macropores on the framework, rather showing shredded structures mainly composed of mesopores (Figure 4-3a and 4-3b). For insufficient reaction time, Mg vapors are only to react with surface silicate layers and cannot proceed internal penetration. Consequentely, it failed to realize the hyperporous structures. From t=30 min, macropores start to show up, however, its structure is not well-defined as well as has thicker frame/larger pores. On the edges of flakes, highly mesoporous region still exist (Figure 4-3c and 4-3d). Eventually, at t=60min, regularly formed macropores are observed and crystal structure successfully grows up to the edges of flakes (Figure 4-3e
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and 4-3f). As expected, the framework can be adopted to exfoliation of talc clays in a level of single layers of hyperporous flakes.
Along with time-dependent structural evolution, we also investigated temperature- dependency on the structure of HPSF at a fixed reaction time for 3 h (see SI, Figure S8).
The TEM images show representative structures prepared at different temperature (500- 700 oC). Regular porous structures start to appear at 550 oC, while randomly formed aggregates are shown at 500 oC (see SI, Figure S8a and 8b). At lower temperature below the melting point of Mg, lateral dimension of each HPSFs becomes smaller than higher one and instead macropores larger than 200 nm are observed. However, higher operating temperature rather gives aggregation on the whole structure by reducing pore size due to exccessive thermal energy (see SI, Figure S8c-S8e). Crystallinity of HPSFs are slightly afftected by reaction temperature except the point that reaction underwent below the melting temperature of Mg did not go to completion (see SI, Figure S8f). Based on its regularity, average particel size and purity, HPSF prepred at 650 oC shows the most proper range of materials for further application.
The X-ray photoelectron spectroscopy (XPS) measurement were carried out to identifty the valence state of silicon in HPSF materials. The XPS spectrum of Si 2p are present in Figure 4-4a, in which several characteristic peaks appearing from 99.5 to 103.5 eV assigned to various oxidation states of silicon. The two primary peaks of atomic silicon (Si0), which consists of 2p1/2 and Si 2p3/2, are displayed along with binding energies. The surface oxides and crystallinity of silicon are also verified by Raman analysis (Figure 4- 4b). Compared to the XPS results, negligible shoulder peaks attributed to amorphous surface oxides are observed between 300 and 450 cm-1. The Raman spectrum shows no amorphous silicon peak at 480 cm-1, while a typical crystalline silicon peak at 516 cm-1 appears. This value is lower than that of bulk silicon at 520 cm-1,38 because HPSF attains enhanced crystallinity from high-temperature magnesiothermic reduction, and many pores on the surface partly reduce the silicon dimension.
On the strength of unparalleled structural originality, we have demonstrated the 3D HPSF as a promising anode material for LIBs. To redeem the low intrinsic electric conductivity of silicon, thin carbon layers were introduced through the thermal decomposition of acetylene gas at 900 oC for 5 min. The carbon-coated 3D HPSF (denoted as HPSF@C) contains carbon layers (with contents of ~10 wt%) covering its whole silicon surface (see SI, Figure S9a-c). After the carbon coating, the HPSFs still preserved their distinctive structure with approximately 9-nm-thick carbon layers (see SI, Figure S9d).
Moreover, the Raman spectrum of the HPSF@C exhibited a disordered (D) and graphene
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(G) band (D/G) ratio of 2.19, indicating that the amorphous carbon layers were coated on the surface of the HPSF (see SI, Figure S9e). Both surface area and pore volumes were slightly reduced to 127 m2 g-1 and 0.317 cm3g-1 after carbon coating (see SI, Figure S9f and S9g). Bulk macroporous silicon materials are also prepared via typical metal-assisted chemical etching (see SI, Figure S10).
