# Wear mechanism of oxide-bonded SiC refractory tiles in the combustion zone of waste-to-energy facilities

A. Tixier, C. Varona, J.M. Brossard, Patrick Ganster, François Valdivieso, J.

Poirier, Emmanuel de Bilbao

► To cite this version:

A. Tixier, C. Varona, J.M. Brossard, Patrick Ganster, François Valdivieso, et al. Wear mechanism of oxide-bonded SiC refractory tiles in the combustion zone of waste-to-energy facilities. Journal of the European Ceramic Society, 2025, 45 (3), pp.117016. 10.1016/j.jeurceramsoc.2024.117016. emse-05014092

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# Wear mechanism of oxide-bonded SiC refractory tiles in the combustion zone of waste-to-energy facilities

Audrey Tixier<sup>a,c,d</sup>, Cyrille Varona<sup>a</sup>, Jean-Michel Brossard<sup>b</sup>, Patrick Ganster<sup>c</sup>,
François Valdivieso<sup>c</sup>, Jacques Poirier<sup>d</sup>, Emmanuel de Bilbao<sup>d,*</sup>

## ARTICLE INFO

### Keywords:

SiC refractories

Waste-to-energy facilities

Oxidation

Corrosion

## ABSTRACT

Refractory tiles used in waste-to-energy facilities are exposed to high temperatures (T<sub>flue-gas</sub> ≈ 1200 °C) and corrosive environment with gases and fly ashes. Post-mortem examination of refractory tiles from the combustion zone of waste-to-energy plants was conducted to understand the damage mechanism of the refractory material. Corroded refractory tiles from this area showed an important loss of matter and large cavities on the hot surface. Investigations highlighted the weakness of the refractory matrix of the material. In contact with slag deposit and corrosive gases, the bonding phase of the refractory material is dissolved and SiC grains are oxidized which causes the spalling of the hot face of refractory tiles.

## 1. Introduction

High temperature (T<sub>flue-gas</sub> ≈ 1200 °C) and corrosive species such as gases, particles, or flying ashes from municipal solid waste combustion expose water wall boiler tubes of waste-to-energy facilities to a severe environment [1-4]. In these conditions, it is required to use refractory materials to protect the boiler while maintaining heat transfer from waste combustion. Thus, shaped silicon carbide tiles are commonly used for their good thermal properties and good abrasion resistance.

According to the height in the combustion chamber, different kinds of refractory lining wears can be distinguished [5,6] (Fig. 1):

- the top area (height > 12 m): the refractory lining is not damaged and very few fly ashes deposed on the tiles;

- the intermediate area (5 m < height < 12 m): fouling type deposit on tiles, volume expansion induced by the mineral phase change, and cracks of SiC refractory tiles;

- the combustion area (grate < height < 5 m): important amount of slag and loss of material of refractory lining.

SiC oxidation (Eq. 1) is one of the phenomena which leads to the wear of SiC-based refractory tiles due to the presence of O<sub>2</sub> and H<sub>2</sub>O in

the flue gas (around 9 vol% and around 16 vol% at the top of the combustion chamber [5]). Below 1000 °C, SiC oxidation kinetic is low [7]. However, SiC oxidation kinetic increases dramatically under water (H<sub>2</sub>O<sub>g</sub>) influence (Eq. 2) [8-12].

$$ \mathrm{SiC}_{(\mathrm{s})} + 3/2 \mathrm{O}_{2} (\mathrm{g}) \rightarrow \mathrm{SiO}_{2} (\mathrm{s}) + \mathrm{CO}_{(\mathrm{g})} \quad (1) $$

$$ \mathrm{SiC}_{(\mathrm{s})} + 3 \mathrm{H}_{2} \mathrm{O}_{(\mathrm{g})} \rightarrow \mathrm{SiO}_{2} (\mathrm{s}) + \mathrm{CO}_{(\mathrm{g})} + 3 \mathrm{H}_{2} (\mathrm{g}) \quad (2) $$

The diffusion of alkalis in the bonding phase of the refractories also contributes to the wear of all the refractory tiles. Penetration of alkali compounds can lead to the formation of low melting compounds [13,14] or also to the formation of leucite KAlSi<sub>2</sub>O<sub>6</sub>, nepheline Na<sub>3</sub>KAl<sub>4</sub>Si<sub>4</sub>O<sub>16</sub>, and β-alumina which are expansive phases. These phase transformations can cause cracks and spalling of the material [15]. Tomita et al. showed that NaCl and KCl promote silica crystallisation into cristobalite at low temperature [16] and basic molten salts can dissolve the protective silica layer from SiC oxidation [17].

