printing 3D models of coral colonies; they generated accurate 3D models suitable for research and education. Ngo et al., in 2018 discussed the revolutionary applications of AM (fused deposition modelling, powder bed fusion, inkjet printing, and contour crafting, stereolithography, direct energy deposition, laminated object manufacturing) in trending applications, i.e. biomedical, aerospace, buildings and protective structures. Materials of survey were metal alloys, polymer composites, ceramics, and concrete. They identified few drawbacks as defaults (voids), inferior and anisotropic microstructure and mechanical properties, divergent from design to execution, layer-by-layer appearance. Additional information can be found in the indexed work of Bourell et al., in 2017. Harris et al., in 2019 reported multimaterial 3D printing by fused filament fabrication and showed the ability of Thermoplastic PolyUrethane (TPU) to adhere to Acrylonitrile Butadiene Styrene (ABS) and acrylonitrile styrene acrylate (ASA). More recently, Jovic et al., in 2019 outlined the most widely used biomaterials derived from land plants and marine algae (nanocellulose, pectin, starch, alginate, agarose, fucoidan, and carrageenan) with an in-depth focus on nanocellulose and alginate. They discussed their potential in 3D bioprinting for tissue engineering, drug delivery, wound healing, and implantable medical devices. And Wasti and Adhikari in 2020 reviewed biomaterials (Poly-Lactid Acid (PLA), PolyHydroxy-Alkanoate (PHA) and blends used in fused deposition modeling technique.

Various ranges of applications are concerned, for instance aerospace, electronics, aeronautics, automobile, art, biomedical, healthcare (Azad et al., 2020; Fan et al., 2020), architecture, and building (Shakor et al., 2019), fashionable food industries (Godoi et al., 2016). Since the pioneer works reported in 2013, 4D printing is introduced and proposes a new strategy that involves the design of a 3D printed part able to undergo a controllable shape change with time as described by (Ge et al., 2013; Tibbits, 2014; González-Henríquez et al., 2019).

The Rapid Prototyping (RP) is based on the local deposition of small volumes of material to directly form as a solderlike shape. The available technologies have been improving constantly since the first French artist Willeme's photosculpture in 1860. The history and evolution of AM technologies are comprehensively discussed by Bourell et al. (Bourell et al., 2009). Among the AM technology, the concept of continuous microextrusion printing is used. The material is selectively dispensed through a nozzle over a movable platform (build stage). When using extrusion, two primary approaches are found. i. First, the commonly FDM (Fused Deposition Modeling) where the material is molten (heat) inside the channel of the liquefier. This step allows flow out through the nozzle and bonding before complete solidification. And ii. second, the revisited use of a supersaturated fluid slurry, as a material for extrusion, is usually carried out without heating. The later includes pressure-assisted methods such as pneumatic (air, compressed gas) and piston methods (Godoi et al., 2016), as both presented in micro-extrusion bioprinters (Bose et al., 2018; Kumar et al., 2019).

Continuous efforts have made possible the use of geo-materials in 3D printing technologies. Issues, potentialities and opportunities are illustrated in works such as (Travitzky et al., 2014; Zocca et al., 2015; Gosselin et al., 2016; Zhong et al., 2017;

Buswell et al., 2018; Lin et al., 2019; Shahsavari and Hwang, 2020; Wang et al., 2020). However, progress is still needed in this geo-materials domain. Indeed, as introduced by (Nair et al., 2020; Shakor et al., 2020a; Shakor et al., 2020b), the printing procedures and dimensional accuracies of printed objects lack details. These details call for a large crossing of knowledge, i.e. interdependent factors of geo-based matrix/control of the process/functional properties.

For this reason, the present research has evaluated kaolin clays in its pristine state. Kaolinite is one of the widely used aluminosilicate minerals, encountered in the manufacturing of traditional ceramics, paper, pottery, paintings and cosmetics. Generally speaking, chemical additions (or improvers) are brought together to propose the best formulation to meet a specific clay-based material; the improvement requiring to answer diverse features as thermal and mechanical properties, rheology, inter-particle interactions, stable phase (Revelo and Colorado, 2018; Revelo and Colorado, 2019).

The regular fused deposition modeling apparatus (FDM) was adapting with a pressure-assisted piston connected to a single screw extruder before the regular nozzle used for printing thermoplastics. For a fundamental approach, the clay-water mixture was selected as a model system, without any chemical additions (or improvers). Concomitant control of the tested process parameters, the kaolin hydration and the complexity of printed structures permitted to understand the effect of the printing conditions on the green mixture. A like-pottery part was fabricated, illustrating the ability, and feasibility of this technology. Post-processed temperature curing, i.e. sintering or as called "high temperature oven-dry," validated the structural stability of the pottery.

# MATERIALS AND METHODS

## Material Properties

### Kaolin Clay and Slurries

Kaolin clay was chosen due to its availability worldwide as minerals found in the Earth's crust, particularly in South of France with its pottery. The porcelain powder (PT010B: composition of kaolin clay 50-100% and quartz 10-25%, density of 2.6 g/cm³) was supplied by Imerys (France).

The eco-responsible character and the low cost of kaolin clays is also an important factor for scaling up and promoting the results, for individual artists to big traditional ceramics industry and specific use. The single clay was used throughout all slurries, without purification. Only the amount of water was varied from one formulation to the other to, in the end, work with a single formulation optimized for processing. Therefore, the chemical composition and the solid particle characteristics (network microstructure, size distribution and shape of particles) may be deemed as constant and practically without any effect on the final printed parts.

## From Atomic Composition to Hexagonal Sheets

The crystallographic structure of kaolin is made of two layers. One is formed by octahedral meshes sharing their oxygen atoms
<figure><img src="image_436.png" /><figcaption>**FIGURE 1** | Hexagonal booklet-like sheets of native (powder) kaolin clay (A) scale of 2 µm, (B) scale of 500 nm. Hexagonal structure forms of steps which reveal successive layers of sheets.</figcaption></figure>

at their corners with other neighbour octahedral meshes. The second one is an assembly of tetrahedral meshes connected to the octahedral layer with hydroxide groups. Aluminum is in the center of the octahedra and silicon in the center of the tetrahedra. Kaolinite is a widely used hydrated alumino-silicate minerals, mainly composed by kaolin clay with a theoretical molecular formula of Al₂O₃·2SiO₂·2H₂O (Bergaya et al., 2006; Loginov et al., 2008). In the observable morphology scale, the assembly of a layer formed of octahedra to a layer formed of tetrahedra forms is called a leaflet. The surfaces of these single sheets have hydroxide and oxygen atoms which allow water molecules to be inserted between the sheets (Awad et al., 2017; Awad et al., 2019). **Figure 1** displays the hexagonal sheets, on native kaolin clay.

## Hydration

The suspensions of kaolin clay were prepared from dry kaolinite powder. These suspensions were conducted by mechanically mixing the clay powder with water (or dispersant solution) at appropriate quantities. Water to clay ratios (in % wt) was of 23 and 27% (see Results part). The water content was checked for each test, by collecting a certain mass of the suspensions and measuring it before and after drying at 37°C. This temperature of 37°C was selected to follow the kaolin clay loss of water mass over time and to avoid any deformation of objects. For a standard consideration, recent work by Shakor et al. (Shakor et al., 2020a) post-processed the objects in an oven for 3 h at 60°C.

## Viscosity and Filament Size

The rheological behavior of the green slurries was directly verified with the 3D printer. The flow rate was determined using a capillarity test (controlled directly by the 3D printer). Appropriate G-code files with different “extrusion screw speed rates” were tested to perform a single linear path with the printer. The nozzle speed was varied while the nozzle diameter stayed constant (1.0 mm). Then a visual appreciation was made to maintain the optimal print flow. This is necessary to calibrate the throughput, which is unique to each individual printer.

The real flow and the flow programmed from a test code in which the printing time was set (and so only the flow was changed) were compared (see Results part, **Figure 2**). Then single line objects were printed to determine the optimal speed

<figure><img src="image_437.png" /><figcaption>**TABLE 1** | Key parameters in 3D printer, constant flow of 1 mm/s, nozzle diameter of 1.0 mm.</figcaption></figure>

<table><thead><tr><td>Material flow</td><td>Printing speed</td><td>Layer thickness</td><td>Wall thickness</td><td>Infill</td></tr></thead><tbody><tr><td>60%<sup>a</sup></td><td>1 mm/s<sup>b</sup></td><td>1 mm</td><td>1 mm</td><td>Concentric</td></tr></tbody></table>

of the nozzle. Again a visual inspection of the surface of the filament was made to verify the homogeneity of the printed lines. A slow speed of 1 mm/s was chosen to produce a very good surface definition of the pottery, with a thickness of 1 mm for all layers (see **Table 1**).

## Methodology and Approach

As identified in the spectrum of works of (Buswell et al., 2018), the clay in all its forms is surrounded: properties of wet materials prior to solidification (i.e. pumpability and extrudability); hardened properties of 3D parts, achievement of geometric conformity.

<footnote><sup>a</sup>Proper to our software controller.</footnote>

<footnote><sup>b</sup>To be tested directly on the printer with the clay extruder.</footnote>
## Pressure-Assisted Deposition Modelling

A commercial Delta WASP 2040 TURBO Printer (Italy) was employed to produce prints of the kaolinite parts. The Delta printer is composed of three mechanical arms powered by three independent motors. In the present work, the 3D printer is a modification of a regular Delta fused deposition modeling (FDM) machine in which the thermo-mechanical system for carrying out the extrusion process was adapted to the clay configuration. In clay configuration, no more temperature control is possible.

