Influence of the mediating behaviour of Sn according to its particle size on a Ni/yttria-stabilised zirconia porous anode structure in a direct carbon fuel cell

Abstract

Direct carbon fuel cells (DCFCs) are devices that convert the chemical energy of a carbon fuel into electrical energy with high efficiency. Although they can theoretically yield a high efficiency, DCFCs have considerable polarisation resistance, resulting from a solid-state interface between the fuel and anode surface. Here, we have employed Sn nanoparticles (<150 nm in diameter) as an interfacial mediator to reduce the polarisation resistance at the interface, through a liquid phase Sn interlayer which can improve both the mass transfer of the solid carbon fuel and electron transfer. We previously found that, when using Sn microparticles (>150 μm in diameter), an interfacial resistance became alleviated with the aid of liquid-state Sn; however, the Sn microparticles tended to accumulate on the anode surface because the Sn microparticles hardly permeated into the submicron anode structure. As a result, we chose Sn nanoparticles to demonstrate the effect of the size of Sn on mediation of the Ni/YSZ (Ni/Yttria-Stabilised Zirconia) surface in this study. Unlike microparticles, the nanoparticles were observed to more deeply permeate into the Ni/YSZ pores without accumulation. Nevertheless, Sn nanoparticles are more prone to oxidation, resulting in an ionically and electronically insulating material. Therefore, the power performance of Sn nanoparticle-mounted DCFCs is, in general, lower than that of Sn microparticle-mounted DCFCs. In addition, we found that there are two possible mediation mechanisms: one is physical mediation by improving both the mass transfer and electron transfer, and the other is a Sn redox looping cycle.

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Introduction

The number of studies on renewable energy systems has grown rapidly due to concerns about the environmental impact of conventional energy sources, e.g. fossil fuels, and their exhaustion.¹ Fuel cells are a part of the recent movement to replace traditional power generation systems and were first pioneered by William R. Grove in 1838.² Fuel cells have been developed, thanks to their benefits including high efficiency and continuous operation. Amongst the various classifications of fuel cells, the use of a direct carbon fuel cell (DCFC) has been suggested as a promising approach and they have advanced considerably in recent years.^3–6 In fact, the concept of a DCFC was firstly patented by William W. Jacques in 1896.⁷ A DCFC is a device that converts the chemical energy of the carbon fuel into electrical energy. Theoretically, DCFCs have superior efficiency compared to other energy generating systems because of their direct energy conversion without intermediate steps.⁵ Moreover, they are not subject to Carnot limitations.⁸ The following equations (eqn (1)–(6)) represent the basic mechanism of DCFCs. It has been recently identified that CO molecules play a dominant role in the DCFC operation.⁹ CO can be generated not only from an electrochemical reaction (eqn (3)) but also from the reverse Boudouard reaction (eqn (5); chemical reaction). Consequently, CO can be electrochemically oxidized into CO₂ with generating electrons (eqn (4)). In addition, since the exhaust gas mainly contains CO and/or CO₂, the exhaust gas can be oxidized into CO₂ through a simple process and then easily captured. This high quality CO₂ can be taken advantage of for CO₂-requiring systems, such as chemical synthesis.¹⁰

Overall: C + O₂ → CO₂ (1)

Anode: C + 2O²⁻ → CO₂ + 4e⁻ (2)

Anode: C + O²⁻ → CO + 2e⁻ (3)

Anode: CO + O²⁻ → CO₂ + 2e⁻ (4)

Anode: C + CO₂ ⇄ 2CO (5)

Cathode: O₂ + 4e⁻ → 2O²⁻ (6)

There are three main DCFC subcategories according to the type of electrolyte material:¹¹ molten hydroxide, molten carbonate, and solid oxide. Amongst these, solid oxide-based DCFCs (SO-DCFCs) have been widely studied due to their simplicity and fuel flexibility.^12–20 They have a broad spectrum of fuel employment; in other words, any carbonaceous materials can be utilised no matter how high a carbon purity the material has, i.e., from carbon black or coke to coal—and even biomass.^16–26 Even though SO-DCFCs are beneficial in many aspects, still there is a big issue to overcome because both the fuel and anode are solid-state materials. In a SO-DCFC, electrochemical reactions can occur at triple phase boundaries (TPBs) at which the current conducting material, ionically conducting electrolyte and fuel simultaneously meet. All these components are solid-state materials and, especially, the anode is made from a composite of the current conducting material and ionically conducting electrolyte. Thus, the contact between the fuel and the anode surface is a matter of concern.

