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Cite this: RSC Adv., 2019, 9, 41228

Silica from diatom frustules as anode material for
Li-ion batteries†
Andreas Nicolai Norberg,a Nils Peter Wagner,ab Henning Kaland,a Fride Vullum-Bruerc
and Ann Mari Svensson a
In spite of its insulating nature, SiO2 may be utilized as active anode material for Li-ion batteries. Synthetic
SiO2 will typically require sophisticated synthesis and/or activation procedures in order to obtaina satisfactory performance. Here, we report on diatom frustules as active anode material without theneed for extensive activation procedures. These are composed primarily of silica, exhibiting sophisticatedporous structures. Various means of optimizing the performance were investigated. These included
carbon coating, the addition of ﬂuoroethylene carbonate (FEC) and vinylene carbonate (VC) to thecarbonate-based electrolyte, as well as activation by an initial potentiostatic hold step. The highest
capacity (723 mA h g 1) was obtained with composite electrodes with pristine diatom frustules andconventional carbon black as additive, with the capacity still increasing after 50 cycles. The capacity wasaround 624 mA h g 1 after subtraction of the contributions from the carbon black. Carbon coated

Received 10th September 2019
Accepted 3rd December 2019

diatom frustules showed a slightly lower but stable capacity after 50 cycles (600 mA h g 1 after
subtraction of contributions from the carbon coating and the carbon black). By the use of

DOI: 10.1039/c9ra07271c

electrochemical characterization methods, as well as post-mortem studies, diﬀerences in reaction

rsc.li/rsc-advances

mechanisms could be identiﬁed and attributed to the operating and processing parameters.

1. Introduction
Rechargeable lithium-ion (Li-ion) batteries are already widelyused and are the dominating solution for automotive as well as
portable consumer devises and high-power applications due tothe superior power and energy density. However, several challenges of Li-ion battery technology are yet to be overcome, e.g.
with respect to cost, resource availability, and safety issues
related to risk of thermal runaway of the batteries.1–3 Anotherissue with Li-ion batteries is the relatively poor environmental
footprint related to the production of the batteries.4 Since Li-ion
batteries are expected to dominate the battery market for
portable electronics as well as electric vehicles in the foreseeable future, it is important to address and solve the abovementioned issues by for example developing and identifying
new electrode materials as well as other components used in the
production of a complete battery. Current state-of-the-art anodea

Department of Materials Science and Engineering, Norwegian University of Scienceand Technology, NTNU, Sem Sælands Vei 12, Realfagbygget K2-132, 7034
Trondheim, Norway. E-mail: andreas.n.norberg@ntnu.no; henning.kaland@ntnu.no;annmari.svensson@ntnu.nob

Department of Sustainable Energy Technology, SINTEF Industry, Strindveien 4, 7034
Trondheim, Norway. E-mail: nils.peter.wagner@sintef.nocDepartment of Thermal Energy, SINTEF Energy, Sem Sælands Vei 11, 7034
Trondheim, Norway. E-mail: fride.vullum.bruer@ntnu.no† Electronic supplementary
10.1039/c9ra07271c

information

41228 | RSC Adv., 2019, 9, 41228–41239(ESI)

available.

See

DOI:

materials (graphites) have a specic theoretical capacity of
372 mA h g 1 and are used in most conventional Li-ion
batteries.5 Approximately 10% of this capacity is irreversibly
lost in the rst cycle due to decomposition of electrolyte on thegraphite surface. Recently, small amounts (5–12 wt%) ofsilicon (Si) or SiOx is mixed into the graphite anode for a small
increase in battery capacity.6–8 Signicant research eﬀort during
the past 10 years has been focused on developing Si-graphite
composites and Si as the active material, due to its extremely
high volumetric and gravimetric capacities, which is up to 10
times that of graphite.9 However, due to the issues caused bylarge volume expansion during cycling and excessive SEI
formation, the implementation of Si anodes is yet to be realizedin commercial batteries. It is foreseen, that the next generationof Li-ion batteries, the so-called generation 3b, will have Sigraphite composite anodes with a higher content of Si thanstate-of-the art Li-ion batteries.10 From an environmental
perspective, there are also challenges associated with Si thatneeds to be addressed. Most of the pure grade Si in the world isobtained from crystalline silica (SiO2) through a reaction withcarbon at elevated temperatures, resulting in a high CO2 footprint in the range 3.5 to 11 kg CO2 per kg Si produced(depending on origin of the energy used). Battery grade graphite
is either natural graphite (39% of the marked), or synthetic
graphite, produced from residues from oil reneries. Naturalgraphite is on the list of critical raw materials for battery
production,11 while synthetic graphite is also associated withThis journal is © The Royal Society of Chemistry 2019