The machine extrusion unit is connected to a piston tank (3 L), which contains the clay dough, and a nozzle for filament layer deposition. The material is directly transported in its paste (i.e. slurries) form with a pressure-assisted piston system (2.5–3.5 bar) just before it is extruded (see Results part). The whole system works at ambient temperature (nozzle, print with tray).

The Delta printer can produce parts of 400 mm in height and 200 mm in diameter. The advantage of such a printer is its aptitude to print large parts in height unlike other printers. The minimum thickness of a layer is about 50 µm and the maximum thickness is 330 µm. The maximum speed at which the paste can be deposited is 300 mm/s just as the speed of moving the print head.

The modelled-parts (digital one) were scaled and orientated in an upright position. They were printed with the amount of infill for all prints set at 100%, without any scaffold support material, i.e. only walls.

Printed parts were developed layer by layer at a thickness of h = 1 mm per layer. The thickness was optimized by reporting a study with thickness variation from 0.3, 0.6 and 1.0 mm respectively. Simultaneously, the speed of the moving of the print head was taken to be 1 mm/s. The speed of 1 mm/s is in adequacy with the flow speed that was used in the Gcode "capillary" test, directly on the 3D machine.

The basement orientation layers, that is, the first layers of the support surface or foundation of the modeled part, have been defined by the printer specifications (concentric or square). No effect was evidenced in the quality of the successive building layers, even if a concentric or squared foundation was chosen.

But it could and can find its importance when looking at drying, sintering and mechanical behavior.

Skin overlap was set to 1, which indicates full overlap between the outer perimeter and the infill.

To ensure good adhesions of the first deposit layer (which supports the rest of the structure), different support surfaces for impression were tested, i.e. aluminum foil, paper, and polishing paper (see Results part).

## Specimen Replication, Software vs. Hardware

Each trial PAM files (gmesh, version 4.1.4) were imported into Cura 3D, an open source slicing software for 3D printers (version 4.4.1) to generate the individual printing slices, paths, or patterns (software). These slices were then compiled and exported as a G-code file ready to be printed on the printer. The file transfer is made via an SD card to the printer (hardware).

## Thermal Stabilization, Sintering

Selected printed parts were open-air dried at 37°C for 24 h. They were then sintered in a furnace (Carbolite type, United Kingdom) at 500°C (campaign 1) or 700°C (campaign 2) for 1 h, and then at 1,200°C for 1 h, respectively in air atmosphere. The conditions were selected by following a pottery process and to control the evolution of the metastable states encountered in clay. The heating and cooling temperature ramps were 10°C/min.

## Structural Characterization and Mechanical Tests

### Nano-Micro-Meso-Scales Structure

The morphology of nano-micro scales structure was evaluated for the raw and in-transition materials. Scanning Electron Microscopy with a high resolution Field Emission in high vacuum mode (MAIA3 Triglav FESEM, TESCAN type, Czech Republic) was used. Samples were mounted on carbon tape and coated with a 7 nm platinum film controlled by a quart crystal nano-balance using a Quorum Q150TES to 30 mA AC. The meso-scale structure was characterized with a binocular microscope (Olympus SZH10, Japan) and an optical microscope (Olympus PMG3, Japan) in reflexion mode.

### Compression

Bulk compression test on sintered parts were performed at 1 mm/min, total deformation of 25%, at 23°C (Dartec HA250 type, United States).

### Hardness

Bulk hardness was tested using a single lap joint specimen (on a sanded surface of sintered parts), by Knoop indentation tests done for 10 s in force loading 1 Kgf, at 23°C (Buehler type, Switzerland) (Ghorbal et al., 2017).

## RESULTS AND DISCUSSIONS

To ensure reproducibility, a methodology has been developed. Itemized protocols concern first the slurries and second the adaptation of the thermoplastic 3D printer to clay configuration.

### Slurries

#### Hydration

In this step, our g-code calibration test maintains two process parameters constant: the flow of 1 mm/s and the nozzle diameter of 1.0 mm. Different hydration %, in H₂O wt, have been tested. The conditions to allow the identification of as named "good printing" slurries are summarized in **Table 2**. The filament diameter and the collapse between filaments were checked post process after the g-code calibration test. The viscosity can play an important role in ensuring defect-free products (Chan et al., 2020). The best solution of the g-code test depends on the conjunction with a suitable combination of printing pressures and speeds. Indeed, at first kaolinite dough should present viscoelastic characteristic to allow extrusion through the fine
<table><caption>**TABLE 2** | Percentage of water in printing paste, constant flow of 1 mm/s, nozzle diameter of 1.0 mm.</caption><thead><tr><th>Water, %wt</th><th>Aspect with illustration</th><th>Consequence</th></tr></thead><tbody><tr><td>&lt;23</td><td>Too solid</td><td>-Poor merger between filaments<br>-Irregular flow</td></tr><tr><td>24.5–26</td><td>Fine line between too solid-too fluid</td><td>-Good performance of the piece<br>-Homogeneous filaments<br>-Regular flow</td></tr><tr><td>&gt;27</td><td>Too fluid</td><td>-Bad behavior of the printed part (slumping)<br>-Non-homogeneous filaments<br>-No regular flow, in the form of drops</td></tr></tbody></table>

<figure><img src="image_438.png" /><figcaption>**FIGURE 3** | Shear-thinning kaolinite dough suspension behaviors, for highest to lowest speed (%).</figcaption></figure>

nozzle, and then turn into self-supporting clay prior post-deposition of the consecutive layers.

## Nonplastic Dry Kaolin Clay

As water is added, the kaolin clay becomes plastic (Table 2A). The water (<27 %wt H₂O) acts as a lubricant, surrounding the clay platelets and allowing them to glide over each other. Surface tension forces acting both on the water and the clay bind them weakly together and allow a printed object to retain its shape (24.5–26 %wt H₂O) (Table 2B). The diameter of filament of kaolinite dough (950 µm checked) respected the nozzle one (1.0 mm). When excess water is added (>27 %wt H₂O), more than the needed minimum to maintain surface tension, water and mass will start to flow (Table 2C). Kaolin clay becomes too fluid to retain the original shape of the object as it was formed from the CAD part.

The development of plasticity in clay is critical to the manufacture of ceramics. Depending on the composition of the clay, plasticity will develop over a narrow range of water content. The hydration window for processing is identified between 24.5 and 26 %wt H₂O; that shows all the importance in preparing valuable slurries.

## Real Flow Rate

In our g-code calibration the flow is saying constant. The flow rate is set on the Cura software as a percentage. It is per default equal to 100%. Data translation between software to hardware jobs needed to make a correction of 60% in hardware flow rate. The average extruder mass as a function of the programmed flow rate was then checked, as collected in Figure 2. The real flow and the programmed flow are perfectly equivalent. The correlation coefficient of 0.999 allows us to validate the adaptation of the throughput in hardware when it is modified in software.

In addition, it can be visually noticed (Figure 3) that for the same movement speed, the flow rate changes the behavior of the kaolinite dough: the higher its speed, the less viscous the fluid. In fact, the supersaturated suspensions of kaolin clay display a shear-thinning (non-Newtonian thixotropic) behaviour (Coussot, 1995; Loginov et al., 2008; Ordonez et al., 2019). The difference in viscosity behavior gives evidence that suspensions of kaolin clay can flow much easier during a print. Thus, the flow of the slurries does not depend merely on its water content but also on the printing speed, and temperature around. This behavior would allow kaolin clay to settle more easily into the n-1 layer, meaning that the cold adhesion would be inherently stronger. The addition of water to raw kaolinite produces a plastic mixture which can be
<table><caption>**TABLE 3** | Tank inlet pressure adjustment, constant flow of 1 mm/s, nozzle diameter of 1.0 mm.</caption><thead><tr><td>Pressure, bar</td><td>Water content, %wt</td><td>Consequence</td></tr></thead><tbody><tr><td>&lt;2.5</td><td>24.5–26</td><td>-No dough flow even when an impression is launched (clogging)</td></tr><tr><td>2.5–3.5</td><td>24.5–26</td><td>-Good flow, dough only came out when the print is launched, setting validated</td></tr><tr><td>&gt;3.5</td><td>24.5–26</td><td>-Dough flowed alone, without the printing being launched</td></tr></tbody></table>

<figure><img src="image_439.png" /><figcaption>**FIGURE 4** | First layer of a cylinder printing of kaolinite dough with different substrates (A) aluminium, (B) paper, and (C) polishing paper (mesh size P320).</figcaption></figure>

shaped without rupture. Kaolinite exhibits plastic properties for primary reasons: kaolin clay particles themselves are crystalline platelet-like structure; with small size of particles; and the presence of water called “water of formation” between the crystalline clay particles. The water acts as a lubricant and allows the clay particles to slide across each other rather than rupture when a shear force is applied.

## Adaptation of Thermoplastic 3D Printer to Clay Configuration

### The Effect of Print Order on Pressure

Two parameters are used to modify the print throughput: the throughput imposed by the software and the pressure imposed in the pressurized tank. To have a total control of the flow, a window for adjusting the imposed pressure in the tank was kept constant, (<2.5; 2.5–3.5, and >3.5) bar, only the software parameter has been changed. It was possible then to fix the following pressure window as given in Table 3, the data are collected on the optimized water content given in Table 2.