In a DCFC operation, high polarisation resistance at the fuel/anode interface is observable, since both the fuel and anode are solid-state materials;²⁷ therefore, the anodic half reactions (eqn (2) and (3)) are sluggish.²⁸ Some researchers have tried to facilitate an effective interface by increasing the TPB sites.²⁹ This has been well studied and it was found that the TPB sites could be no longer extended beyond 10 μm even if a gas-state fuel was employed.^30–37 To increase the number of TPB sites, a liquid metal was introduced into the anode site.^28,38–52 Among the various liquid metal candidates, Sn has been widely studied because of its low melting temperature, lesser toxicity, and resistance to ash and/or sulfur poisoning.^{28,41–44,46,48,50,51,53} The applications of Sn span from low temperature fuel cells, as a favoured electrocatalyst for organic fuel utilization,^54,55 to high temperature fuel cells including DCFCs. In high temperature fuel cells, the role of Sn is not limited to an anodic substitute or liquid anode; that is, it can also work as an additive,⁶ dopant,⁵⁶ or catalyst.⁵⁷ One of the other approaches using Sn is to make a Ni/Sn alloy which is less vulnerable to carbon deposition than Ni, which is the most commonly utilised solid oxide fuel cell (SOFC) anode catalyst.^58–60 In addition, we previously demonstrated the role of Sn as an interfacial mediator between a solid-state fuel and anode.¹⁵

As an interfacial mediator, Sn improves the contact between the fuel and anode, because Sn becomes a liquid-state component at DCFC working temperatures. Therefore, mass transfer and/or electron transfer can occur within the liquid-state Sn, and then, the TPBs can become extended.¹⁵ Our previous study utilised Sn microparticles (>150 μm in diameter, Table 1), and they were not able to deeply penetrate into the anodic pores due to their large dimensions, causing accumulation over the anode surface. In other words, the Sn microparticles become redundant since they fully covered the anode surface. We observed that 15 mg of Sn microparticles fully covered the anode surface (3.14 cm² of geometric surface area). When less than 15 mg was used, this was insufficient to cover the surface; when more than 15 mg was used, the excess amount of Sn became accumulated over the surface, leading to a decrease in the power performance. Therefore, we concluded that 15 mg of Sn microparticles was an optimal quantity for DCFC operation. In this contribution, on the contrary, we have introduced Sn nanoparticles (<150 nm in diameter) into fuel/anode interfaces to demonstrate the Sn behaviour according to particle size in a DCFC operation.

Table 1 Summary of our previous¹⁵ and present work at 850 °C

Carbon (Sn quantity)^a/mg		300 (0)	300 (15)	300 (30)	300 (60)	0 (300)
a In the case of the previous work, the Sn quantity consists of microparticles; whilst it consists of nanoparticles for the present work. Note: the carbon quantity does not include the Sn quantity, i.e., the total mass of 300 (60) is 360 mg, not 300 mg.
Maximum power density/mW cm^−2	Previous work	67	136	133	93	42
	Present work	—	97	102	115	38

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Experimental section

Pre-experimental testing details and preparations

Carbon black powder (Ensaco 350G, Timcal) was selected as the fuel due to its high carbon purity. Commercially available anode-supported button-type cells (Ceramic fuel cell power, South Korea) were employed for the fuel cell tests, comprising a Ni/YSZ (Ni/Yttria-Stabilised Zirconia) porous anode, 8YSZ electrolyte and LSM (Lanthanum Strontium Manganite) cathode. The porosity of the anode was measured using mercury intrusion porosimetry and it was 28.4%. A Pt (99.9%, 52 mesh, Alfa Aesar) current collector was attached to the cell with Ag paste (Fujikura Kasei). Later, the cell was gently air-sealed with a sealant (Thermiculite 866).