The electrochemical performance of both the HPSF and HPSF@C was evaluated by using CR2016-type coin cells in the potential window from 0.01 to 1.2 V. All the displayed capacities in Figure 4-5 were calculated based on only active silicon materials excluding carbon blacks and binders. The first cycle discharge capacities of the HPSF and HPSF@C anodes were 2620 and 3055 mAh g-1, respectively, in the initial galvanostatic measurements at a current density of C/20 (150 mA g-1), which correspond to remarkably high initial coulombic efficiencies of 92.6% and 92.7%, respectively, comparing to that of bulk macroporous silicon (88.8%, denoted as p-Si@C) (Figure 4-5a). This exceptionally high reversibility of the HPSF is attributed to the fact that it retains its macro-sized structure with dominant macropores. As mentioned, nanostructured materials have inevitably higher surface areas primarily occupied by either meso- or micropores, which enables the accessibility of lithium ions and simultaneously could be an origin of the irreversibility of lithium ions. Even though the HPSF shows a relatively higher surface area than conventional bulk silicon materials, it can avoid large irreversibility by minimizing the portion of meso- and micropores, as shown in Figure 4-2i. Moreover, high crystallinity of HPSF also contributes to high initial coulombic efficiency of HPSF anode.
The synergistic coupling of suitable porosity/pore size and high crystallinity of HPSF can significantly enhance the initial coulombic efficiency. In contrast, other alloy-based anodes including Si, Ge, Sn and those of nanostructures, usually display much lower initial coulombic efficiency (50-80%) than commericial graphite-based anodes as illustrated in Figure 4-5b.14,39-44
Through the differential capacity (dQ/dV) plots of the HPSF during prolonged cycles, we can estimate the juncture of the phase transition from crystalline to amorphous Si (see SI, Figure S11). A sharp peak at 0.08 V is related to the phase transition of crystalline silicon to amorphous lithium silicide (LixSi) during the first discharge (lithiation) reaction.45-47 From the subsequent discharge cycles (2nd, 5th, 10th, and 100th), two apparent peaks at 0.27 and 0.09V are observed, indicating the successive transition of LiSi to Li7Si3 and finally the amorphous Li15Si4 phase.45-47 Meanwhile, charge plots show characteristic peaks related to the dealloying reaction at 0.27 (Li15Si4 to Li7Si3) and 0.47V (Li7Si3 to LiSi).45-47
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The cycling stability of HPSF anodes at a rate of C/5 (600 mA g-1) is shown in Figure 4-6a. Both the HPSF and HPSF@C anodes exhibited superior cycle retention of 89.1%
and 94.6% after 100 cycles, respectively, unlike the high irreversible capacity and drastic capacity fading observed in typical micro-Si anodes at the early stage of cycling.48-51 Even though macropores are uniformly formed on p-Si@C samples along with carbon coating, it also shows severe degradation and failed before 100 cycles. It is typically due to a break- down of the electronically conductive network from the large volume expansion of Si anodes, theoretically over 300%. Most of the active materials form a wrapping-like network with polymeric binders by solely covering the exterior surfaces. Then, it cannot be restored after volume contraction during the dealloying reaction unless highly soft binders are used. However, the HPSF has a very durable network during cycles, because it has open spaces over 100 nm bored through the framework and a thinner frame size of less than 50 nm. These characteristics allow the polymeric binders to not only cover the entire exterior surfaces but also penetrate the interior framework. Accordingly, it can form a more effective linkage with neighboring stacked silicon flakes and a sustainable network (see SI, Figure S12). To demonstrate this aspect, we fabricated an electrode composed of the HPSF and polymer binders on a hard substrate. As depicted in Figure S13a, the binders effectively covered the entire surface of the HPSF particles, and they were intertwined with the macropores of the HPSF (see SI, Figure S12b-S12d).
The HPSF anodes were further monitored for extended cycles (Figure 4-6b). In the case of the bare HPSF anode, its capacity retention was rapidly decayed at a rate of C/2 (1500 mA g-1) after 200 cycles, which corresponded to a capacity retention of 42.8%, owing to the typical problem of unstable solid-electrolyte-interphase (SEI) formation and poor conductivity shown in the micro-sized silicon anode. In contrast, the HPSF@C delivered a high charge capacity of ~1619 mAh g-1 after 200 cycles at a rate of C/2, corresponding to an improved capacity retention of 95.2% compared to its initial charge capacity.
Surprisingly, the HPSF@C electrode cycled at a rate of 1C retained 74.4% after 400 cycles, still outperforming commercial graphite anodes in terms of capacity, despite the micro-Si anode. These remarkable results are attributed to the synergistic coupling of the unconventional structural network and the formation of a stable SEI induced by carbon coating layers.