### The Effect of Print Order on First Layer Impression Support

The very first layer was printed and only the optimal condition was tested (24.5–26 wt% H₂O, nozzle diameter of 1.0 mm, nozzle

speed of 1 mm/s, software rate 100% (60% in hardware), thickness layer 1 mm, skin overlap set at 1). **Figure 4** illustrates the first layer (foundation) of a cylinder printing of kaolin clay with different substrates A aluminium, B paper, and C polishing paper (mesh size P320, which corresponds to a grain size of 46.2 µm). Aluminum adhesion is the worst; the material slides over the surface and gives heterogeneous dimensional control. For the paper, the adhesion in the early stages of printing is correct, but after a while the adhesion becomes of poor quality. In fact, it has been observed that the surface of the paper does not stay dry but absorbs water (moisture) from the slurry of kaolinite; the surface curls. The best compromise is found for the polishing paper (P320 mesh size). It is mainly due to the surface roughness and probably to its impermeable surface or poor water absorbing character, giving a better anchorage to the deposited material. A P1200 mesh size (equivalent to 15.3 µm) was also tested and did not give better adhesion. Generally speaking, on one hand, the adhesion must be sufficient to allow the slurries to follow the required printing conditions. On the other hand, the adhesion must leave certain freedom to the layer of slurries to avoid any cracking as the slurries start to dry during printing. Indeed, once printed, water can evaporate freely and the slurries will contract (see *To Access to the Effect of Print Order on Structural 3D Making and Sintering*).

## To Access to the Effect of Print Order on Structural 3D Making and Sintering

Kaolin needs post “high temperature oven-dry” treatment to produce solid ceramic objects. This additional processing should be carefully checked and carried out to prevent further defects in the part such as its explosion or change of dimension during sintering (such as an unwanted deformation) and to ensure controlled phase change of the hydrated kaolin (Njoya et al., 2017). The sintering profile is illustrated in **Figure 5A** with morphology in **Figure 5B**. The chemical changes can be summarized as here after:

**Step 1:** heating ramps 10°C/min

**Step 2:** annealing 500°C (or 700°C), 1h

$Al_2(OH)_4Si_2O_3 \rightarrow Al_2O_3 + 2SiO_2 + 2H_2O$

(hydrated kaolin) (metakaolin) (water)

**Step 3:** heating ramps 10°C/min

350-750°C, $Al_2O_3 + 2SiO_2 \rightarrow Al_2O_3 + SiO_2$

(metakaolin) (alumina) (silica)

950-1050°C, $Al_2O_3 + 2SiO_2 \rightarrow MgAl_2O_4 + SiO_2$

(metakaolin) (spinel) (amorphous silica)

**Step 4:** annealing 1200°C, 1h

$Al_2O_3 + 2SiO_2 \rightarrow 2(3Al_2O_3 + 2SiO_2) + 5SiO_2$

(metakaolin) (mullite) (amorphous silica)

**Step 5:** cooling ramp 10°C/min
<figure><img src="image_440.png" /><figcaption>**FIGURE 5** | From hydrated kaolin to pseudo-ceramic: (A) thermal history, (B) morphology after thermal sintering. Depending on the nature of the system, variety of higher temperature mineral phases may crystallize.</figcaption></figure>

## Chemical Reaction-Scale

Before sintering process of ceramics, the first stage is the “air drying” (Figure 5A). It is generally carried out at temperature around 150°C, but in our case 24 h at 37°C (see Material part, Material Properties). At this stage of “air drying,” the water is in the form of films surrounding the kaolin sheets; the mechanically admixed water is lost. The kaolin mixture shrinks and loses its plasticity. The solid to solid transition gives specific results, here the kaolin sheets draw closer together, until they may eventually come into contact and start to adhere. The next water loss takes place more slowly, coming from the pore spaces between the clay particles. Induced-stresses may later lead to structural damage (cracking or warping) of the printed part. The printed layers experience maximum shrinkage while exhibiting maximum of water loss. Water first evaporates from the surface of the printed part, then by capillary action, water rises from the interior to the surface and in turn evaporates. One can notice that this property makes clay a candidate for environmental remediation applications (Garrisson, 2016; Awad et al., 2019).

When the temperature rises to 200–400°C①, any mechanically bound or interstitial water, still present after the ambient drying, is vaporized and driven off.

In addition to water, as the temperature rises to over 200°C, any carbon or organic matter in the clay will start to oxidize, migrate to the surface and finally escape as CO or CO₂. The oxidation rate and completeness depend on the temperature scan (rate), the time hold at the maximum annealing temperature, the size (wall width) of the parts, the clay and paste nature.

Two plateaus then follow the “air drying” stage. Indeed, when the clay is sintering at temperatures over 600°C②, and in particular at higher temperatures of ~1,200°C④, the clay undergoes major structural and chemical changes due to dehydroxylation (i.e. hydroxyl ions (OH⁻) lost which are chemically bound within the clay lattice) (Johnson and Kessler, 1969). Chemically bound water is lost. These changes dictate the conversion of clay into a useful ceramic. High temperatures dislodge ions from original positions held in the clay structure into more favorable sites, thus converting the clay into chemically stable and structurally stronger material.

Plateau at 450–500°C② permits a complete evaporation of the water contained between the sheets of the kaolin structure. The water loss occurs abruptly during the structural transformation of hydrated kaolin to metakaolin with minor shrinkage but with increased porosity. Heating kaolinite above about 500°C results in a disordered crystalline form which is still reversible by rehydration back to its ordered kaolin structure. At 550°C, the organic matter is fully oxidized. At 750–850°C③, mineral water was successfully expelled from the clay. Clay starts to lose its characteristic crystalline structure, the crystal lattice itself will break down. At 950–1,100°C, the structure of kaolin is irreversibly lost, but some illite may persist. With the breakdown in clays, high-temperature alteration and reaction products can appear, generally accompanied by pronounced shrinkage. Si-Al kaolin chain collapses on sintering into metakaolin, followed by the rejection of amorphous SiO₂ during spinel formation. Above 1,000°C④, high temperature new mineral mullite phase (3Al₂O₃ 2SiO₂) is expected to form (Yamuna et al., 2002); >1,100°C is the glassy phase of the ceramic where porosity diminishes rapidly. Notice that starting about 600°C, vitrification or “sintering” can start in the more fusible impurities. At the plateau of 1,200°C, the microstructural arrangements take place during the processing transformation of kaolin to ceramics. In theory, the piece should be cooked over 1,400°C to be completely transformed, but we only had at our disposal an oven reaching 1,200°C (Revelo and Colorado, 2018).

## Morphology-Scale

Morphology of kaolin particles after sintering at 1,200°C for 1 h is illustrated in **Figure 5B**. Compared to the morphology of native kaolin (Figure 1), the hexagonal shapes have given way to more rounded shapes with a densification (Figure 5B). The densification would be done by two mechanisms, the volume distribution or/and the grain boundaries diffusion (Johnson, 1969). At this stage (it is not the scope of the works to discuss the processus that will require further analysis) diffusion in volume therefore in the solid state is a time dependent processus sensitive to the sintering speed. Secondly, the diffusion at interfaces (grain boundaries) would be plausible since observations with Scanning Electron Microscopy shows that the pseudo-hexagonal plate structure seems to be retained during densification and after transformation of metakaolin. This morphology as well as small size (about 2 µm) would favor obtaining a large specific surface of solid-solid contacts necessary for efficient diffusion at grain boundaries. It is also necessary to consider the simultaneous existence of a mechanism of densification by viscous flux which has been demonstrated in the literature during the sintering of compacts or kaolinite-based
<figure><img src="image_441.png" /><figcaption>**FIGURE 6** | Printed pawns after sintering, first campaign with Step 2: annealing at 500°C, 1 h. The illustrations (9, 12, and 21 mm) show the dimensional stability of the parts. Right side: facies of rupture after compression test (1 mm/min 25% deformation, T 23°C) on 9 mm pawn. Medium side: size of the inter-filament on 12 mm pawn. Printed-layer thickness of 1 mm. Cylinders of 20 mm in diameter.</figcaption></figure>

ceramics (Cambier et al., 1984). This point considers additional open questions, the application of the speed method, the constant annealing, the growth of primary mullite above 1,150°C and, but around 1,250°C the crystallization of new phases (cristobalite and secondary mullite).

## 3D Printed Part-Scale

Printed pawns of different size (3, 9, 12, 15, 18, and 21 mm corresponding to the superposition of the number of layers) after sintering were observed. Bearing at Step 2: 500°C② (campaign 1) was of interest because the pawns did not break during sintering, whatever their size, with dimension stability and conformity. **Figure 6** shows 9, 12, 21-layered pieces fabricated using kaolin clay, after the first sintering campaign. The inter-layers coexist with one another much more smoothly than the top adhesion layer does (**Figure 6**, 12-pawn Top and Bottom and graph), thickness variation 970–730 µm. The layer on the top adhesion is flush, thickness of 250 µm.

While observing the pawns, their difference to the touch clearly shows that the microstructural arrangement expected during baking has taken place, even with a color change from beige to light pink. The kaolin is then supposed to be found in a frozen metakaolin-mullite transient state (**Figure 5B**).

To examine the dimensional stability idea more closely, optical microscopy was used to look at the cross-section of a rupture

facies of a pawn after a compression test (1 mm/min, 25% deformation, 23°C) (**Figure 6**, 9-facies of rupture). The facies of rupture revealed that kaolin clay indeed well settled, in the bulk and outer shell. The interlayer connection shows no defaults, only few small individual porosities are found within layer thickness itself. The hardness of the bulk material has been checked, by indentation Knoop test (10's in force loading 1 Kgf, 23°C); a HK of [300–350] was revealed, typical values for hard materials. Enough vitrification has taken place to give a moderate degree of hardness, the clay is considered baked.