Sn nanoparticles (<150 nm, Aldrich) were evenly sprayed onto the anode surface. Unlike our previous work wherein a brush was utilized to apply the Sn,¹⁵ a spraying method was selected to provide pneumatic pressure, through the nozzle, to facilitate Sn permeation and to evenly distribute the Sn on the Ni/YSZ surface. Respectively, 15, 30 or 60 mg of Sn nanoparticles was distributed over a 3.14 cm² anode surface (geometric surface area). Subsequently, 300 mg of carbon black was mounted onto the Sn-covered surface. An additional reference experiment was also carried out, using 300 mg of Sn nanoparticles without carbon. Afterwards, the cell was connected to optimally designed apparatus (NARA Cell-Tech) – consisting of an alumina ceramic reactor, a furnace and an electrochemical workstation.

Fuel cell operation and post-experimental testing

Prior to a power performance test, the cell was air-sealed once more using ceramic bond (Aremco ceramabond 668). The apparatus was programmed to elevate the temperature at a ramping rate of 5 °C min⁻¹, and to hold the temperature for a few minutes for the power performance test. Power performance tests were carried out at 700 °C, 750 °C, 800 °C and 850 °C. Ar gas was continuously supplied to the anode chamber from start-up to 600 °C at a rate of 30 mL min⁻¹ to purge out any atmosphere-derived gas (N₂, O₂, CO₂) within the chamber. For the cathode chamber, on the other hand, O₂ gas was always supplied at a rate of 50 mL min⁻¹, as an O²⁻ precursor. Since a sudden temperature drop could affect the cell resistance,^61,62 all the gas inlets were heated and maintained at 100 °C.

Model experiments using a Si wafer were carried out under the same experimental conditions as detailed above – 15, 30 or 60 mg of Sn nanoparticles with temperature elevation. The purpose of the model experiments was to understand the internal behaviour of the Sn nanoparticles within the anode chamber, since in situ observation during DCFC operation is difficult due to the high operation temperature. Subsequently, the model systems were analysed after the experiment with a field emission scanning electron microscope (FE-SEM; S-4700, Hitachi, Japan) to obtain the morphological features. In addition, thermogravimetric analysis (TGA-50H, Shimadzu) was performed to clearly identify the effect of the size of the Sn at experimental temperatures.

Results and discussion

Fuel cell performance

Direct carbon electrochemical oxidation reactions (C + nO²⁻ → CO_n + 2ne⁻, n = 1, 2; eqn (2)–(4)) are sluggish in DCFCs because both the fuel and anode are solid-state materials. To overcome the sluggish kinetics, Sn nanoparticles were employed as a liquid-state interfacial mediator between the solid-state materials (fuel/anode). Current density–potential (j–V) and current density–power (j–P) tests were carried out with and without the Sn nanoparticles. The resultant curves are plotted in Fig. 1, showing the following maximum power densities (MPDs) at 850 °C: 97, 102 and 115 mW cm⁻², obtained with 15, 30 and 60 mg of Sn nanoparticles respectively. The results were additionally tabulated, Table 1, and compared to our previous work with Sn microparticles. Our present results show that an increase in the Sn quantity leads to enhancement of the power performance. However, this is an entirely inverse trend compared to our previous work:¹⁵ the previous work exhibited that 15 mg was optimal, followed by 30 mg and 60 mg. In our previous study,¹⁵ the decrease in the power performance was due to over-growth and accumulation of the excess quantity of Sn over the anode porous surface; thus permeation into the pores became hindered and therefore there were less mediated TPB sites. As shown in Table 1, the performance values for the present work are, in general, slightly lower than those for the previous work.

Fig. 1 Current density–potential (j–V) and current density–power (j–P) curves at different temperatures: (a) 700 °C, (b) 750 °C, (c) 800 °C and (d) 850 °C. Each symbol in the graphs represents a different quantity of Sn nanoparticles: 15 mg (● green circle), 30 mg (▲ blue triangle), 60 mg (▼ red upside-down triangle) and Sn nanoparticles only (■ black square). Open symbols designate cell voltage values (left-hand side ordinate) and closed symbols power density values (right-hand side).