Investigations of the failure mode of SiC refractory tiles, especially in the intermediate area where expansion and cracks on the tiles were noticed, show that alkaline and calcium sulphates can condense near the hot face of tiles [5,18]. Acid-base reactions occur between CaSO<sub>4</sub>, which is dissociated in Ca<sup>2+</sup> and SO<sub>4</sub><sup>2-</sup>, and free silica from the refractory matrix

<footnote><sup>a</sup> Bony SAS – Produits Réfractaires, Saint-Étienne 42000, France</footnote>

<footnote><sup>b</sup> Veolia Recherche et Innovation, Limay 78520, France</footnote>

<footnote><sup>c</sup> Mines Saint-Étienne, LGF CNRS UMR 5307, Saint-Étienne 42023, France</footnote>

<footnote><sup>d</sup> CNRS, CEMHTI UPR3079, Univ. Orléans, Orléans F-45071, France</footnote>

<footnote>* Correspondence to: 1D Avenue de la Recherche Scientifique, Orléans 45071, France.</footnote>

<footnote>E-mail address: emmanuel.de-bilbao@univ-orleans.fr (E. de Bilbao).</footnote>
to form parawollastonite CaSiO₃. Sulphation of KCl and NaCl in the refractory tiles can cause the formation of feldspar phases such as microcline KAlSi₃O₈, cristobalite SiO₂, and liquid phases in the material. In these conditions, expansion variation from silica crystallisation and the formation of expansive phases lead to the failure of the material [7, 19].

In the combustion zone of waste-to-energy plants, an important amount of slag covers the refractory tiles, and a significant loss of material is pointed out. In contact with slag deposit, the refractory lining reacts to form liquid phases and to promote the dissolution and diffusion of oxygen and contaminant elements on the surface of tiles [6]. Contrary to the intermediate area, the wear mechanism of the SiC-based refractory tiles is more related to the liquid slag.

The aim of this study is a better understanding of the wear and the failure mode mechanism for the oxide-bonded SiC refractory tiles located in the combustion zone of waste-to-energy facilities. Post-mortem characterisation, thermodynamic calculations, and laboratory corrosion experimentations were performed to propose a specific damage mechanism and to suggest some material improvements.

## 2. Material and methods

### 2.1. Refractory material

SiC-based refractory tiles used in waste-to-energy facilities contain around 90 wt% SiC bonded with an alumino-silicate matrix (Fig. 2). The chemical composition and main properties of the material are presented in Table 1. Due to the complex shape (important thickness variations, holes) and the way of shaping (uniaxial pressing) of the refractory tiles, the density distribution of the material is heterogeneous. This is illustrated by the density and pore size measurements determined on different spots in the tile. Fig. 3 shows the pore size distribution determined by intrusion mercury porosimetry on samples (10x10x20 mm³) taken in the cold face (sample 1) and in the hot face (sample 2) of the same tile. In sample 1, a bimodal distribution between 0.02 and 1 µm

and 10–100 µm is observed. Its open porosity is evaluated to 19 vol%. In sample 2, a monomodal distribution ranging from 1 and 10 µm is observed and its open porosity is evaluated to 15 vol%.

### 2.2. Experimental methods

Microstructure examinations were carried out by scanning electron microscopy (Philips XL40 and JEOL JSM-6500F) in the backscattering mode on samples beforehand polished (up to a diamond suspension of 1 µm) and coated with carbon. Chemical compositions were analysed by X-ray fluorescence (Shimadzu EDX 720) and energy-dispersive spectroscopy (Oxford Instruments and Princeton Gamma-Tech) coupled with SEM. Powder X-ray diffraction (pXRD, Bruker D8 Advance using CuKα radiation and Ni-filter) was conducted on samples ground in a tungsten carbide mortar (Fritsch Pulverisette) to study possible phase transformations. Pore size distribution was determined by mercury intrusion porosimetry (Micromeritics AutoPore IV 9500) and open porosity was measured by Archimedes method due to the limitation of the mercury penetration in the very small pores. The oxidation of SiC grains was measured by thermogravimetric analyses (Netzsch STA 409 CD).