Nota, a second sintering campaign, with a Step 2: annealing at 700°C②, 1 h was also experimented. This second campaign confirmed that the first level annealing step at 500°C is essential to allow the water to fully evaporate. At 700°C, some water would still be confined and blocked before processing the kaolin. The pressure exerted by the water during heating would cause the fissuration or even the explosion of the object with temperature increasing. It also allowed studying the influence of the thickness of the pawn on the quality of sintering. The 3 mm high pawn became curled after sintering, no crack is visible showing strong inter-layer contact (**Figure 7**, 3-pawn). The pawns of 9, 12, 15, and 18 mm exhibit slight cracks. This clearly evinces concentric failure along the filament interfaces and inter-layer distortion of the outer shell (**Figure 7**, 18-pawn Top and Bottom and graph), thickness variation range of 800–700 µm. The layer on the top adhesion is very flush, thickness
<figure><img src="image_442.png" /><figcaption>**FIGURE 7** | Printed pawns after sintering, second campaign with Step 2: annealing at 700°C, 1 h. The illustrations (3, 18, and 21 mm) show the three tendencies of dimensional loss: 3 mm with curled deformation, 18 mm with cracks formation and inter-layer delamination batches, and 21 mm broken parts with a zoom showing voids between inter-filament inside the bulk part due to the printing 3D process. Printed-layer thickness of 1 mm. Cylinders of 20 mm in diameter.</figcaption></figure>

less than 250 µm. We can speak about anisotropic shrinkage considering that the border effect or inter-layer discontinuity makes these areas the weakest points of the printed paste, where local constraints are concentrated while water evaporates. The thickest pawn of 21 mm completely exploded (Figure 7, 21-pawn and facies of rupture). The facies of rupture revealed additional weakest mesoscale points between inter-layers with rectangular sometimes accentuated air gaps in the printed-bulk; sometimes compact area. Noting that these cohesive peel failures resulted from each individual pass of the printer. Internal stresses are set up in the ceramic if it loses water too fast according to the kinetics of vitrification of the kaolin. This information is very useful. The quality of the print depends on the type of structure you want to print.

# A 3D-PADM CERAMIC AS A PROOF OF CONCEPT

## From a CAD Idea and Using Hydrated Kaolin

Of these observations, an artistic pottery sculpture was designed, printed, dried and sintered. The main steps of the AM process are combined in Figure 8. Step 1. Software—3D model is designed by using the specialized computer-aided gmesh (version 4.1.4) modeling software. It is digitized and sliced into a number of consecutive layers (Cura 3D version 4.4.1). Step 2. Software to Hardware—The file is subsequently converted to an appropriate format for printing and exported as a G-code file to be transferred to the PADM set up. Step 3. Optimal process parameters—the printing conditions are adjusted depending on the printing methodology, i.e. slurries, pressure, air temperature, and nozzle material flow.

## To an Artistic Pottery

Step 4. 3D printing—The designed structure is processed. Step 5. Drying and sintering—The structure is dried and sintered at 500 and 1,200°C (campaign 1, Figure 5A, Steps 2 and 4). Successfully a pottery part was fully designed. Kaolinite of low plasticity undergoes no shrinkage upon drying, the printed sculpture did not show deformation. This method permits the fabrication of fully customized and personalized products with a geometrical complex structure in an economical manner (mostly for limited productions). It is worth mentioning that ceramics PADM offers environmental responsible implications, kaolinite slurries can be recycled before complete sintering.

# SUMMARY AND OUTLOOK

Manufacturing industries are being forced to create and deliver quality products while decreasing their environmental impact. Kaolinite dough taken in its pristine state is offered to bring insight to help frame and direct future research around the technology.

Specifically:

* Applying PADM technology to print complex structures in geo-materials does not rely on the concentration of manufacturing processes of the product in a single step. It is associated with the synergetic combination of the states of the materials:

- properties of wet materials prior to solidification (i.e. pumpability and extrudability), i.e. the essential constituents of ceramics (design of slurries with enhanced formulation value),
<figure><img src="image_443.png" /><figcaption>Step1. Software: to sliced-design</figcaption></figure>

<figure><img src="image_444.png" /><figcaption>Step1. Software: from computer
aided design</figcaption></figure>

<figure><img src="image_445.png" /><figcaption>Step 2. Software to Hardware:
transfert to 3D printer</figcaption></figure>

<figure><img src="image_446.png" /><figcaption>Step 3. Optimal materials
& process parameters</figcaption></figure>

<figure><img src="image_447.png" /></figure>

<figure><img src="image_448.png" /><figcaption>Pottery</figcaption></figure>

<figure><img src="image_449.png" /><figcaption>Step 4 3D printing</figcaption></figure>

<figure><img src="image_450.png" /><figcaption>Step 5. Drying and Sintering</figcaption></figure>

<caption>**FIGURE 8** | From a CAD idea and using hydrated kaolinite to an artistic pottery.</caption>

- binding mechanisms and hardened properties of parts
during 3D deposition of layers

- dehydration

- "high temperature oven-dry";

* together with the set of PADM parameters such as:

- material flow

- printing speed

- layer thickness.

* Achievement of geometric conformity is demonstrated. Ceramic parts have no cracks nor shape-instability (warping) before and after sintering.

There are still research gaps to be explored in the optimization and sustainability of AM processes. Challenges in the 3D printing of ceramics are not only attributed to controlling slurries properties to have sufficient workability and open time for extrusion but also the metrics to improve the precision and reliability of the manufacturing process, the efficient post-processing structural properties e.g. strength, inter-layer adhesion, deformation, build-ability.

Finally the non-toxic and reusable materials with the
development of new innovative materials are nowadays a
fundamental consideration. Future works will concern the
formulation of geo-materials especially applying for 3D-printing.

DATA AVAILABILITY STATEMENT

The original contributions presented in the study are included in
the article/Supplementary Material, further inquiries can be
directed to the corresponding authors.

AUTHOR CONTRIBUTIONS

SAEB: literature review, designing and drafting the content of the
manuscript. AB: reading the content of the manuscript. LJ with
SAEB and AB: study's work.

FUNDING

Delta WASP 2040 TURBO Printer was bought through the ANR
project 'Western Mediterranean Impressed Wares: An
interdisciplinary research on Early Neolithic-CIMO' (ANR-14-
CE31-0009).

ACKNOWLEDGMENTS

The authors would like to acknowledge the MINES Paris
Engineering School of PSL Research University, and especially
Michel Bellet, Professor in charge of the Specialty “Materials” for
graduate students, for giving us the nice opportunity to propose and work on this exploratory subject. The authors would also thank Suzanne Jacomet, Gilbert Fiorucci, Thierry Colin and

Cyrille Collin, respectively for SEM, sintering with compression, and hardness technical support. CNRS institution is also acknowledged.

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**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2021 Boyer, Jandet and Burr. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
<figure><img src="image_451.png" /></figure>

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# Characterization of Fly Ash Alkali Activated Foams Obtained Using Sodium Perborate Monohydrate as a Foaming Agent at Room and Elevated Temperatures

Lidija Korat and Vilma Ducman*

Slovenian National Building and Civil Engineering Institute (ZAG), Ljubljana, Slovenia

Alkali activated foams have been extensively studied in recent years, due to their high performance and low environmental footprint compared to foams produced via other methods. Three types of fly ash differing in chemical and mineralogical composition and specific surface were used to synthesize alkali activated foams. Sodium perborate monohydrate was added as a foaming agent and sodium dodecyl sulphate as a stabilizing agent. Foams were characterized at room temperature and after exposure to an elevated temperature (1,000°C). Densities from 1.2 down to 0.3 g/cm³ were obtained, depending on the type of fly ash and quantity of foaming agent added. Correspondingly, compressive strength ranged from 1 to 6 MPa. Comparing all three fly ashes the most favorable results, in terms of density and corresponding compressive strength, were achieved from the fly ash with the highest amounts of SiO₂ and Al₂O₃, as well as the highest amorphous phase content i.e., RI fly ash. Furthermore, after firing to 1,000°C, the density of samples prepared using fly ash RI remained approximately the same, while the compressive strength increased on average by 50%. In the other two types of fly ash the density increased slightly after firing, due to significant shrinkage, and compressive strength increased by as much as 800%. X-ray powder diffraction analysis confirmed the occurrence of a crystallization process after firing to 1,000°C, which resulted in newly formed crystal phases, including nepheline, sodalite, tridymite, and gehlenite.

**Keywords:** alkali activated foams, fly ash, sodium perborate monohydrate, microstructural evaluation, mechanical properties

## INTRODUCTION

Alkali-activated (AA) materials have been widely investigated in recent years (Provis and Bernal, 2014; Provis et al., 2015; Luukkonen et al., 2018; Provis, 2018) within the context of the construction and building industry. More specifically, research in the area of sintered and non-sintered foam materials, based on the process of alkali activation, is currently extremely active (Zhao et al., 2010; Chen et al., 2011; Hajimohammadi et al., 2017a; Rabelo Monich et al., 2018; Rincon, 2019). The aforementioned foams present one of the most promising types of AA materials, given their high performance and low environmental footprint in comparison to foams produced by other methods (Dal Pozzo et al., 2019). They are known to offer relatively high strength and good insulation, as well

## OPEN ACCESS

**Edited by:**
Jacqueline Anne Johnson,
University of Tennessee Space
Institute (UTSI), United States

**Reviewed by:**
Ping Duan,
China University of Geosciences,
Wuhan, China
Sriramya Duddukuri Nair,
Cornell University, United States

*Correspondence:
Vilma Ducman
vilma.ducman@zag.si

**Specialty section:**
This article was submitted to *Structural Materials*, a section of the journal *Frontiers in Materials*