Feasibility of Sn permeation according to its particle size

We obtained microscopic images (Fig. 2) to investigate the size-effect of Sn. In Fig. 2a and b, the morphologies are analogous each other, which were respectively obtained with 15 mg and 60 mg of Sn nanoparticles. In other words, 60 mg of Sn nanoparticles can readily permeate throughout the porous anode surface without being accumulated over the surface. Our previous work¹⁵ with Sn microparticles, on the contrary, showed severe over-growth and accumulation when 60 mg of Sn was employed because it was not able to deeply permeate. The pristine cell was magnified and the images are exhibited in Fig. 2c and d. Fig. 2e is an image of the Sn nanoparticles, whilst Fig. 2f is of Sn microparticles. The particle sizes of the Sn particles are tabulated in Table 2, and the pore size of our Ni/YSZ anode structure is in the submicron range. Fig. 2c and e are at the same magnification, implying that the Sn nanoparticles are tiny enough to easily permeate into the anode pores. In addition, Sn microparticles are too large to permeate into the anode pores. Moreover, the size of the Sn microparticles makes it difficult for them to pass through a submicron anode structure. In other words, the Sn nanoparticles could relatively easily permeate into deeper pores, building more mediated TPB sites, and hence enhancing the power performance.

Fig. 2 SEM images of the Ni/YSZ porous anode surface and Sn particles: (a) anode surface after an experiment using 15 mg of Sn nanoparticles and (b) using 60 mg; (c) pristine cell (×10 000 of magnification); (d) pristine cell (×1000); (e) Sn nanoparticles, (f) Sn microparticles. Note: images (c) and (e) have a different magnification in order to facilitate comparison; herein, a pristine cell means a cell being not ever used, namely, a cell before being mounted with Sn powders or carbon fuel.

Table 2 Morphological information for the Sn particles, obtained from SEM observation

	Particle size distribution	Average particle size	Particle morphology
Sn nanoparticles	<150 nm	130 (±68) nm	Sphere
Sn microparticles	1.0–32.7 μm	4.1 (±3.2) μm	Ellipsoid–oblong

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It seems that the permeation of molten Sn is more demanding than for solid-state Sn particles. Tao et al. suggested that liquid Sn permeation throughout porous ceramic materials is restricted due to its high surface tension (>400 dyn cm⁻¹) and reduced wettability with ceramics (>90°).²⁸ In other words, once the Sn melts, further permeation is difficult, compared to the permeation of solid-state Sn particles. In addition, it has been reported that the initial size of the solid-state Sn might influence the permeation behaviour, implying that nanoparticles relatively easily permeate.⁵³ These behaviours provide the reason why the initial particle size is important and why power performance becomes enhanced along with an increase in the quantity of Sn nanoparticles. Namely, Sn nanoparticles more facilely permeate into anode pores than Sn microparticles. In Fig. 3, the depth of the Sn particle permeation into the pores is presented along with additional inset schematic diagrams and results for the model experiments. The model experiments provide information about the height of the Sn nanoparticles when sprayed onto a Si wafer. When using 15 mg of Sn nanoparticles, as shown in Fig. 3a, the permeation depth was ca. 15 μm; whilst ca. 50 μm was observed for 60 mg (Fig. 3b). The model experiments (Fig. 3 right) show a good agreement with the depth results (Fig. 3 left). In addition, the depth results show how well the Sn nanoparticles permeate into anode pores.

Fig. 3 SEM cross-sectional images, showing (left) Sn permeation into the porous Ni/YSZ anode microstructure and, (right) the height of Sn when sprayed onto a Si wafer. The different amounts of Sn particles were evenly applied using a spray onto the 3.14 cm² anode surface: i.e., (a) 15 mg and (b) 60 mg, respectively. The insets are schematic diagrams representing the depth of Sn permeation (dark grey) into pores (white) within the Ni/YSZ anode (green).

Model experiment for Sn distribution

A model experiment was carried out to identify the agglomeration behaviour of the Sn nanoparticles. Fig. 4 shows that the surface was sufficiently covered with Sn nanoparticles even though some of the particles became overgrown and agglomerated. In our previous study, on the contrary, the Sn microparticles tended to only overgrow and tended not to cover the surface,¹⁵ implying that Sn permeation with microparticles is considerably hindered. In other words, once the Sn nanoparticles have permeated gently into the porous structure, the Sn nanoparticles would be well distributed over the anode surface without Sn accumulation, as shown in Fig. 2 and 3. Since Sn becomes a liquid-state material during fuel cell operation at a high temperature, liquid Sn could cover the YSZ sites where Ni is absent and, at these sites, the Sn could work as an electron conducting material, instead of Ni. In other words, when Sn nanoparticles are spread deeper into the pores, this could result in broader TPB sites. In addition, since the Sn is liquid, the solid/solid interfaces could be mediated by the liquid Sn layer and therefore the mass transfer of the solid fuel could become enhanced. As a result, the power performance becomes improved along with an increase in the Sn nanoparticle quantity (Fig. 1).