The post-mortem examination was completed with lab experiments for a better understanding of the local phenomena between raw materials of the refractory material and slag deposit. To study the behaviour of raw materials in contact with slag deposit, pellets (Ø 20 mm, 6 mm) containing 50 wt% of real slag deposit from the plant designated as “plant B” and 50 wt% of refractory raw material were shaped with a hand press into a metallic mold. For the tests, two-grain size distributions of silicon carbide (Traxys, 95 % purity, 1–3 mm, and 0–250 µm) and bentonite clay (Ankerpoort Portaclay A90) were used as refractory raw materials. Before the experiments, the bentonite clay powder was sintered at 1250 °C for 4 h to transform montmorillonite (the main constituent of bentonite clay) into mullite 3Al₂O₃·2SiO₂, quartz, and cristobalite to represent the alumino-silicate matrix. After heat treatment, the fireclay pieces were ground in a tungsten carbide mortar to be mixed with slag from plant B. Powders were dried in an oven before

<figure><img src="image_4.png" /><figcaption>**Fig. 1.** Condition of the refractory lining in the intermediate and lower zones of a waste-to-energy facility. Intermediate area: irreversible volume expansion of tiles volume expansion and cracks. Combustion zone: slag deposit on lining and loss of matter on the hot face of tiles.</figcaption></figure>
<figure><img src="image_5.png" /><figcaption>Fig. 2. Oxide-bonded SiC refractory tile and its microstructure.</figcaption></figure>

<table><caption>**Table 1**

Chemical composition and main properties of the material (data from supplier).</caption><thead><tr><th></th><th></th><th></th><th></th></tr></thead><tbody><tr><td rowspan="4">Chemical composition (wt%)</td><td>SiC</td><td>90 ± 2</td><td>Cold crushing strength (MPa)</td><td>110 ± 20</td></tr><tr><td>SiO<sub>2</sub></td><td>7.3</td><td>Refractoriness under load at 0.2 MPa (°C)</td><td>> 1700</td></tr><tr><td>Al<sub>2</sub>O<sub>3</sub></td><td>2.5</td><td>Thermal expansion coefficient (K<sup>-1</sup>)</td><td>4.8 × 10<sup>-6</sup></td></tr><tr><td>Fe<sub>2</sub>O<sub>3</sub></td><td>0.2</td><td></td><td></td></tr><tr><td>Bulk density (g.cm<sup>-3</sup>)</td><td></td><td>2.65 ± 0,05</td><td>Thermal conductivity (W.m<sup>-1</sup>.K<sup>-1</sup>)</td><td>14</td></tr><tr><td rowspan="2">Open porosity (vol%)</td><td rowspan="2"></td><td rowspan="2">15 ± 2</td><td>Mean specific heat (kJ.kg<sup>-1</sup>.K<sup>-1</sup>)</td><td>13</td></tr><tr><td>20–1100 °C</td><td>1.13</td></tr></tbody></table>

being mixed in a tumbler mixer (Turbula T2C) for 2 h. Then, the powders were pressed and the pellets were fired at 1200 °C for 5 h in alumina crucibles.

refractory tile. SEM observations on slag deposit samples show multi-phase systems with large pore size distribution. Local chemical analyses are consistent with pXRD measurements and confirm the presence

## 3. Results

### 3.1. Slag deposit characterisation

Samples of slag deposit were collected in the combustion zone of three waste-to-energy facilities in France after a year of operation at least. Slag deposits are formed from successive layers and the presence of spherical pores indicates a gaseous activity (Fig. 4). The chemical composition of slag deposits from the different facilities was analysed by X-ray fluorescence after grinding and homogenising.

Despite the different sampling points, chemical compositions are close and are mainly composed of CaO-Al₂O₃-SiO₂ with the presence of alkaline, alkaline earth, and metal oxides (Table 2). Quite similar chemical composition of slag deposits is found by Pilate et al. [6]. Additionally, pXRD patterns show the same main phases in samples from all three facilities: akermanite Ca₂MgSi₂O₇, gehlenite Ca₂Al₂SiO₇, wollastonite CaSiO₃, and anorthite CaAl₂Si₂O₈. Although slag deposit on the refractory lining has to be distinguished from bottom ash (final residues at the release of the grate), these results show some similarities with studies of bottom ash from Vehlow et al. [20] and Eighmy et al. [21] where akermanite, gehlenite, and anorthite are also found.