**Received:** 15 June 2020
**Accepted:** 18 August 2020
**Published:** 03 September 2020

**Citation:** Korat L and Ducman V (2020) Characterization of Fly Ash Alkali Activated Foams Obtained Using Sodium Perborate Monohydrate as a Foaming Agent at Room and Elevated Temperatures. *Front. Mater.* 7:572347. doi: 10.3389/fmats.2020.572347
as being non-combustible up to high temperatures (Bajare et al., 2019). The key element likely to influence the use of alkali-activation is the local availability of appropriate raw materials (precursors and activators) (Provis, 2018), typically metakaolinite, slags and ashes. The wide availability of fly ash has forced researchers to find novel applications suitable for larger-scale production, one option being lightweight foams with a bulk density of <1 g/cm³. The mechanical properties, i.e., compressive strength of such lightweight materials depends on the density of the foamed AA materials, and is found to range between 1 and 10 MPa for densities from 0.36 to 1.4 g/cm³ (Zhao et al., 2010; Masi et al., 2014a; Sanjayan, 2015). Foaming agents for producing AA foams are selected from various compounds, especially hydrogen peroxide (Masi et al., 2014; Abdollahnejad et al., 2015; Palmero et al., 2015; Hajimohammadi et al., 2017b; Szabo, 2017), NaOCl (Nyale et al., 2013; Boke, 2015; Łach et al., 2018), metal powders (e.g., elemental aluminum or zinc) (Masi et al., 2014a; Hlaváček et al., 2015; Kamseu et al., 2015; Hajimohammadi et al., 2017c), perborate (Abdollahnejad et al., 2015; Davidovits, 2015), as well as silica fume and some others (Zhao et al., 2010; Henon et al., 2013; Guo et al., 2016; Kioupis et al., 2018). Foaming agents react or decompose in alkaline conditions and form gas which remains trapped inside the structure, thus expanding the material, increasing its volume and forming a cellular structure (Masi et al., 2014a; Sanjayan, 2015). Dried AA foams (already successfully stabilized at room or slightly elevated temperature) (Masi et al., 2014b; Sanjayan, 2015; Kioupis et al., 2018) can also be subjected to heat treatment at higher temperatures, for example in a furnace at temperatures between 800 and 1,000°C, in order to improve mechanical properties (Chen et al., 2011; Badanoiu et al., 2015; Carabba et al., 2017; Bajare et al., 2019; Rincon Romero, 2019). Rincon Romero et al. (2019) presented glass-ceramic foams, based on soda lime glass and coal fly ash, where dried foams were later subjected to heat treatment at 700-900°C. The glass-ceramic foams produced had a high porosity (58-78%), a thermal conductivity of 0.163 ± 0.005 W m⁻¹ K⁻¹ (measured at 25°C, samples treated at 800°C), and reasonable compressive strength (1.8-8.7 MPa). The results of Abdollahnejad et al. (2015) show that sodium perborate outperforms hydrogen peroxide, leading to a lower overall thermal conductivity of foam materials (approximately 0.1 W m⁻¹ K⁻¹) and a compressive strength of approximately 6 MPa. Zhao et al. (2010) fabricated foams with coal fly ash and 13 wt% foaming agent (sodium dodecyl benzene sulfonate and gluten) sintered at 1,050°C for 2 h after alkali-activation. The resulting foams exhibited values for water absorption, apparent density, and compressive strength of approximately 126.5%, 0.414 g/cm³, and 6.76 MPa, respectively. Results from other researchers to date have nevertheless shown that the stability of foams at high temperatures (especially with regard to shrinkage) is strongly affected by the composition of raw materials, not only when fly ash is used (Martin et al., 2015) but also when other additives such as metakaolin and slag are present in the AA material (Mierzwinski et al., 2019; Tran et al., 2019). Some of the authors concluded that shrinkage occurs due to the increased density or change in volume induced by the crystallization of new phases, rather than due to a melting

point or viscous creep (Martin et al., 2015). In 2006, Bakharev wrote that fly ash activated with activators containing sodium exhibited shrinkage cracking and a rapid decline in strength at 800°C (but an increase in average pore size), where the loss of strength on firing may have been associated with the deterioration of the aluminosilicate gel and feldspar crystallization. Martin et al. (2015) studied AA material (fly ash with 15% sodium silicate + 85% 10-M NaOH) from 25 to 600°C and their findings showed that the increase in temperature led to an increased compressive strength (with the specimens exhibiting elastic behavior and brittle failure), increased pore size distribution, and the presence of hydroxysodalite (nepheline at 800°C), and that the dimensional stability of the material was affected at temperatures over 1,000°C. According to Chen-Yong et al. (2017), the mechanical strength of foamed AA materials (class F fly ash with hydrogen peroxide) deteriorated to 3 MPa at 400°C but increased up to 11 MPa at 800°C, due to the formation of crystalline phases (nepheline, anorthite, and cristobalite) at higher temperatures, which potentially act as fillers to reinforce the matrix.

The aim of the present research was to evaluate the usability of sodium perborate monohydrate (SPM), which is inexpensive and widely available, as a foaming agent in AA systems. Three different types of fly ash were used, the main difference being their chemical composition, i.e., the content of SiO₂ and Al₂O₃, which might subsequently influence the properties of the foams produced. The performance of foams was determined after curing at 70°C but also assessed after exposure to an elevated temperature (1,000°C), in order to verify their mechanical properties and dimensional stability at an elevated temperature and thus expand the potential applications of inorganic AA foams to areas where resistance to high temperatures is required, such as for the insulation of furnaces, chimneys, and heat resistant doors.

# MATERIALS AND METHODS

## Materials

Three types of fly ash (designated as RI, TS, and TT) originating from different sources were investigated, including two types (RI, TS) originating from coal thermal power plants, and one (TT) derived from the co-combustion of coal and biomass material at a heat and power station. RI fly ash is F type, while the other two are C type fly ash. The Brunauer-Emmet-Teller (BET) specific surface areas and chemical compositions (determined by X-ray fluorescence) of the investigated fly ashes are shown in Table 1. Sodium silicate solution (sodium silicate crystal 0112; produced by Tennants Distribution, Ltd., SiO₂:Na₂O = 1.97, 54.2 mass% aqueous solution) and NaOH (produced by Donau Chemie, 41.7 mass% aqueous solution) were used as activating agents. The foaming agent was SPM (Sodium Perborate Monohydrate produced by Belinka) and the stabilizing agent was sodium dodecyl sulphate (SDS) (produced by Acros Organics).

## Preparation of Alkali-Activated Foams

The investigated AA foams were synthesized from sodium silicate solution, NaOH, and the selected fly ash, to which
<table><caption>**TABLE 1** | BET specific surface area and chemical composition of the different types of fly ash (wt%) which were used in the study of alkali-activated foams.</caption><thead><tr><th>Fly ash</th><th>BET specific surface area (cm²/g)</th><th>Loss on ignition (at 950°C)</th><th>SiO₂</th><th>Al₂O₃</th><th>Fe₂O₃</th><th>CaO</th><th>MgO</th><th>Na₂O</th><th>K₂O</th><th>SO₃</th></tr></thead><tbody><tr><td>RI</td><td>1.7</td><td>2.3</td><td>52.5</td><td>23.3</td><td>7.5</td><td>6.1</td><td>2.5</td><td>0.8</td><td>2.2</td><td>0.7</td></tr><tr><td>TS</td><td>16.7</td><td>5.0</td><td>39.4</td><td>17.0</td><td>8.6</td><td>18.6</td><td>3.9</td><td>0.8</td><td>1.8</td><td>2.7</td></tr><tr><td>TT</td><td>33.4</td><td>10.2</td><td>27.3</td><td>10.1</td><td>16.7</td><td>19.7</td><td>11.1</td><td>0.2</td><td>0.9</td><td>1.5</td></tr></tbody></table>

SPM (as the selected foaming agent) and SDS (as the stabilizing agent) were then added. SDS has already been confirmed as a suitable stabilizing agent in a previous study (Korat and Ducman, 2017). Samples were prepared with varying compositions, the sample designations and sample compositions of the various mixtures of the AA foams are presented in **Table 2**. Three mixtures were prepared for each selected fly ash, with designations 1, 2, and 3 representing the amount of SPM and/or SDS added. The amount of SPM and SDS is equal, determined according to the mixture of water glass, NaOH, and fly ash being investigated. First NaOH was dissolved in the water glass solution, and then this solution was added to the fly ash. Components were then mixed together. Samples were mixed for 5 min in a laboratory mixer. At the end of mixing process the SPM and SDS were added and the foaming process started. The mixture was then placed in molds with dimensions of 20 mm × 20 mm × 80 mm and the samples expanded further. The molds were then placed in an oven for 24 h at a temperature of 70°C. The resulting hardened test specimens were demolded and stored for a further 3 days at a controlled temperature of 20 ± 2°C.

## Thermal Treatment of Alkali-Activated Foams

The investigated AA foams were subjected to thermal treatment using a Porotherm furnace fired to 1,000°C at a heating rate of 600°C/h.

## Methods

The density of the AA foams were calculated by dividing the measured weight of the samples by their dimensions, firstly at the age of 4 days and later after firing. The dimensions, which were also used to determine shrinkage, were measured using digital Vernier calipers (Mitutoyo, Japan), and the sample mass was measured using a laboratory scale (Sartorious, Germany). Reported density values were obtained from the average of three separate measures per sample.

The compressive strength of the test specimens was determined at the age of 4 days and later on after firing, by means of a Toninorm testing machine (Toni Technik, Germany), using a force application rate of 0.005 kN/s. The reported values of mechanical strength were obtained from the average of 10 sample measures per mixture.

The back-scattered electrons image mode of a low vacuum Scanning Electron Microscope (JEOL 5500 LV equipment) was then used to examine the microstructure of selected cross-sections of the hardened foams.

X-ray diffraction analysis was performed using an Empyrean (PANalytical) X-ray diffractometer equipped with Cu Kα radiation. The data was collected at room temperature at an acceleration voltage of 45 kV using a current of 40 mA, from 4 to 70° 2θ in steps of 0.031° 2θ.