Fig. 4 SEM images of (a) the surface morphology and (b) a cross-sectional view, after a model experiment using a Si wafer, where 60 mg of nano-Sn was sprayed onto a 3.14 cm² geometric surface area. The results are distinct from those with micro-Sn;¹⁵ in that, Sn overgrowth is not predominant and therefore the surface is almost fully covered with Sn.

Sn oxidisability according to the particle size

The overall power performance in the present work with Sn nanoparticles is slightly lower than for the previous work with Sn microparticles (Table 1). This seems to be due to a sensitive oxidation behaviour of Sn nanoparticles, compared to that of Sn microparticles, since the surface area per unit mass of the nanoparticles is much greater than that of the microparticles.⁶³ In addition, Sn nanoparticles are known for their prompt oxidation to SnO and further oxidation to SnO₂.^64,65 Sn oxides are insulating materials and therefore decrease the TPBs. Sn oxidation could occur because O²⁻ ions are provided at the TPBs through the YSZ electrolyte at elevated temperatures.⁶⁶ The Sn particle size affects the initiation temperature of Sn melting and oxidation (Fig. 5 and Table 3^67,68), and smaller particles tend to react at even lower temperatures. In other words, Sn nanoparticles are more prone to oxidation into Sn oxides than Sn microparticles, and therefore the power performance of the Sn nanoparticle-mounted DCFC is, in general, lower than that of the Sn microparticle-mounted DCFC.

Fig. 5 Thermogravimetric analysis of various Sn particles. The weight of the Sn particles becomes heavier as they are oxidised into Sn oxides. Particles with a larger size tend to start being oxidised later.

Table 3 Detailed information for different Sn–O species^67,68

	Molar mass/g mol^−1	Ratio^a	Melting point/°C
a Ratio of mass of a certain species (Sn, SnO, SnO2) to Sn.
Sn	118.71	1	231.93 (bulk particles)
			224.4 (for 29.1 nm particles)
			203.1 (for 14.5 nm particles)
			117.3 (for 11.3 nm particles)
SnO	134.71	1.135	1080
SnO2	150.71	1.270	1630

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Evidence of Sn being a liquid metal

Even though Sn oxides could be formed during the DCFC operation (eqn (7)) and therefore cause decreased power performance, they could be reduced back into liquid-state Sn, which mediates the fuel/anode interface. For example, it has been reported in metallurgy that SnO₂ can be carbothermally reduced at an elevated temperature⁶⁹ – as well as by CO.⁴⁶ In addition, it was recently identified that CO is dominantly present within the anode chamber during a SO-DCFC operation⁹ and, since gaseous species have more mobility than solid carbon particles, it seems that the possibly formed SnO₂ can be reduced by CO, as well. The following carbothermal reduction mechanisms (including reduction by CO) are the most widely known reactions.

Sn + 2O²⁻ → SnO₂ + 4e⁻ (7)

SnO₂ + 2C ⇄ Sn + 2CO (8)

SnO₂ + C ⇄ Sn + CO₂ (9)

SnO₂ + 2CO ⇄ Sn + 2CO₂ (10)

In a temperature region higher than 641 °C, the free energy values of equations (eqn (8)–(10)) are all negative,⁷⁰ and therefore these above reactions would occur spontaneously during a DCFC operation. In other words, even though Sn oxides can be formed, they can be reduced into liquid-state Sn during the DCFC operation, once the anode chamber adequately lacks oxidants. Fig. 6 supports that Sn exists in a reduced state. The turnover period for 300 mg of Sn to be fully oxidised into Sn oxides is less than ca. 1 h;¹⁵ nevertheless, the discharge tests lasted several hours, Fig. 6, implying that not all the Sn particles are oxidised and carbothermal reduction may happen. In addition, a battery mode (the peak of the red line) was observed when using 60 mg of Sn nanoparticles, suggesting that a considerable amount of Sn remains reduced until the battery mode begins.