Fig. 5 shows optical and SEM observations of the slag deposit on a

<figure><img src="image_6.png" /><figcaption>Fig. 3. Pore size distributions and open porosity values on two samples from a slice of a refractory tile. Measurements of pore size distribution by mercury porosimetry and open porosity by Archimedes method.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Fig. 4. Cross section of a slag deposit presenting a succession of porous layers.</figcaption></figure>

of akermanite, gehlenite, and wollastonite which are the main phases. The slag deposit presents a microstructure that evolves along its thickness and it can be decomposed into different layers which are separated by dotted lines in the figure. The innermost part layer (microstructure no. 1 in Fig. 5) shows the presence of SiC grains surrounded by an alkaline-enriched vitreous phase (5 wt% Na₂O and 6 wt% K₂O) and no silica layers around the particles were observed. In layer no. 2, the microstructure features wollastonite CaSiO₃ crystals and diopside-type CaMgSi₂O₆ crystals contained in a fissured alkaline-enriched phase. The microstructures of the layers no. 3 and no. 4 show bigger crystals of wollastonite, gehlenite, and akermanite. Small crystals of perovskite CaTiO₃ and (Al,Ti,Fe)-enriched calcium silicate phase (bright phases) are also visible in the last two layers.

## 3.2. Post-mortem SiC-based refractory tiles

Macroscopic observations of the post-mortem SiC-based refractory tiles show an important loss of matter and the presence of crater-like cavities on the hot face (Fig. 6). This surface feature is found in every collected sample. No irreversible volume expansion or cracks are observed on the corroded refractory tiles from the combustion zone. Fig. 7 represents the interface area between the slag deposit and the tile on a cross-section. As detailed in the previous section, macropores and some fragments of refractories located on the side of the slag deposit are pointed out.

The microstructure of corroded tiles near the hot surface is presented in Fig. 8. It shows lots of pores and an attack of the bonding phase. EDS measurements of the bonding phase and X-ray maps near the hot surface of tiles highlight the presence of alkaline elements in the refractory matrix. The microstructure in this area is similar to the one observed on the slag deposit sample near the interface (Fig. 5 layer no. 1). The observations point out the presence of a silica layer revealing the oxidation of SiC fines. This layer can appear with microcracks as observed in the orange-framed microstructure in Fig. 8.

SiC oxidation as well as S and Cl traces are found beyond 10 mm in depth into the material. Under NaCl, KCl, Na₂SO₄, and K₂SO₄ influence, the transformation of amorphous silica into cristobalite is possible in the temperature range in the tiles [16,22]. Volume change due to the crystallisation of amorphous silica into cristobalite induces cracking of the layer [23] promoting oxygen transport and, reducing the passivation layer of the SiC grains. Fig. 9 presents powder X-ray diffraction patterns of a post-mortem refractory tile cut in three layers. Results show the presence of cristobalite SiO₂ and sanidine K₁ₓNaₓAlSi₃O₈ which is the high-temperature form of potassium feldspar [24] in the whole tile. The comparison of XRD measurements point also out a higher intensity of cristobalite peaks nearby the cold face of the refractory tiles which can

<table><caption>**Table 2**

Chemical composition of slag deposit collected in three waste-to-energy facilities. XRF measurements.</caption><thead><tr><th>wt%</th><th>SiO<sub>2</sub></th><th>CaO</th><th>Al<sub>2</sub>O<sub>3</sub></th><th>Na<sub>2</sub>O</th><th>MgO</th><th>P<sub>2</sub>O<sub>5</sub></th><th>Fe<sub>2</sub>O<sub>3</sub></th><th>TiO<sub>2</sub></th><th>K<sub>2</sub>O</th></tr></thead><tbody><tr><td>Plant A</td><td>42.5</td><td>29.7</td><td>16.7</td><td>2.3</td><td>2.2</td><td>2.1</td><td>2.5</td><td>1.4</td><td>0.6</td></tr><tr><td>Plant B</td><td>39.1</td><td>33.9</td><td>13.0</td><td>2.9</td><td>2.9</td><td>2.9</td><td>2.7</td><td>2.3</td><td>0.3</td></tr><tr><td>Plant C</td><td>35.1</td><td>36.9</td><td>14.8</td><td>3.2</td><td>2.5</td><td>2.9</td><td>2.5</td><td>1.6</td><td>0.5</td></tr></tbody></table>