For the porosity analysis X-ray micro-computed tomographic imaging was conducted using an Xradia µCT-400" tomograph (XRadia, Concord, CA, United States). The 3D images were produced using a working voltage of 80 kV and an energy of 125 µA, with a spatial resolution of 29 µm, using a ×0.39 magnification optical objective. Regions of interest (ROIs) were selected from the center of the sample, at a size of 15 mm × 15 mm × 15 mm. The projection data for each tomographic scan consisted of 2,000 projection images, which were taken from different view-points with an exposure time of 1 s per projection. These images were reconstructed into 3D tomographic volumes using Avizo Fire 3D-image analysis software. The segmentation procedure and quantification, for the determination of overall porosity and pore size distribution of the hardened foams, was performed using the procedure previously described (Korat et al., 2013).

Thermal conductivity tests on the samples were performed using a Hot Disk TPS 2500S instrument with a Kapton sensor 7577 at a temperature of 22–23°C. The measure was tested on square samples (20 mm × 20 mm × 80 mm), with 10 measurements performed on each samples.

<table><caption>**TABLE 2** | The sample designation and sample compositions of the different mixtures prepared for the investigation (in wt%).</caption><thead><tr><th>Sample designation</th><th>Fly ash</th><th>Water glass</th><th>NaOH</th><th>Addition of sodium perborate monohydrate</th><th>Addition of sodium dodecyl sulfate</th></tr></thead><tbody><tr><td>RI1</td><td>64.9</td><td>9.2</td><td>24.0</td><td>1.0</td><td>1.0</td></tr><tr><td>RI2</td><td>63.6</td><td>9.0</td><td>23.5</td><td>1.9</td><td>1.9</td></tr><tr><td>RI3</td><td>62.4</td><td>8.8</td><td>23.1</td><td>2.8</td><td>2.8</td></tr><tr><td>TS1</td><td>51.7</td><td>15.5</td><td>31.0</td><td>0.9</td><td>0.9</td></tr><tr><td>TS2</td><td>50.8</td><td>15.3</td><td>30.5</td><td>1.7</td><td>1.7</td></tr><tr><td>TS3</td><td>50.0</td><td>15.0</td><td>30.0</td><td>2.5</td><td>2.5</td></tr><tr><td>TT1</td><td>46.8</td><td>14.2</td><td>37.1</td><td>0.9</td><td>0.9</td></tr><tr><td>TT2</td><td>45.9</td><td>13.9</td><td>36.4</td><td>1.9</td><td>1.9</td></tr><tr><td>TT3</td><td>45.1</td><td>13.7</td><td>35.8</td><td>2.7</td><td>2.7</td></tr></tbody></table>
# RESULTS AND DISCUSSION

## Analysis of Samples Cured at 70 and 1,000°C

**Figure 1** shows the AA foams from three types of fly ash (designated as RI, TS, and TT), prepared using a foaming agent and stabilizing agent in the range 0.9–2.8 wt% (designated as 1, 2, and 3). **Figure 1A** depicts cured samples at 70°C and **Figure 1B** shows samples after being thermally treated at 1,000°C. The change in color observed in the thermally treated samples can be explained by oxidation of the iron present in the raw fly ash, presenting from greyish to brownish color at higher temperatures. This explains why samples with TT fly ash are darker and have a more intense brown color (see **Table 1**).

**Table 3** shows the density and compressive strength of the AA foams investigated, after curing at 70°C and when thermally treated to 1,000°C. Sample RI1 exhibits the highest density of all the samples prepared with RI fly ash when cured at 70°C (0.67 g/cm³). As the amount of foaming agent and stabilizing agent added increased, the density of the RI1 sample decreased by more than half, to 0.3 g/cm³. A similar trend is observed for samples from the TS and TT fly ash, but both have a higher density than the samples produced from RI fly ash. Nevertheless, even with the highest amount of foaming and stabilizing agents added, their densities remain above 0.5 g/cm³. Compressive test results for samples from RI, TS, and TT fly ash (cured at 70°C) are also shown in **Table 3** and **Figure 2**. Values are the highest for samples from RI fly ash (around 6 MPa) and begin to decline as the amount of foaming and stabilizing agent added increased, which decreases density. Values for the samples from TT fly ash are slightly higher than for the sample from TS.

Compressive test results for samples from RI, TS, and TT fly ash exposed to temperatures of 1,000°C are also shown in **Table 3**. A tendency to increase compressive strength upon thermal treatment was observed in all samples. After heat treatment, sample RI1 still has the highest density among all three samples from RI fly ash and the lowest amongst all the samples when compared to those produced from the other two types of fly ash.

Compressive strength values for RI fly ash began to decline (from 10 to 2 MPa) as the amount of foaming and stabilizing agent added increased, and the density decreased. A similar trend is observed for samples from TS and TT fly ash. There are no significant differences in density between cured and thermally treated samples for samples from RI fly ash (only a few percent), or for samples from TS fly ash, but there is a big difference for samples from TT fly ash (from 10 to 50%), which can be attributed to the increased density resulting from the significant shrinkage.

Dimensional changes, i.e., shrinkage, occurred in samples after exposure to high temperatures of 1,000°C. For samples with the most foaming agent added, the biggest changes were detected in the direction of length, as shown in **Figure 3**. The results show that the difference in dimensional changes (percentage change in length) is approximately 4% for sample RI3, 9% for sample TS3, and 19% for sample TT3. Lin's work (Lin et al., 2009) revealed thermal shrinkage values of 45% for geopolymer and 30% for geopolymer with the addition of 10 vol% alumina filler, when heated to 1,000°C. Shrinkage could be attributed to the presence of alkaline earth flux and other oxides in the fly ash and activators, which enter the glassy phase, decreasing the viscosity of these samples resulting in higher shrinkage (Dana et al., 2004), or, as suggested by Martin, to volume changes caused by the formation of new crystalline phases (Martin et al., 2015). Weight loss after exposure to temperatures at 1,000°C is approximately 10% for sample RI3, 14% for sample TS3, and 20% for sample TT3, as shown in **Figure 3**.

Thermal conductivity tests were conducted on AA foams cured at 70°C, the results of which are listed in **Table 4**. Thermal conductivity properties relative to density are given in **Figure 4**.

The results show the thermal conductivity of RI samples to be 0.205 W/m K for the least porous sample (RI1), falling to 0.123 W/m K for the most porous one (RI3; **Table 4**), where the range of apparent densities for the corresponding samples obtained is 0.67–0.33 g/cm³. The results for TS samples show that thermal conductivity ranges from 0.151 W/m K for the most

<figure><img src="image_452.png" /></figure>

<figure><img src="image_453.png" /><figcaption>**FIGURE 1** | Alkali-activated foams (A) cured at 70°C and (B) thermally treated to 1,000°C (two samples per mixture).</figcaption></figure>
<table><caption>**TABLE 3** | Density and compressive strength of the alkali-activated foams investigated when cured at 70°C and thermally treated at 1,000°C (the measured standard deviations are shown in brackets).</caption><thead><tr><th rowspan="2"></th><th colspan="2">Density (g/cm³)</th><th colspan="2">Compressive strength (MPa)</th></tr><tr><th>Cured samples</th><th>Thermally treated</th><th>Cured samples</th><th>Thermally treated</th></tr></thead><tbody><tr><td>RI1</td><td>0.67 (0.01)</td><td>0.68 (0.01)</td><td>6.33 (0.79)</td><td>9.73 (0.28)</td></tr><tr><td>RI2</td><td>0.43 (0.02)</td><td>0.43 (0.01)</td><td>1.93 (0.26)</td><td>3.02 (1.25)</td></tr><tr><td>RI3</td><td>0.33 (0.03)</td><td>0.34 (0.02)</td><td>1.02 (0.12)</td><td>2.17 (0.04)</td></tr><tr><td>TS1</td><td>1.02 (0.02)</td><td>0.95 (0.01)</td><td>3.61 (0.46)</td><td>12.89 (0.51)</td></tr><tr><td>TS2</td><td>0.75 (0.01)</td><td>0.75 (0.01)</td><td>1.81 (0.44)</td><td>8.15 (0.69)</td></tr><tr><td>TS3</td><td>0.57 (0.01)</td><td>0.63 (0.01)</td><td>1.20 (0.31)</td><td>4.98 (0.52)</td></tr><tr><td>TT1</td><td>1.18 (0.02)</td><td>1.03 (0.03)</td><td>4.12 (0.33)</td><td>9.21 (0.42)</td></tr><tr><td>TT2</td><td>0.87 (0.01)</td><td>0.97 (0.03)</td><td>2.31 (0.35)</td><td>7.19 (0.96)</td></tr><tr><td>TT3</td><td>0.71 (0.02)</td><td>1.07 (0.02)</td><td>2.23 (0.39)</td><td>15.51 (2.84)</td></tr></tbody></table>

porous sample (TS3) to 0.262 W/m K for the least porous one (TS1), with the range of apparent densities being 0.57–1.02 g/cm³. The thermal conductivity measured in the TT samples treated at 70°C ranges from 0.232 to 0.363 W/m K, with a density of 0.71 g/cm³ in the least porous sample (TT3). As expected, the higher the density, the higher the thermal conductivity (Figure 4). The overall behavior of the thermal conductivity is generally controlled by the pore volume, as the most porous sample exhibits the lowest thermal conductivity, as has been found by other researchers (Henon et al., 2013).