Fig. 6 Discharge tests for the experiments with Sn nanoparticles at a constant voltage of 0.4 V. The calculated total charges for 15 mg of Sn (solid black line) and 60 mg of Sn (solid red line) are, respectively, 1038.5 mA h and 884.5 mA h.

Overall discussion and suggested mechanisms

So far, we have discussed why the power performance of a Sn nanoparticle-mounted SO-DCFC differs from that of a Sn microparticle-mounted SO-DCFC. To sum up, the permeation, mediation action and oxidation of Sn are key steps that determine the power performance. First, in terms of permeation, Sn microparticles hardly permeate into the Ni/YSZ pores because the permeation depth of the Sn particles is determined by the initial particle size and, once the particles become molten, further permeation is difficult. On the contrary, Sn nanoparticles easily permeate into the Ni/YSZ pores at room temperature and therefore an increase in the amount of mounted Sn enhances the electrochemical reaction sites, accordingly. This different permeation behaviour is illustrated in Scheme 1. In addition, when molten, the Sn microparticles form a thicker Sn layer, e.g. Scheme 2b-1, and diffusion of the reactants through the Sn layer becomes rather difficult. This behaviour occurred in our previous work: the Sn layer being much thicker and having an overgrown Sn layer led to a decrease in the power performance (Table 1). However, in the case of Sn nanoparticles, as shown in Fig. 2a and b, they can easily permeate into the anode pores. Also, they can be well distributed, thinly, thanks to a spray method, as shown in Fig. 4. Namely, as reflected in Scheme 2c-1, the carbon particles can take advantage of the enhanced TPBs.

Scheme 1 Overall comparison between Sn microparticles and nanoparticles according to the particle size, and the effect of increasing the applied quantity on the temperature regions of (upper) the Sn permeation and (lower) the fuel cell operation steps. In this study, we employed a cell, the anode part of which was Ni/YSZ, that had a submicron porous structure.

Scheme 2 Overall mechanism occurring within an anode chamber when a Sn interlayer is introduced. (a) General features of the anode surface. (a-1) Anode/fuel interface with Sn particles in the middle, i.e. before temperature elevation; (a-2) molten Sn layer without the carbon fuel in order to facilitate reading, i.e. during fuel cell operation in elevated temperature regions. (b) Oxidation and carbothermal reduction mechanisms of Sn. (b-1) Sn gets oxidised by oxide ions when reactants such as C and CO are not reaching the reaction sites, cf. this is an electrochemical reaction and generates electrons; (b-2) SnO₂ gets reduced by carbon particles; (b-3) SnO₂ gets reduced by CO molecules, cf. the CO generation mechanism is explained in (c); (b-4) resultant Sn recovery. (c) Mediation mechanism of Sn. (c-1) Direct electrochemical oxidation of carbon particles at a Sn-mediated Ni/YSZ interface, cf. as a result, CO and/or CO₂ is produced; (c-2) electrochemical oxidation of CO; (c-3) the reverse Boudouard reaction, cf. by looping together the (c-2) and (c-3) processes, both the CO and CO₂ partial pressure within the anode chamber can increase.

When it comes to mediation, a thin liquid Sn layer can mediate the fuel/anode interface and therefore improve the power performance. This is because the liquid Sn interlayer can enhance the mass transfer of the fuel and alleviate the high polarisation resistance of the solid/solid interface. The mediation occurs as shown in Scheme 2c-1, that is, liquid Sn provides broader TPBs because the liquid Sn can transfer both electrons and solid carbon particles. However, if the Sn layer is too thick, i.e. as in the case of the Sn microparticle-mounted DCFC, the mass transfer of the carbon particles becomes difficult due to a low wettability of the Sn in the cell. When the Sn layer is sufficiently thin and mediates the electrochemical reactions of the fuel, i.e. the advantage of using Sn nanoparticles, both CO and CO₂ can be generated through electrochemical reactions (eqn (2) and (3); Scheme 2c-1). Also, CO can be generated through a reverse Boudouard reaction (eqn (4); Scheme 2c-3). Once CO is present within the anode chamber, it can play a dual role in the SO-DCFC: one role is to generate electricity (Scheme 2c-2), and the other is to reduce SnO₂ back into Sn (Scheme 2b-3).