<figure><img src="image_8.png" /><figcaption>Fig. 5. Polished cross-section of a slag deposit sample in contact with a refractory tile. Observations with an optical microscope in dark-field mode (centre). SEM microstructures (backscattered electron mode) and EDS analyses of the different layers (no. 1 to no. 4) of the sample. Chemical composition is given in wt%.</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Fig. 6.** Corroded refractory tile from the combustion zone in a waste-to-energy facility.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>**Fig. 7.** Interface between slag deposit and refractory material. Presence of corroded refractory fragments in the slag deposit.</figcaption></figure>

be explained by silica crystallisation in presence of alkali chlorides and sulphates. Mullite peaks from the sintering of the bonding phase are only visible on the slice of the cold zone of the used tile. Although corrosion products were found deep in the tiles, the dissolution mechanism of the refractory matrix is rather superficial where the main microstructure changes occur up to 3 mm below the hot face of the tiles.

### 3.3. Thermodynamic calculations

Observations and analyses are completed with thermodynamic calculations using FactSage™ v8.2 software (FactPS and FToxid databases)

[25] to better understand the thermochemical behaviour of the slag deposit and the slag/refractory interactions. The viscosity of the liquid phases is calculated with FactSage™.

For the slag deposits, calculations are done from chemical compositions detailed in Table 2. From the results, it is possible to get the liquid content at 1200 °C which is the service temperature in the combustion zone of plants, the solidus and liquidus temperatures, the primary phase of crystallisation, and the viscosity of the molten phase at 1200 °C. These results are summarised in Table 3.

At 1200 °C, slag deposits from the three plants are mainly liquid (73 wt% for the lowest value). The solidus temperatures of samples from plants A and B are similar. As well, the liquidus temperatures between the two samples are close with a difference of 15 K. The predicted primary crystalline phase of every sample is a calcium silicate such as anorthite CaAl₂Si₂O₈, akermanite Ca₂MgSi₂O₇, or a phase of the melilite group which is consistent with XRD results and microscopic observations. The viscosity of the liquid phase is estimated between 3 and 12 Pa. s. As expected, the more basic (CaO/SiO₂), the less viscous.

The interface between the slag deposit and the refractory at 1200 °C was simulated with thermodynamic calculations. The results are presented in Fig. 10. The figure shows the phases according to the ratio between the refractory (R) and the slag deposit (S). R/(R+S) = 0 gives the slag mineralogic composition while R/(R+S) = 1 gives the refractory composition. The composition of the refractory material (R) given in Table 1 and the chemical composition of the slag deposit from plant B (S) given in Table 2 were retained for the calculations. Results show a large amount of liquid phase and significant quantities of wollastonite CaSiO₃, melilite Ca₂MgSi₂O₇/Ca₂Al₂SiO₇, feldspars (CaAl₂Si₂O₈ and (Na,K)AlSi₃O₈) and, silicon carbide SiC. These solid phases are consistent with post-mortem analyses of the refractory material and the slag deposit. Calculations also exhibit several minor solid phases such as iron, iron phosphides, perovskite CaTiO₃, diopside CaMgSi₂O₆, cordierite Mg₂Al₄Si₅O₁₈, Ti and Fe carbides, silica and, mullite.

From R/(R+S) = 0–0.044, silicon carbide is not present. However, wollastonite CaSiO₃ increases widely in this range. The calculations showed a release of CO gas from SiC oxidation along the interface and the maximum gas volume is estimated at 6.2 litres for R = 0.023. Silicon carbide SiC reacts in contact with slag deposit to produce SiO₂ which is dissolved in the liquid phase. The increase of silica tends to form wollastonite with CaO contained in the liquid slag deposit. The oxygen supply for the SiC oxidation would come from the reduction of oxides such as iron and phosphorus oxides which then combine. These calculations do not consider the atmosphere of the combustion chamber. In real conditions, silicon carbide is exposed to other gaseous oxidising compounds (O₂, H₂O...).