## Microstructural Examination of Samples Cured at 70°C

Microstructural characteristics have a significant influence on material characteristics, especially thermal and mechanical properties, and because the pore structure inside the material strongly affects the cell (solid) structure, tests by scanning electron microscope (SEM) and X-ray micro-computed tomography (XCT) were performed to gain a more detailed understanding. Firstly, Figure 5 shows the microstructural evolution of AA foam samples attained through SEM before exposure to higher temperatures. It is evident that samples from RI fly ash have a porous structure with evenly distributed round pores, clearly developed and connected to each other. No micro-cracks are observed within and nearly all of them are cell pores with

no mineral crystallization observed inside or on the walls. Contrarily, samples from TS and TT fly ash have a different pore distribution, especially TT foamed samples, which exhibit high structural heterogeneity and contain hardly any rounded pores. These differences in the type of porosity are most probably a reflection of the previously mentioned thermal conductivity—for example TS2 samples have a higher density than RI1 (0.75 and 0.67 g/cm³, respectively) and exhibited slightly lower thermal conductivity (0.200 vs. 0.205 W/m K), and sample TT3, with a density of 0.71 g/cm³, showed a thermal conductivity of 0.232 W/m K.

SEM images can only, however, provide the pore characteristics on a single cross section of the specimen, meaning representative shapes or the number of pores cannot be appropriately characterized. More detailed microstructural 3D data can be obtained from high resolution 3D images achieved through X-ray microcomputed tomography. 3D data is displayed on the basis of threshold segmentation and 3D reconstructions, and where the pore structure quantitative visualization and characterization of pore structure of AA foams cured at 70°C are presented on Figure 6. Figure 6 illustrates the ROI (Region Of Interest) of the analyzed AA foams cured at 70°C, including the grayscale image and the biggest pores (by volume), and after the watershed algorithm, where the labeled images show distinct colors for similar pores (purple, green, and yellow pores, as determined according to the EqDiameter). The overall porosity of samples is presented in Table 5.

Overall porosity, the local volume ration of the cell (solid) in the ROI, is the highest for samples with RI fly ash. The higher percentages for the RI2 and RI3 samples are illustrated by the porous nature of the AA foams analyzed, due to the addition of the largest amount of agents. RI2 and RI3 both consist of a defined connected pore network with larger individual pores, some of which are even larger than 1 mm. A similar effect is observed in samples from the TS fly ash, but with a smaller pore size and porosity up to approximately 40% in TS3. In the case of TT, the effect of poor porosity development is more notably pronounced. The differences between TT1 and TT3 can be seen from the SEM images (Figure 5), which is coherent to the structure detected by JACK scanning, revealing deficient pore development. The structure of samples is similar to one another, with the porosity ranging from 10 to 20%.

<figure><img src="image_454.png" /><figcaption>**FIGURE 2** | Density and compressive strength of the alkali-activated foams investigated when cured at 70°C and thermally treated at 1,000°C.</figcaption></figure>
<figure><img src="image_455.png" /><figcaption>**FIGURE 3** | Changes in sample dimensions of the alkali-activated foams investigated, expressed in length and weight loss.</figcaption></figure>

3D geometrical parameters (especially equivalent diameter) are relevant for the determination of pore morphology and can indicate the thickness of the asymmetric and/or elongated pore structures. Equivalent diameter is reported as the diameter of a sphere having the same volume as the feature (Figure 7), showing the percentage of pores (as well as number and volume) distributed in each sample.

As shown in **Figure 7A**, the highest number of pores in the RI samples is in the 100–850 µm pore diameter range, where RI1 has the larger number of pores ranging in diameter from 100 to 500 µm, due to the addition of the least amount of foaming agent. The number of pores with a diameter below 100 µm is similar across all three RI samples, with the RI1 sample containing the highest number of pores in this category. In TS samples, the greatest number of pores fall in the pore diameter range below 500 µm (Figure 7D), where TS1 has the largest number of pores with a pore diameter less than 100 µm, due to the addition of a smaller amount of foaming agent. Despite the larger percentage of foaming agent added to samples TS2 and TS3, the number of pores in the 100–500 µm pore diameter range is still high in sample TS1. TT samples (Figure 7G) have the highest number of pores with a diameter less than 500 µm, compared to other types of fly ash, where sample TT1 having the largest number of pores with a diameter less than 100 µm.

In the RI sample, most pores are in the 100–850 µm pore diameter range (Figure 7B), where RI1 has the largest amount/percentage of pores in the pore diameter range 100–500 µm, reaching almost 70%. Due to the addition of a greater amount of foaming agent, the percentage of number of pores in the range 500–850 µm is higher for samples RI2 and RI3, with approximately 38% of pores in sample RI2 falling in this range. Sample RI3 has a larger percentage of pores with a diameter above 500 µm, approximately 60%. The percentage of

pores with a diameter below 100 µm is about 10% for all three RI samples. In TS samples most pores are below 500 µm in diameter, with sample TS1 having the largest percentage of pores with a diameter less than 100 µm, reaching almost 70% (Figure 7E). Due to the addition of a greater amount of foaming agent, the percentage of number of pores in the diameter range 100–500 µm is higher (above 60%) for samples TS2 and TS3. Only sample TS3 has a larger percentage of pores in the pore diameter range above 500 µm (about 35%) and the least in the pore diameter range below 100 µm (about 10%). In the TT samples, most pores are in the pore diameter range below 500 µm (Figure 7H). Due to the addition of a greater amount of foaming agent, the percentage of number of pores in the range 100–500 µm is higher only for TT2 (almost 60%). The percentage of pores in the pore diameter range below 100 µm is large, especially in the RI3 sample (70%).

The volume of pores in the pore diameter range from 100 to 850 µm is between 40 and almost 90% for all three RI samples (Figure 7C), with the greatest volume of pores observed in the RI1 sample, showing the maximum volume of pores for the RI1 sample. In the RI2 sample the volume of pores in this pore diameter range is 60%, with 40% of pores having a diameter greater than 850 µm. In the RI3 sample the largest percentage of pore volume has a pore diameter of greater than 500 µm, almost 95%. TS samples have the greatest volume of pores in the pore diameter range between 100 and 850 µm (Figure 7F), whereas sample TS1 has almost 80% in the pore diameter range between 100 and 500 µm, and sample TS2 more than 80%. In the TS3 sample the maximum volume of pores is in the pore diameter range above 500 µm (75%). In samples TT1 more than 70% of the volume of pores lie in the pore diameter range from 100 to 500 µm (Figure 7I). A high volume of pores in this pore diameter range is also seen in the TT2 sample (80%), whilst in sample TT3

<table><caption>**TABLE 4** | Thermal conductivity of the investigated alkali-activated foams cured at 70°C (the measured standard deviations are shown in brackets).</caption><thead><tr><th>Sample designation</th><th>RI1</th><th>RI2</th><th>RI3</th><th>TS1</th><th>TS2</th><th>TS3</th><th>TT1</th><th>TT2</th><th>TT3</th></tr></thead><tbody><tr><td>λ (W/m K)</td><td>0.205 (0.006)</td><td>0.143 (0.002)</td><td>0.123 (0.007)</td><td>0.262 (0.008)</td><td>0.200 (0.002)</td><td>0.151 (0.003)</td><td>0.363 (0.006)</td><td>0.303 (0.005)</td><td>0.232 (0.007)</td></tr></tbody></table>
<figure><img src="image_456.png" /><figcaption>**FIGURE 4** | Thermal conductivity properties vs. density in the analyzed alkali-activated foams cured at 70°C (RI, TS, and TT fly ash used, 1, 2, and 3 represent amount of foaming and stabilizing agent used).</figcaption></figure>

the maximum percentage of the volume of pores is in the pore diameter range above 850 µm, due to poor porosity development.

Another 3D geometrical parameter relevant for the determination of pore morphology is sphericity, which describes how spherical an object is, or how rounded it is. Sphericity is a measure of the degree to which an object approximates the shape of a sphere, where the sphericity of 1 represents a perfect spherical shape. X-ray computer tomography has given the possibility of evaluating the results and the calculated sphericity values, as a function of pore size, are shown in **Figure 8**. The sphericity of the samples cured at 70°C is also given in **Table 5** and represents the quantitative statistical data measured from the reconstructed images. The average value, and minimal and maximal sphericity of the tested samples, are

also written in **Table 5**. This analysis was conducted on pores larger than 850 µm diameter, as presented in **Figure 8**.

The results for RI samples thermally treated at 70°C show that the majority of pores were close to spherical in shape as their sphericity values are above 0.50, between 0.55 and 0.95, with no huge difference between mixtures (Figure 8A). There are less spherical pores, i.e., with a sphericity smaller than 0.65, in the TS1 samples and more spherical pores in TS2 and TS3 (up to 0.85). The sphericity of the pores in TT samples thermally treated at 70°C is relatively low (0.25-0.75), confirming their irregular shape, especially for sample TT3, which has the lowest sphericity due to poor pore shape (Figure 8C). It is also observed that the sphericity is equal between mixtures in RI samples, whereas sphericity is seen to slightly change between the mixtures in TS samples (mixtures with a higher addition of foaming agent have higher sphericity values) and in TT samples there are huge changes between the mixtures. Notably, no large differences in the sphericity of samples occur due to an increase in the equivalent diameter of pores.