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signicant CO2 emissions, as it is made from the residues fromoil reneries.
As an alternative to Si, SiO2 has emerged as a potential
abundant and environmentally benign anode material with anestimated capacity between 749 and 1673 mA h g 1.12 Initially,
SiO2 was believed to be electrochemically inactive towards Li.
However, in 2001 Gao et al. reported that SiO2 nanoparticles
were reactive towards Li-ions at potentials between 0 V and1 V.13 Sun et al. was the rst to propose a lithiation mechanism,
which was further expanded on by Guo and Chang et al.
resulting in the following proposed mechanisms:12,14,15
5SiO2 + 4Li+ + 4e 4 2Li2Si2O5 + Si

(1)

2SiO2 + 4Li+ + 4e / Li4SiO4 + Si

(2)

SiO2 + 4Li+ + 4e / 2Li2O + Si

(3)+

Si + xLi + xe / LixSi(4)

By far the most common method for preparing SiO2 forbattery anodes in the literature is the preparation of porous SiO2
structures by wet chemical methods.16–24 However, Chang et al.
have also demonstrated that it is possible to achieve a capacityof approximately 750 mA h g 1, simply by milling quartz.12 In
practice, nanostructures of silica have shown stable cycling
capacity of up to 1000 mA h g 1 without suﬀering from similar
performance issues as Si.17 The improved cycling stability of
SiO2 compared to Si, is believed to originate from the formationof lithium oxide (Li2O) and lithium silicates, which buﬀer the
expansion of Si during lithiation.14,25,26 To further limit thevolume expansion of Si and improve the otherwise sluggish
electronic conductivity of SiO2, diﬀerent carbon coatings, orattempts to embed the SiO2 particles in a carbon matrix, areoen utilized.16,21,24,27–30
In the work of Lepoivre et al.,31 two batches of pure andmonodisperse SiO2 spherical particles with a diameter of 150/
200 nm and 500 nm were prepared. A capacity of 400 mA h g 1
was obtained for electrodes made from the material with
200 nm particles, aer a so-called potentiostatic discharge(potential held at 2 mV for 250 h prior to cycling). Thus, partial
conversion of the material to silicon and Li15Si4 and Li4SiO4 wasachieved. Full conversion was not observed, which was attributed to the insulating character of SiO2, preventing the reduction beyond a certain penetration depth (45/50 nm).32,33
The drawback with all these materials is that they require quite
complex and expensive synthesis methods. On the other hand,diatom frustules found in the seabed are composed primarily ofsilica, exhibiting porous and sophisticated structures ranging insize from 10 nm to 1 mm.34,35 These structures might be benecial
to Li-ion batteries by easing electrolyte penetration and shortening the diﬀusion path for Li-ions between the electrolyte andthe active material.36,37 One research group has previously reported the use of diatom frustules as active material in Li-ion
batteries.38–40 Here, the diatom frustules were used in combination with red algae as anode material. Half-cells were tested witha capacity of 500 mA h g 1 aer 80 cycles.

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In this work, we investigate the use of a diﬀerent kind of
diatom frustules as anode material in Li-ion batteries andinvestigate various means of optimizing the performance.
These include carbon coating, the addition of uoroethylene
carbonate (FEC) and vinylene carbonate (VC) to the carbonatebased electrolyte, as well as activation by an initial potentiostatic hold step. The highest capacity (723 mA h g 1) was achieved with composite electrodes with pristine diatom frustules,
no electrolyte additives and conventional carbon black
conductive additive. Using electrochemical methods, as well as
post-mortem studies, diﬀerences in reaction mechanisms could
be identied and attributed to the operating and processing
parameters.

2.