From R/(R+S) = 0–0.95, the liquid phase is noticed. Mullite (e.g. phase formed from the sintering of the alumino-silicate matrix) is noted from 0.96. Before 0.96, alumina, which is necessary for the formation of

<figure><img src="image_11.png" /><figcaption>**Fig. 8.** Microstructure of the hot face of a corroded refractory tile. Fissured silica layer around a SiC fine and EDS analysis of corroded bonding phase (orange-framed microstructure). Chemical composition is given in wt%.</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>**Fig. 9.** X-ray diffraction patterns of a refractory tile before and after use. Distances from the hot face are indicated on the right. Indexed phases: silicon carbide SiC (X), cristobalite SiO₂ (C), mullite Al₆Si₂O₁₃ (M), quartz SiO₂ (Q), sanidine K₀.₈₃Na₀.₁₇AlSi₃O₈ (S) and tungsten carbide grinding contamination WC (W) from grinding contamination.</figcaption></figure>

<table><caption>**Table 3**
High-temperature behaviour of slag deposit simulated by thermodynamic calculations (FactSage™ v8.2).</caption><thead><tr><th>Slag deposit</th><th>Plant A</th><th>Plant B</th><th>Plant C</th></tr></thead><tbody><tr><td>T<sub>solidus</sub> (°C)</td><td>968</td><td>971</td><td>1060</td></tr><tr><td>Liquid content at 1200 °C (wt%)</td><td>96</td><td>88</td><td>73</td></tr><tr><td>T<sub>liquidus</sub> (°C)</td><td>1220</td><td>1235</td><td>1293</td></tr><tr><td>Primary phase of crystallisation</td><td>Anorthite CaAl<sub>2</sub>Si<sub>2</sub>O<sub>8</sub></td><td>Akermanite Ca<sub>2</sub>MgSi<sub>2</sub>O<sub>7</sub></td><td>Melilite Ca<sub>2</sub>MgSi<sub>2</sub>O<sub>7</sub>/Ca<sub>2</sub>Al<sub>2</sub>SiO<sub>7</sub></td></tr><tr><td>Basicity CaO/SiO<sub>2</sub></td><td>0.7</td><td>0.9</td><td>1.1</td></tr><tr><td>Viscosity of liquid phase at 1200 °C (Pa.s)</td><td>11</td><td>5</td><td>4</td></tr></tbody></table>

mullite, is contained in the liquid phase and, feldspars (CaAl₂Si₂O₈ mainly). Thus, the refractory bonding phase is dissolved in contact with the slag deposit. Feldspars, diopside CaMgSi₂O₆ and, cordierite Mg₂Al₄Si₅O₁₈ are formed as reaction products. The calculations also highlighted the increase of the liquid phase viscosity to reach a value of 615 Pa.s due to the rise of network formers such as SiO₂. Despite these results do not take the corrosive flue gas and the dynamic environment into account, they are consistent with experimental observations with the presence of feldspar, melilite, and wollastonite and the oxidation of silicon carbide.

## 4. Wear mechanism of oxide-bonded SiC tiles in the low area of waste-to-energy facilities

The post-mortem examination of oxide-bonded SiC refractory tiles revealed interactions between solids, liquids, and gases. The origin of gases may be attributed to the combustion, the gas condensation at the slag deposit/tile interface, the residual gases trapped in the viscous liquid phase of the slag deposit, and the oxidation of SiC particles. The contribution of every cause could not be determined in this study. Nevertheless, a large contribution from SiC oxidation was assumed due to the significant amount of SiC in the refractory tiles. Observations and thermodynamic calculations exhibited the dissolution of the refractory matrix. The formation of liquid and gaseous phases coupled with mechanical and thermal stresses leading to the weakening of the interface and the spalling of the refractory material are illustrated in Fig. 11 and discussed in the following sections.