The rate capability of the HPSF anodes was evaluated at various rates of C/5, C/2, 1 C, 2 C, 5 C, and 10 C for each of the five cycles. The obtained results are shown in Figure 4- 6c. The distinctively structured HPSF with carbon layers on its surface provides the HPSF@C anode a high current density in spite of its relatively large particle size. This is
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attributed to the reduced Li ion diffusion length from the thinner frame size (< 50 nm) and the accessibility of electrolytes through the large number of huge macropores and meso- /micro-pores on its framework. In principle, the reduced dimensions increase the rate of lithium ion transport into the active materials considerably. The diffusion constant for alloying reactions increases with the square of particle size in nanometer.6 In this regard, the HPSF has significantly reduced dimension in thin flake structure, which leads to enhancement of both lithium ion and electrolytes. Furthermore, high surface area permits a high contact area with the electrolyte to increase the high lithium-ion flux across the interface, which is related to accessibility of lithium ions. In addition, carbon layers help the formation of thin and uniform SEI layers over the HPSF surface. Interestingly, it delivered a reversible charge capability of ~580 mAh g-1 at a rate of 10 C, the most exceptional results among the micro-Si anodes.
In order to further demonstrate the stability of HPSF anodes in repeated cycles, the cells were disassembled after the cycling test in the Ar-filled glovebox. After rinsing the HPSF anodes with high purity dimethyl carbonated (DMC) solvent and completely drying them, electrode swelling and morphological changes were characterized by SEM and TEM analyses. The cross-sectional SEM images of bare HPSF electrodes showed a huge volume expansion (152%) after 100 cycles at a rate of C/5. Thereafter, it continued to expand up to 200% after 200 cycles (see SI, Figure S13a-S13c). In contrast, the pulverization problem was remarkably suppressed in the case of HPSF@C electrode (66% and 77%) after 100 and 200 cycles, respectively (see SI, Figure S13d-S13f), reflecting a great consistency in electrochemical performance. This can be attributed to several factors such that, firstly, HPSF@C electrode can accommodate large volume expansion with a help of macropores on the framework retaining structural integrity. Second, HPSF has multi-stacked structure even though neighboring layers are not chemically bonded. Stacked system can mitigate pulverization in some degree. In addition, polymeric binders can penetrate the open channels on flakes and then make an effective linakage with neighboring flakes to suppress the pulverization unlike bulky porous Si materials (see SI, Figure S12). From the top view SEM images of cycled electrodes, bare HPSF electrodes were severely pulverized (see SI, Figure S14a-S14c), while HPSF@C electrodes retained stable networks with binders and conductive carbon showing a correspondingly flat surface (see SI, Figure S14d-S14f).
This different result originates from the stabilized interface at the early stage of cycling, which prevents the continuous decomposition of electrolytes. Non-coated HPSFs were transformed into sponge-like structures with all the closed pores. We could not observe 100-nm-sized macropores or the meso-/micro-pores anymore or even their framework
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(Figure 4-7a-c). It was assumed that the distinctively structured HPSF had undergone repetitive expansion and contraction in an isotropic orientation and then the macropores totally collapsed as different morphologies emerged (Figure 4-7d-f). Meanwhile, flake- like structures were retained in the HPSF@C samples with reduced-size macropores. It may be attributed to surface SiOx layers and thin carbon coating layers on the HPSF@C, which can act as good promoter to form stable SEI layers. Furthermore, newly developed mesopores were observed on the framework. Since it did not lose its structural originallity, it could lead to outstanding electrochemical performance in terms of both cycling stability and rate capability for long-term tests. Compared with other porous Si anodes, the HPSF anode has superior cycling performance and rate capability, and more importantly, these were realized in a scalable manner (see SI, Table S1). The disadvantage of HPSF@C materials synthesized in this work is low tap density of materials owing to highly porous and thin flake structure. Even though it is lower than conventional bulk Si, the HPSF has a few micrometer dimension in lateral direction with hundreds of nanometer in their thickness. Thus, it has much higher tap density than commercial nanoparticles or other porous structures.