Experimental

2.1 Pre-processing of diatom frustules and materials
preparation
The pre-processing and carbon coating of diatoms described inthis work are based on the process described in a patent by
Vullum-Bruer et al.41 Sea-hauled coscinodiscus diatoms (Planktonic AS) were separated from seawater by a 36 mm mesh sieveand dried for 24 h at 90  C. The dried diatoms were rinsed under
running deionized-water (DI water), heated on a hotplate to130  C and stirred at 500 rpm for two hours at a diatom to DI–
water ratio of 1 : 100 wt%. Following heating, fresh DI–water wasadded at the same ratio and the beaker was sonicated for 30 min
before the diatoms were washed under running DI–water for
5 min. Finally, the diatoms were dried at 90  C for 24 h followed
by 150  C for 24 h and calcined in a tube furnace at 650  C for 2 h
in a synthetic air atmosphere. Calcined diatoms were milled byplanetary milling at 600 rpm for 40 min at 10 min intervals witha 5 min break between the intervals. To investigate the eﬀect ofcarbon coating, milled diatoms were mixed with sucrose (Sigma
Aldrich) dissolved in DI–water as a carbon precursor with a ratioof 20 to 80 wt% silica and sucrose, respectively. The solvent wasevaporated on a hotplate at 40  C while stirring at 300 rpm, followed by a heat-treatment in a tube furnace at 650  C underowing Ar atmosphere. The pristine and carbon coated frustules
are named SiO2/P and SiO2/C, respectively.
2.2

Materials characterization

Surface area and porosity data were determined by nitrogen
adsorption (Micrometrics Tristar 3000). Samples were degassed
for 12 h under vacuum at 250  C prior to analysis. Particle sizeand particle size distribution were measured in solution by laser
diﬀraction (Horiba Partica LA-960). To limit agglomeration, thesample solution was sonicated for 30 min prior to themeasurement. The carbon content was measured by thermogravimetric analysis (TGA, Netzch STA 449C Thermal Analysis
System) with a heating rate of 10  C min 1.
Scanning electron microscopy (SEM, Hitachi S-5500) wasused to study the morphology of the diatom frustules. Focusedion beam (FIB, FEI Helios Nanolab Dual Beam) cross-sections
were prepared and electrodes analyzed by SEM (FEI Apreo)
before and aer cycling.

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2.3

Electrode preparation

A slurry was prepared by mixing the active material, pristine
SiO2 (SiO2/P) or carbon coated SiO2 (SiO2/C) with 35 wt% carbon
black (CB, Timcal C-Nergy C65) and 15 wt% sodium alginate(Sigma Aldrich Sodium alginate). In order to investigate thecapacity contribution of carbon, two reference electrodes wereprepared with pyrolyzed sucrose and CB as the active material,otherwise the electrodes were the same composition. The slurries were mixed by ball milling for 45 min in stainless steel
containers using a RETSCH mixer mill, casted on a 10 mm thickcopper foil and dried overnight at 90  C in a vacuum oven.
Capacity is reported with respect to the mass of the SiO2 or the
carbon coated SiO2, unless otherwise stated.

2.4

Paper
carbonate (VC, Sigma Aldrich) with the ratio 94 : 5 : 1 in vol%,respectively. Cells with the additives are named SiO2/P_FEC-VC
and SiO2/C_FEC-VC. Galvanostatic cycling was performed using
BCS 805 (Biologic) and CT2001A (Lanhe) galvanostats between
0.01 and 2.0 V vs. Li/Li+. In addition, an electrochemical reduction program was used by introducing a 48 h potentiostatic stepat 2 mV vs. Li/Li+ in the second discharge cycle, in accordancewith the procedure suggested by Lepoivre et al.31 This potential
was chosen because it is below the lithiation potential of SiO2,
but above the potential of Li-plating. The current density was
50 mA g 1 for the rst cycle and the electrochemical reduction
step, and 200 mA g 1 for the subsequent cycles. All cells had an
active material mass of 0.40 mg, apart from the reference cell not
subjected to a potentiostatic step, which had an active material
mass of 0.45 mg.

Electrochemical characterization

Electrochemical characterization was carried out in CR2016 coincells. Li-foil was used as counter electrode and a Celgard 2400 as
separator. Two electrolytes were used; a commercial 1 M LiPF6 in
50 : 50 vol% EC : DEC (Sigma Aldrich) and an in-house madeelectrolyte prepared by mixing the commercial electrolyte with
uoroethylene carbonate (FEC, Sigma Aldrich) and vinylene

3.