### 4.1. Dissolution of the refractory matrix and alkaline contamination

Analyses of the corroded refractories pointed out significant concentrations of sodium and potassium in the bonding phase. The presence of sanidine K₁ₓNaₓAlSi₃O₈ (high-temperature form of potassium feldspar) was detected with XRD measurements and local traces of sulphur and chlorine in the refractory matrix were measured with EDS. As demonstrated by Prigent et al. [18], the formation of microcline KAlSi₃O₈ (low-temperature form of potassium feldspar [24]) in the pores of the refractory tiles can arise due to the sulphation of alkaline chlorides (Eq. 3) [26]. In the combustion area, the flue gas temperature is higher than that in the intermediate area [5] which can explain the presence of the high-temperature potassium feldspar polymorph. High concentrations of alkaline elements were also observed close to the deposit/material interface. In this area, no new phase precipitation was observed and a direct dissolution of the refractory matrix by liquid oxides occurred. The impregnation of molten slag deposit leads to the enrichment of alkali compounds in the refractory matrix also assisted by the condensation of salts.

$$ (Na,K)Cl_{(g)} + 1/2 SO_{2} (g) + 1/4 O_{2} (g) + 1/2 H_{2}O_{(g)} = 1/2 (Na,K)_{2}SO_{4} (s) + HCl_{(g)} \quad (3) $$

The behaviour of the fired clay bonding phase of SiC-based refractory in contact with slag was studied at the laboratory scale. After 5 h at 1200 °C, the mix (50 wt% slag – 50 wt% refractory matrix) was fully melted and gas bubbles were visible due to residual carbonates and sulphates contained in the slag (Fig. 12). Microscopic observations revealed the presence of anorthite crystals surrounded by a silica-rich phase liquid at 1200 °C. The dissolution of the alumino-silicate matrix, therefore,
<figure><img src="image_13.png" /><figcaption>**Fig. 10.** Representation of the interface between the slag deposit and the refractory material at 1200 °C obtained by thermodynamic calculations with FactSage™ v8.2 software and the evolution of the liquid phase viscosity. “R” and “S” stand for refractory material and slag deposit respectively.</figcaption></figure>

contributes to the increase of the liquid phase viscosity estimated at 90 Pa.s with FactSage™ software.

## 4.2. Formation of pores by oxidation of SiC grains

The interface between the refractory tiles and the slag deposit is characterised by the presence of macropores. The formation of these macropores may be explained by the oxidation of SiC grains. To confirm it, the behaviour of SiC grains in contact with slag was studied at the lab scale according to the experimental procedure presented in Section 2.2. After 5 h at 1200 °C, no silica layer is observed around SiC aggregates (1–3 mm). Small spherical particles of ferrophosphorus coming from the reduction of iron and phosphorus oxides by silicon carbide are present. Numerous pores around SiC grains are observed and a higher SiO₂ content (50 wt% initially) is noticed in the slag containing SiC grains (Fig. 13a) which shows the oxidation of SiC. These observations are consistent with post-mortem examination and thermodynamic calculations. In these conditions, silica from SiC oxidation is dissolved in the slag which influences the viscosity of the liquid phase and affects gas

<figure><img src="image_14.png" /><figcaption>**Fig. 11.** Cycling wear mechanism of refractory tiles in the low area of waste-to-energy facilities. Projections of slag deposit on the refractory lining (step 1). Dissolution of the bonding refractory matrix and SiC oxidation by liquid slag (step 2). Interface separation and spalling of refractory material (step 3). Damaged refractory lining exposed to new projections of slag deposit (step 4).</figcaption></figure>

release from silicon carbide oxidation.

Silicon carbide oxidation is influenced by the temperature, the oxygen partial pressure, the water vapor [27], as well as the grain size [28]. The latter influences significantly the oxidation rate and its kinetics. This was confirmed with thermogravimetric measurements on the two different grain size distributions (1–3 mm and 0–250 µm) that are generally used for the formulation of the refractory material. The measured mass variations corresponding to the oxidation of SiC particles after 3 h at 1200 °C under air atmosphere were 0.07 % for the coarse distribution and 3.4 % for the finer distribution. As the grain size distribution of silicon carbide in the refractory tiles is wide, fine particles represent an important parameter in the oxidation process and formation of silica. In contact with the slag, silica released from the oxidation of SiC fines reacts with CaO contained in the slag which leads to the formation of wollastonite crystals (Fig. 13b). These observations are similar to those observed in the second layer of the slag deposit stuck to the refractory lining (Fig. 5 microstructure no. 2) and are consistent with thermodynamic calculations presented in Section 3.3.