Last but not least, tin oxide formation affects the Sn mediation behaviour and hence decreases the power performance. As shown in Scheme 2b-1, when reactants such as C and CO do not reach the TPBs, the Sn can get oxidised into SnO₂, which prevents further electrochemical reactions as well as mass transfer and electron transfer. SnO₂ can be carbothermally recovered to form Sn (Scheme 2b-4) by carbon particles (Scheme 2b-2) and/or CO (Scheme 2b-3). However, the Sn redox cycle does not seem fully reversible. Even though both Sn microparticles and Sn nanoparticles can sufficiently cover the anode surface with an amount of 15 mg, the power performance values (Table 1) for the two cases are different. The gap between the two cases is primarily due to the fact that Sn nanoparticles are more vulnerable to oxidation (Fig. 5, 6 and Table 3). If reduction of the SnO₂ was fully obtainable, the gap would not be this great. In this sense, the Sn oxidation is not fully reversible and results in an insulating layer at the fuel/anode interface. However, it is still clear that the majority of the Sn exists in a metallic liquid state because the power performance with C + Sn was greater than that of C or Sn only (Table 1) and because the discharge tests lasted longer than the Sn oxidation turnover time (Fig. 6).

In fact, until now, we have focused on the physical mediating behaviour of the liquid Sn, i.e. formation of a liquid interlayer between the solid fuel and solid anode in order to improve both the mass transfer and electron transfer, which can alleviate polarisation resistance resulting from the solid-state interface. However, there is one more possible pathway to generate electrons in our system, that is, a Sn redox looping cycle: Scheme 2b-1 → Scheme 2b-2/Scheme 2b-3 → Scheme 2b-4 → Scheme 2b-1. In the case of a Sn redox looping cycle occurring in deeper pores, the role of CO seems more important than that of the carbon particles because solid carbon has less mobility than gas-phase CO and hardly reaches the reaction sites. The redox cycle for solid carbon particles still substantially occurs in areas beneath the solid carbon fuel. Even though, as we already discussed, this cycle is not fully reversible, it is still possible, and we believe the two pathways are both important. However, because of technical difficulties, in situ observation of DCFCs is still difficult and, if in situ techniques are developed for DCFCs, then we can ascertain which pathway (either physical mediation or an electrochemical looping cycle) plays a dominant role in DCFC operations.

Conclusion

We have investigated the interfacial mediation behaviour between a fuel and anode using Sn nanoparticles. To sum up, the permeation, mediation action and oxidation of Sn are key steps that determine the power performance. Our previous work using Sn microparticles¹⁵ showed that the best power performance is attainable when Sn sufficiently permeates into the anode pores and then covers the anode surface without overgrowth and accumulation. An overall comparison of the permeation behaviour of Sn microparticles and nanoparticles is schematically provided in Scheme 1. It exhibits that Sn microparticles hardly permeate into the Ni/YSZ pores, whilst Sn nanoparticles do so easily. Therefore, Sn nanoparticles can reach deeper into the pores and then form a thinner interlayer. However, the formation of a thinner interlayer is like a double-edged sword: it can result in better mediation behaviour due to the enhanced TPBs; but it is more vulnerable to oxidation into SnO₂ and the oxide layer acts as an insulator between the fuel and anode (Fig. 2, 5, Tables 2 and 3). SnO₂ reduction is available via the pathway in Scheme 2b. Nevertheless, the Sn redox cycle does not seem fully reversible, and therefore, unrecovered SnO₂ causes a decreased power performance. In this sense, the power performance of a Sn nanoparticle-mounted DCFC is, in general, lower than that of a Sn microparticle-mounted DCFC, Table 1. In addition, we have found that there are two possible mediation pathways in terms of the chemistry of Sn. One is physical mediation by improving both the mass transfer and electron transfer through the liquid Sn interlayer between the solid fuel and solid anode, because the liquid interlayer can alleviate polarisation resistance resulting from the solid-state interface (Scheme 2c). The other is a Sn redox looping cycle (Scheme 2b). However, because of technical difficulties in in situ analysis of DCFCs, we cannot ascertain which one plays a greater part in the electrochemical reactions, and we need to further develop DCFC in situ analytical techniques for clearer identification. Nevertheless, it is clear that the Sn interlayer mediates the electrochemical reactions and this behaviour can be taken advantage of for practical DCFC operation.

Acknowledgements

This work was supported by the GIST Research Institute (GRI) in 2016.

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