Results and discussion

Results from the materials characterization of the preprocessed, as well as pre-processed and carbon-coated frustules are shown in Fig. 1. The XRD patterns in Fig. 1(a) show theabsence of crystalline phases in the samples. The broad peakFig. 1 (a) XRD patterns of SiO2/P and SiO2/C extracted from diatom frustules, (b) size distribution of milled and un-milled SiO2/P obtained bylaser diﬀraction, (c) surface area of SiO2/P_unMilled, SiO2/P and SiO2/C, (d) thermogravimetric curves of SiO2/C.

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centered at 22 is associated with amorphous SiO2.21,42 EDX
analysis of the frustules conrmed that the frustules mainlycomprise of Si and O (Fig. S1†). However, small amounts of Na
and Fe was identied. The presence of Na can be explained by
the Na-alginate binder, while diatom frustules are known tocomprise of small amounts of Fe.43 However, it is also possiblethat the Fe was introduced from the stainless steel jar used forslurry mixing. The particle size distributions of the milled andun-milled SiO2 determined by laser diﬀraction can be seen in
Fig. 1(b). The mean and the median size of the milled materialis 3.1  0.2 and 3.8  0.2 mm, respectively. As evident from
Fig. 1(b), the milling results in a bimodal size distribution, withthe maximum peak intensity of the two size distributions at
0.445 mm and 5.87 mm. The BET surface area and the t-plot
areas, obtained from nitrogen adsorption data, are shown in
Fig. 1(c). Milling of the SiO2 increased the BET surface area from
8.2 to 17.2 m2 g 1, and this increase was almost exclusively inexternal surface area. The carbon coated SiO2 shows a signicant increase in BET surface area, to 289 m2 g 1. The increase insurface area is a result from an increase in both micropores andexternal area. The carbon content of the coated SiO2 wasdetermined by TGA to be 46.5 wt% (Fig. 1(d)).

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The microstructure and morphology of the extracted diatoms can be seen in Fig. 2(a) and (b). A magnied view of thesurface in Fig. 2(b) is provided in Fig. 2(c). Fig. 2(d) showsa high-resolution micrograph of the porous surface of a frustulelayer. A comparison to AFM micrographs obtained by Losic
et al.,34 indicates that a large fraction of the structures is of thespecies coscinodiscus. The main structures of the frustules
appear to be in the range of 50 nm to 3 mm, which is in agreement with Losic et al.34 In the high-resolution micrograph inFig. 2(d), even smaller structures with dimensions below 50 nm
can be seen. The frustules exhibits features ranging fromroughly 3 mm to 40 nm, where the larger features originatesfrom the frustule macro structure (diﬀerent layers that make up
the frustule) seen in Fig. 2(a)–(c), while the smaller features seenin Fig. 2(d) could be the fundamental building blocks of silica,
as described by Schmid et al.35 In contrast to the frustule
macrostructure, the smaller features might provide an interconnected mesoporous structure that goes through the entire
layer.
Fig. 3(a) compares the cycling performance of electrodes
fabricated from pristine milled diatom frustules (SiO2/P) andmilled carbon coated diatom frustules (SiO2/C), cycled in 1 M
LiPF6 EC : DEC electrolyte with and without additions of 5 andFig. 2 (a) Overview of multiple frustule layers, (b) cross-section of one layer of a frustule, (c) surface of frustule in (b), (d) high-resolution
micrograph of frustule surface.

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SiO2 (SiO2/P and SiO2/P_FEC-VC) exhibit lower capacity thancells with carbon coated SiO2 (SiO2/C and SiO2/C_FEC-VC).
However, the cells with pristine SiO2 experience a greater
increase in capacity over the subsequent 50 cycles. Aer 50
cycles, the cell with pristine SiO2 cycled without the FEC and VC
electrolyte additives was the best performing, with a capacity ofOpen Access Article. Published on 12 December 2019. Downloaded on 9/24/2024 1:02:35 AM.
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1 wt% of FEC and VC and subjected to a potentiostatic step
(2 mV vs. Li/Li+ for 48 h). The potentiostatic step was introduced
in order to activate the SiO2 as previously suggest by Lepoivre
et al.31,35
The potentiostatic step in the second cycle improved thecapacity of all cells signicantly. Initially, cells with pristine