## 4.3. Spalling of the refractory material and cyclic wear

At the service temperature, the bonding phase of the refractory material is dissolved in contact with the liquid slag deposit, and SiC particles from the refractory material are oxidised. The viscosity of the liquid phase is impacted due to high silica content, resulting in gas sequestration at the deposit/tile interface and the small depth of impregnation by the molten slag in the tiles. Beyond the macropores, the slag deposit continues to react with the fragments of corroded refractories which leads to the formation of a thin layer containing wollastonite CaSiO₃ and diopside-type CaMgSi₂O₆ crystals that might tend to transform into wollastonite. Fig. 14 illustrates the dissolution of the fired clay bonding phase of the refractory, the oxidation of SiC grains, the formation of macropores, and the formation of silicate crystals.

The separation of the interface and the spalling of the refractory material occur afterward and can be attributed to several phenomena:

- the superficial embrittlement of the refractory tile around the interface related to the gaseous and liquid phases;

- the presence of cracks due to the formation of expansive phases as wollastonite in a thin layer near the interface;
<figure><img src="image_15.png" /><figcaption>**Fig. 12.** Pellets of slag and slag mixed with fired clay (refractory matrix) after 5 hours at 1200 °C (left). Microstructure of the slag/fired clay mixture containing anorthite crystals in a silica-rich phase (right). Chemical composition is given in wt%.</figcaption></figure>

<figure><img src="image_16.png" /><figcaption>**Fig. 13.** Microstructures of the reactivity between slag and 1–3 mm SiC grains (a.) and 0–250 µm (b.) after 5 hours at 1200 °C under air atmosphere. Chemical compositions are given in wt%.</figcaption></figure>

<figure><img src="image_17.png" /><figcaption>**Fig. 14.** Illustration of dissolution and oxidation steps. Dissolution of the bonding phase (left). Oxidation of SiC grains and formation of macropores due to the gas release (centre). Formation of silicate crystals due to the reaction between refractory fragments and the slag deposit (right).</figcaption></figure>

- the accumulation of the slag deposit layers;

- thermal variations during waste combustion inducing a mismatch of coefficients of thermal expansions.

As a result, a loss of material in the hot face of refractory tiles occurs.

After the first wear cycle, the dissolution of the bonding phase is promoted by the presence of contaminant elements at the interface and alkali feldspars in the depth of the tiles from gas condensation. The damage process continues until a complete wear or the replacement of the tile.
## 5. Conclusion

The damage mechanism of the oxide-bonded SiC refractory tiles in the combustion zone of waste-to-energy facilities was characterised. They are governed by the following cyclic steps which lead to a loss of material:

- dissolution of the bonding phase and oxidation of SiC grains by the molten slag deposit;

- formation of macropores at the interface between slag deposit and refractory lining;

- spalling of the refractory material.

This study highlights the importance of silica in the wear mechanism because of its impact on viscosity and its high reactivity with others components. Although the viscosity is not experimentally measured, observations and calculations showed its influence on the liquid phase in the damage process owing to the gas sequestration from SiC oxidation.

Some refractory material improvements can be suggested from these results:

- the modification of the composition of the refractory matrix to get better corrosion resistance;

- the use of a SiC grain size distribution and sintering temperature adapted to limit silicon carbide oxidation;

- the reduction and the homogenisation of open porosity and pore size to limit gas and liquid penetration and to avoid the acceleration of the wear mechanism rate.

### CRediT authorship contribution statement

**Jean-Michel Brossard:** Writing – original draft, Resources, Investigation. **Patrick Ganster:** Writing – original draft, Supervision. **Emma-nuel de Bilbao:** Writing – original draft, Validation, Supervision, Software, Methodology, Funding acquisition, Conceptualization. **François Valdivieso:** Writing – original draft, Supervision. **Jacques Poirier:** Writing – original draft, Supervision, Methodology, Funding acquisition, Formal analysis, Conceptualization. **Audrey Tixier:** Writing – original draft, Visualization, Resources, Methodology, Investigation, Formal analysis, Conceptualization. **Cyrille Varona:** Writing – original draft, Supervision, Resources, Funding acquisition.

### Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgements

This work was supported by the French National Association for Research and Technology (ANRT) and BONY SAS (no. 2019/0226). The authors thank also operators from waste-to-energy facilities for their help in collecting samples and their expertise.

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