Paper

Fig. 3 (a) Speciﬁc discharge capacity obtained for pristine (SiO2/P) and carbon coated (SiO2/C) diatomic frustules cycled in 1 M LiPF6 EC : DEC
electrolyte or 1 M LiPF6 EC : DEC with 5/1 wt% FEC/VC (FEC-VC) after a 48 h potentiostatic hold step at 2 mV vs. Li/Li+ in the second cycle (withthe exception of SiO2/P_NPS, which was not subjected to any potentiostatic hold step), (b) capacity of cells in (a) with the capacity contribution ofcarbon is subtracted, (c) coulombic eﬃciency of same cells with the ﬁrst 5 cycles as insert, (d) potential proﬁle of cycle 1, 3, and 50 for SiO2/P vs.
Li/Li+, (e) measured current through the respective cells during the potentiostatic step.

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723 mA h g 1. In contrast, FEC and VC appears to improve thecapacity of the carbon coated SiO2. The capacities demonstrated
here are higher than previously shown with silica derived fromseawater diatom algae, where stable cycling at 523 mA h g 1 wasdemonstrated over 20 cycles (30 wt% carbon black in the electrodes), and 500 mA h g 1 aer 80 cycles.38–40
For a fair comparison of the electrodes, in Fig. 3(b), the total
contribution of the carbon coating (pyrolyzed sucrose) andcarbon black was subtracted from the capacity and speciccapacity was normalized to the mass of SiO2 in the electrodes.
The capacity of carbon black as well as pyrolyzed sucrose wasdetermined in separate experiments with electrodes made fromthe respective materials, showing stable cycling for 100 cycles at
a capacity of 274 mA h g 1 and 560 mA h g 1, respectively(Fig. S2†). Comparing Fig. 3(a) and (b), the signicant capacity
contribution from carbon black and the amorphous carbon
coating is evident, emphasizing the importance of accounting
for the capacity contribution of carbon when dealing withamorphous carbon coatings. With the exception of an overall
reduction in capacity, the same trends can be observed inFig. 3(b) as in Fig. 3(a). Initially, electrodes made from themilled and carbon coated frustules cycled with FEC and VC
containing electrolyte exhibits the highest capacity, which israther stable upon cycling (maintaining a capacity of
600 mA h g 1 aer 50 cycles). Pristine milled frustules cycled in
an additive free electrolyte have lower initial capacity. However,these electrodes show improved capacity upon further cycling,increasing to 625 mA h g 1 aer 50 cycles. The materials show
remarkably good stability over the rst 50 cycles with all cells
exhibiting a nal capacity higher than the initial. This stability
is believed to originate from the formation of Li2O, Li4SiO4, andLi2Si2O5, which have been suggested to provide protection
against the destructive volume expansion during lithiation ofsilicon.14,25,26
Fig. 3(c) show the corresponding coulombic eﬃciencies (CE)of the cells over 50 cycles. In the rst cycle, CE is low for all cells,
ranging from 38.9% for pristine SiO2 cycled in an electrolytewith FEC and VC (SiO2/P_FEC-VC), to 46.1% for carbon coated
SiO2 cycled in an additive-free electrolyte (SiO2/C). The initially
low CE may be attributed to the formation of SEI and the irreversible conversion of SiO2 to Si, Li4SiO4, Li2Si2O5 and Li2O. In
the second cycle, the CE is still low, ranging from 50.3% to67.9%, most likely due to additional conversion of SiO2 in thepotentiostatic step. In the subsequent cycling, the CE increases
rapidly, stabilizing around 99% aer 10 cycles. It should also be
noted that the cells subjected to a potentiostatic step havea signicantly higher CE aer 10 cycles compared to the reference cell not subjected to any potentiostatic step (SiO2/P_NPS).
A possible explanation is that the potentiostatic step ensuresthat most of the SiO2 available for conversion indeed is converted in the rst cycles, while the conversion process continues
incrementally over tenths of cycles in SiO2/P_NPS. If SiO2
anodes are to be used in full cells, this is a crucial observation,
as continuous depletion of Li+ in a full cell will rapidly deplete
the reservoir of Li+, rendering the cell useless. In contrast, thepotentiostatic step allows for the possibility of developing a pre-This journal is © The Royal Society of Chemistry 2019

RSC Advanceslithiation process, which can fully convert SiO2 prior to full cell
assembly.
The evolution of the potential prole over cycling for SiO2/P
can be seen in Fig. 3(d). The gure clearly displays how thepotentiostatic hold step in the second cycle, results in an
emerging plateau at approximately 0.2 V in the third cycle.
Corresponding gures for the other cells can be found Fig. S3.†
Fig. 3(e) displays the current going through the cells during the
aforementioned potentiostatic step, and the integral of thesecurves represents the charge passed. As seen from Fig. 3(e),although the initial currents are higher for the carbon-coatedsamples, a higher total charge is passed during the activation
step for the uncoated electrodes, and slightly higher for theconventional electrolyte compared to that with FEC and VC. The
increased charge going through the cells with pristine SiO2 isalso reected in Fig. 3(c), showing lower CE in the 2nd cycle forthe cells with pristine SiO2 compared to the carbon coated SiO2.
Further insight into the eﬀect of carbon coating, electrolyte
additives and the potentiostatic step can be gained from thediﬀerential capacity plots of the cells in Fig. 4. All cells in thisgure have the same mass. In the rst discharge cycle (Fig. 4(a)),
electrolyte reduction peaks can be seen for all cells. In agreement with the literature, cells with FEC and VC have electrolyte
reduction peaks at higher potentials, ranging between 1.4 and
1.5 V vs. Li/Li+, which appears to limit the SEI formation at 0.8 V
vs. Li/Li+ associated with the reduction of EC and DEC,15,44,45 as
may be inferred by the reduced size of these peaks. By
comparing the peaks in the anodic and cathodic direction of therst cycle (Fig. 4(a)) and third cycle (Fig. 4(b)), formation of new
peaks can be observed aer the potentiostatic step in cycle 2. In
the subsequent cycles (Fig. 4(c)), these peaks continue todevelop, resulting in peaks at approximately 0.29 and 0.48 V vs.
Li/Li+ in the anodic direction and 0.20 and 0.03 vs. Li/Li+ in thecathodic direction. A close-up of the rst peak in the cathodic
direction can be seen in Fig. 4(d). These peaks correspond quite
well to the known lithiation/delithiation potentials of Si in theliterature.46 For all cells, the peak intensities in the diﬀerential
capacity plot increase over the following cycles, with the highest
increase observed for the pristine sample in the conventional
electrolyte, in line with the observed increase in capacity(Fig. 3(a)). Moreover, the peaks are found at the same position
for each of these cycles, which is also in line with the highreversibility observed during cycling and the high CE observedaer the rst cycles.
Fig. 5 shows the SEM micrographs of the electrode surfacesaer 50 cycles, and Fig. 6 shows the FIB cross-section of thesame electrodes. The micrographs clearly show the deposits
resulting from the SEI formation on the samples. However, theamount and morphology of the SEI appears to be dependent on
both the use of carbon coating as well as the FEC and VC
additives. At this point it should also be emphasized thatuncycled carbon coated frustules were analyzed in the SEM, andthe carbon coating was not visually evident. This suggests that
all the deposits observed in Fig. 5 and 6 are SEI formation andnot the carbon coating. Fig. 5(a)–(c) show electrodes with pristine SiO2. In this case (Fig. 5(a)), some of the SiO2 particles show
now sign of SEI formation. Interestingly, the images of the SEI-RSC Adv., 2019, 9, 41228–41239 | 41233

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free SiO2 indicates that the microstructure of the frustules is
intact aer cycling. To investigate the microstructure of SiO2
particles with substantial SEI formation, the electron beam wasused to remove the layer. Fig. S4† demonstrate that the microstructure indeed is intact under the SEI layer. By comparing theindicated SiO2 particles in Fig. 5(b) and (c), a slight diﬀerence in
the SEI morphology, depending on the electrolyte, is observed.
The SEI on the electrode cycled with the electrolyte additives
included (Fig. 5(c)) appears to have a rougher surface consistingof spherical particles, while the electrode cycled in the additivefree electrolyte (Fig. 5(b)) has a more uniform surface. In thecase of electrodes made with carbon coated SiO2, the same
diﬀerence in SEI morphology can be observed in cells with
(Fig. 5(e)) and without (Fig. 5(c)) the additives. However, in bothcases, the carbon coating appears to increase the thickness of
the SEI layer signicantly. The increased SEI formation on
carbon coated SiO2 can also be observed within the electrode in
the FIB cross-section micrographs in Fig. 6. Comparing theelectrodes with pristine SiO2 in Fig. 6(b) with the electrodes withcarbon coated SiO2 in Fig. 6(d), the structure of the pristine SiO2
particles can be easily observed, while the increased SEI
formation makes it diﬃcult to spot individual carbon coated
SiO2 particle. Moreover, the overall porosity of the electrodeappears to be better preserved in the electrodes cycled in thePaperelectrolyte containing additives (Fig. 6(c) and (d)), compared tothe additive-free electrolyte (Fig. 6(b) and (e)). The observationthat the CE for pristine SiO2 is lower than that of the carbon
coated SiO2, and that less deposition products can be observedon the pristine SiO2, suggest that the conversion of SiO2 could
be higher in electrodes with pristine SiO2. This hypothesis isfurther strengthened by the data from galvanostatic cycling in
Fig. 3(a), where SiO2/P exhibits the highest capacity aer 50
cycles. Moreover, the SEM images might provide insight into
the eﬀect of FEC and VC as electrolyte additives. In cells withpristine SiO2, FEC and VC appears to limit the performance ofthe cells, while the opposite is true for carbon coated SiO2. One
possible hypothesis is that the initial decomposition of FEC and
VC observed in Fig. 4(a) increases the thickness and/or densityof the SEI formed on the pristine SiO2, which again might
hamper the conversion of SiO2. On the other hand, carbon
coated SiO2 shows quite extensive SEI formation in Fig. 5(d) and(e). In this case, the initial decomposition of FEC and VC might
limit the overall SEI formation. This hypothesis is supported by
the diﬀerential capacity plots in Fig. 4(a), which clearlyshow that the initial reduction peak of FEC at approximately
1.45 V vs. Li/Li+ suppress the reduction peak of EC/DEC at
0.8 V vs. Li/Li+.

Fig. 4 Diﬀerential capacity plots of all cells in (a) cycle 1, (b) cycle 3, (c) cycle 3–50 and (d) close-up of ﬁrst cathodic peak in (c).

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The good electrochemical performance of the diatom frustules is most likely related to the specic porous structure of the
diatom frustules, expected to be benecial to the anode
performance. The pores might ease the electrolyte penetration
into the electrode and improve the accessibility of the electrolyte and the reacting Li+ ions towards the active material. In
addition, the structure might also buﬀer any expansion oflithiated Si during cycling and thereby improve the cyclingstability. This is actually observed in Fig. 5, where the SEM
micrographs show very little distortion of the superstructure forthe cycled electrodes. Similar ndings have also been reported

RSC Advancesin the literature, for instance for amorphous SiO2/C composite
anodes.21 In this work, the good electrochemical performance
was attributed to the porous structural characteristics of the
SiO2/C material, allowing more electrolyte to be absorbed into
active materials, and also contributing to good rate-capability
and cycling stability of the electrode. Similar conclusions weremade in ref. 36, where the low charge transfer resistance of
C@Si–SiO2 electrodes was attributed to the mesoporous structure, resulting in good electrolyte–electrode contact.
In an eﬀort to identify which of the three proposed conversion reactions of SiO2 is taking place, the specic capacity gain

Fig. 5 SEM micrographs of SiO2 particles on; (a) and (b) SiO2/P, (c) SiO2/P_FEC-VC, (d) SiO2/C and (e) SiO2/C_FEC-VC, after 50 cycles. Arrows in
red indicate SiO2 particles.

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RSC Advances(QG), given by the increase in capacity for a set of cycles, wasplotted against the irreversible capacity loss (ICL), given by thesum of the CE losses for the same set of cycles. This method isbased on the work of Lepoivre et al. and the data is plotted inFig. 7.31 In addition, Fig. 7 includes the theoretical capacity gainvs. irreversible capacity loss calculated for the conversion reactions (1) to (3), representing the maximum possible increase in
specic capacity possible for a certain ICL. These theoreticalvalues are found by calculating the theoretical specic capacity
for reaction (1) to (3) (equaling 335, 836 and 1673 mA h g 1,
respectively) and the corresponding theoretical ICL for total
conversion (equaling 1784 mA h g 1 for all the reactions). It

Papershould be noted that experimental results will always include
some contribution from irreversible losses due to side reactions, thus the experimental QG values will always be lower than
the theoretical limit for a given ICL. The experimental valuesplotted as dots in Fig. 7, represent the increase in capacity fromcycle 2 to cycle 3 (hollow), as well as over cycle 3 to cycle 50(lled), plotted against the ICL for the same cycles. Cycle 1 was
excluded, as a large fraction of the ICL loss in this cycle may be
assumed to be caused by SEI formation. In Fig. 7, all the points
representing QG for the cycle 2 to 3 are below the theoretical lineof conversion reaction (2), but above the line of conversion
reaction (3). It is therefore reasonable to believe that theFig. 6 SEM micrographs of FIB cross-sections; (a) uncycled SiO2, (b) SiO2/P, (c) SiO2/P_FEC-VC, (d) SiO2/C and (e) SiO2/C_FEC-VC. Arrows in
red indicate SiO2 particles.

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material upon cycling, with a gain in capacity compatible witha high fraction of material converted to Si and Li2O. Aer the
initial formation cycles a high coulombic eﬃciency wasobserved, corresponding to low formation of electrolyte
decomposition products as compared to the carbon coatedmaterial.

Conﬂicts of interest
We declared that there is no conict of interests and all thelisted authors fully agree with publishing the manuscript.

Acknowledgements
Fig. 7 Increase in speciﬁc capacity (QG) as a function of ICL. The threelines (yellow, red and blue) show the theoretical capacity gain asa function of ICL for the three proposed lithiation mechanisms of SiO2.
The scatter dots show the measured increase in capacity as a functionof the ICL over the electrochemical reduction step, cycle 2 to cycle 3(hollow), and over cycle 3 to cycle 50 (ﬁlled).

capacity gain during the potentiostatic step is mainly driven byreaction (2). From cycle 3–50, QG for the cells with pristine SiO2
is well above the theoretical limit of reaction (2), implying
a signicant contribution from reaction (3). In the case of cellswith carbon coated SiO2, virtually no capacity increase is seen incycle 3–50. Thus, it is likely that most of the ICL is caused by SEI
formation. SEM micrographs of the electrode surface aer
cycling (Fig. 5(d) and (e)) showing decompositions products on
the surface, also support this.
In previous work on SiO2 frustules as an anode material,
where stable cycling (521 mA h g 1 for 20 cycles) was demonstrated with carbonized diatom algae, Li2O could not be
detected on the surface aer cycling.38 Thus, indicating that theconversion most probably occurs via a lithium silicate, implying
a lower theoretical capacity.

4. Conclusions
In this work, it is demonstrated that diatom frustules from theseabed, subject only to a simple cleansing and milling procedure, exhibit good performance as anodes for Li-ion batteries.
The best performance was obtained aer an initial activation
procedure, implying a potential hold step at 2 mV for 48 h. The
highest capacity (723 mA h g 1) was obtained with composite
electrodes with diatom frustules and conventional carbon black
conductive additive, with the capacity still increasing aer 50
cycles. The capacity was around 624 mA h g 1 aer subtractionof the contributions from the carbon black. Carbon coateddiatom frustules showed a slightly lower and stable capacityaer 50 cycles (679 mA h g 1, 600 mA h g 1 aer subtraction ofcontributions from the carbon coating and the carbon black).
Addition of FEC to the electrolyte improved the performance ofthe carbon coated frustules, but not the pristine material. The
good performance was attributed to an eﬃcient activation of theThis journal is © The Royal Society of Chemistry 2019

The authors gratefully acknowledge Planktonics AS for
supplying the diatoms used in this project. The ResearchCouncil of Norway is acknowledged for the support to theNorwegian Micro- and Nano-Fabrication Facility, NorFab,
project number 245963/F50. The Norwegian University of
Science and Technology is acknowledged for funding the work.

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