## Phase Changes by X-Ray Powder Diffraction

X-ray powder diffraction (XRD) (Figure 9) revealed the variable phase composition of the studied ashes and AA foams cured at 70 and 1,000°C (results only for samples with the smallest and largest addition of agents). Based on the height of the background presenting the amorphous phase it can be seen the RI fly ash contains the highest amount of amorphous phase of all three precursors, followed by the TT fly ash, and then the TS. The amount of amorphous phase was determined in a previous study (Kramar and Ducman, 2018), as 84.4% for RI, 71.7% for TT, and 70.0 for TS (designated in the previous research as K3, K1, and K2, respectively). The main crystalline phases identified in the designed geopolymers were similar to those found in the raw precursor, such as quartz and hematite, which were also present in all samples cured at 70°C. Lime, anhydrite, and gehlenite were

<figure><img src="image_457.png" /><figcaption>**FIGURE 5** | Representative SEM image of alkali-activated foams cured at 70°C (RI, TS, and TT fly ash, 1 and 3 represent the amount of foaming and stabilizing agents used).</figcaption></figure>
<figure><img src="image_458.png" /><figcaption>**FIGURE 6** | Region of interest of the analyzed alkali-activated foams cured at 70°C (RI, TS, and TT fly ash, 1, 2, and 3 represent the amount of foaming and stabilizing agents used) presenting grayscale images and the biggest main pores. Colors denote a range of size classes with respect to EqDiameter: purple (0.85–1 mm), green (1–1.6 mm), and yellow (more than 1.6 mm).</figcaption></figure>

found in all samples, with the exception of one from RI ash (Figure 9). It is known that an increase in quartz concentration positively influences mechanical properties, due to the capacity of its particles to create barriers for crack propagation (Nergis et al., 2020). Brownmillerite, a hydraulic phase, which could contribute to the increase in strength, was present in the samples with TT ash derived from the co-combustion of coal and biomass, while feldspar was identified in samples with TS ash, periclase in RI and TT ashes and mullite in RI samples. Mullite is a less commonly encountered compound formed from aluminum, silicon, and oxygen, which increases the refractivity of a geopolymer due to its very high melting temperature of 1,840°C (Nergis et al., 2020). As written by Diaz et al. (2010), amorphous compounds are easier to dissolve than crystalline compounds during the first step of geopolymerization (dissolution of species), yielding higher amounts of reactive SiO₂ and Al₂O₃ to combine during the transportation/

<table><caption>**TABLE 5** | Porosity by SCT of the investigated alkali-activated foams cured at 70°C and thermally treated at 1,000°C.</caption><thead><tr><th rowspan="3"></th><th colspan="2">Porosity by SCT (%)</th><th colspan="6">Sphericity by SCT (due to sample diameter range above 850 µm)</th></tr><tr><th rowspan="2">Cured samples</th><th rowspan="2">Thermal treated</th><th colspan="3">Cured samples</th><th colspan="3">Thermal treated</th></tr><tr><th>Average</th><th>Min</th><th>Max</th><th>Average</th><th>Min</th><th>Max</th></tr></thead><tbody><tr><td>RI1</td><td>39.7</td><td>32.2</td><td>0.71</td><td>0.50</td><td>0.94</td><td>0.70</td><td>0.54</td><td>0.90</td></tr><tr><td>RI2</td><td>66.4</td><td>56.5</td><td>0.81</td><td>0.54</td><td>0.96</td><td>0.78</td><td>0.57</td><td>0.95</td></tr><tr><td>RI3</td><td>76.5</td><td>72.4</td><td>0.83</td><td>0.60</td><td>0.95</td><td>0.83</td><td>0.58</td><td>0.95</td></tr><tr><td>TS1</td><td>18.3</td><td>15.7</td><td>0.50</td><td>0.28</td><td>0.65</td><td>0.47</td><td>0.33</td><td>0.56</td></tr><tr><td>TS2</td><td>19.8</td><td>18.7</td><td>0.72</td><td>0.56</td><td>0.90</td><td>0.58</td><td>0.40</td><td>0.80</td></tr><tr><td>TS3</td><td>38.9</td><td>38.0</td><td>0.63</td><td>0.46</td><td>0.83</td><td>0.56</td><td>0.38</td><td>0.74</td></tr><tr><td>TT1</td><td>16.4</td><td>12.7</td><td>0.53</td><td>0.32</td><td>0.76</td><td>0.64</td><td>0.53</td><td>0.75</td></tr><tr><td>TT2</td><td>15.5</td><td>17.9</td><td>0.76</td><td>0.57</td><td>0.88</td><td>0.67</td><td>0.51</td><td>0.86</td></tr><tr><td>TT3</td><td>15.9</td><td>15.0</td><td>0.32</td><td>0.20</td><td>0.51</td><td>0.29</td><td>0.20</td><td>0.41</td></tr></tbody></table>

coagulation phase of the geopolymeric reaction, therefore resulting in a higher degree of geopolymerization and consequently higher mechanical strength. When the sodium based alkalis are used with fly ash, the reaction product in the system is associated with the formation of sodium aluminosilicate gel (Bhagath Singh and Subramaniam, 2016).

## Microstructural Evolution of Samples Heat Treated at 1,000°C

Figure 10 shows the microstructural evolution of AA foam samples after exposure to higher temperatures at 1,000°C, attained through SEM. Samples with RI fly ash still have a nice porous structure with evenly distributed round pores, connected and with no micro-cracks, whereas samples with TT fly ash exhibit poor microstructural evolution. Figure 10 shows that no noticeable changes were observed in samples after exposure to higher temperatures. The more detailed results of X-ray microcomputed tomography further confirm the findings obtained by SEM, as presented in Figure 10.

Overall porosity determined by SCT is still the highest for samples with RI fly ash, followed by TS samples and TT with poor porosity development (Table 5). The difference in overall porosity between cured samples and samples exposed to higher temperatures is noticeable, varying between samples from 1 to 10%.

Porosity determined by SCT is visually presented in Figure 11, depicting a range of size classes above 0.85 mm, and pore size distribution of the samples is shown in Figure 12. The pore diameter ranges and the number of pores observed in RI samples show a similar trend to the same samples cured at 70°C, but the number of pores in the 100–850 µm pore diameter range, as well as below 100 µm, is smaller in samples cured at 1,000°C (Figure 12A). TS samples cured at 1,000°C also
<figure><img src="image_459.png" /></figure>

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<figure><img src="image_461.png" /></figure>

<figure><img src="image_462.png" /></figure>

<figure><img src="image_463.png" /></figure>

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<figure><img src="image_466.png" /></figure>

<figure><img src="image_467.png" /><figcaption>**FIGURE 7** | Graphical representation of the pore size distribution of the analyzed alkali-activated foams cured at 70°C (RI, TS, and TT fly ash used, 1, 2, and 3 represent the amount of foaming and stabilizing agents used; (A)–(C) RI fly ash, (D)–(F) TS fly ash, (G)–(I) TT fly ash), determined by means of X-ray micro-tomography. The results present the number, percentage and volume of pores with respect to pore size diameter (EqDiameter).</figcaption></figure>

show similar porosity development to samples cured at 70°C, where sample TS1 cured at 1,000°C has large amount of pores in the pore diameter range below 100 µm (Figure 12D), although still less than samples cured at 70°C. Despite the greater percentage of foaming agent added to samples TS2 and TS3, the number of pores in sample TS3 (in the 100–500 µm pore diameter range) is still high, unlike in sample TS3 cured at 70°C. The highest number of pores for TT samples is in the pore diameter range below 500 µm (Figure 12G), where sample TT1 still has the largest number of pores in the pore diameter range below 100 µm, due to the addition of a smaller amount of foaming agent. A similar trend to in samples cured at 70°C can be seen, although a smaller number of pores with a pore diameter below 100 µm are present in sample TT1 when cured at 1,000°C. Despite the larger percentage of foaming agent added to samples TT2 and TT3, TT1 has the highest number of pores in the pore diameter range between 100 and 500 µm, unlike in the sample cured at 70°C.

In R1 samples cured at 1,000°C the majority of pores are in the pore diameter range 100–850 µm (Figure 12B), where RI1 has the largest percentage of pores in the pore diameter range 100–500 µm, reaching almost 74%, which is more than in those samples cured at 70°C. Due to the higher amount of foaming agent, the percentage of pores in the range 500–850 µm is higher in RI2 and RI3 samples, about 33 and 10%, respectively, which is slightly lower than in those samples cured at 70°C. The percentage of pores in the pore diameter range below 100 µm is approximately 10% for all three RI samples. In TS samples cured at 1,000°C (Figure 12E) the majority of pores have a diameter below 500 µm, with samples TS1 and TS2 having a large percentage of pores in the pore diameter range below 100 µm, approximately 54 and 35%, respectively. A similar trend as in TS samples cured at 70°C was observed, but with a smaller amount of pores. Due to the higher addition of foaming agent in samples TS2 and TS3, the percentage of pores in the 100–500 µm range is higher (above 60%) when cured at 1,000°C compared to when
<figure><img src="image_468.png" /><figcaption>**FIGURE 8** | Graphical representation of the sphericity values as a function of pore size of the analyzed alkali-activated foams cured at 70 (A)–(C) and 1,000°C (D)–(F), determined by means of X-ray micro-tomography.</figcaption></figure>

cured at 70°C. Sample TS3 has a larger percentage of pore sizes both in the pore diameter range 100–500 µm about (approximately 60%) and in the pore diameter range below 100 µm (approximately 30%), unlike the samples cured at 70°C. In TT samples cured at 1,000°C (Figure 12H), the majority of pores have a diameter below 500 µm. Due to the higher addition of foaming agent, the percentage of pores between 100 and 500 µm is higher only for TT1, similar to in samples cured at 70°C. Sample TT3 cured at 1,000°C has a large percentage of pores in the pore diameter range below 100 µm (70%).

When compared to other RI samples cured at 1,000°C, sample RI1 has the highest volume of pores in the pore diameter range 100–850 µm (approximately 80%; Figure 12C), which is slightly less than in the samples cured at 70°C. In the RI2 samples almost 60% of the volume of pores lies in the 500–850 µm pore diameter range, while 25% of pores have a diameter larger than 850 µm. In the RI3 sample the largest percentage of the volume of pores is above a diameter of 500 µm (almost 92%), slightly less than in the samples cured at 70°C. At least 75% of the volume of pores in samples TS1 cured at 1,000°C have a pore diameter ranging between 100 and 500 µm (Figure 12F), which is less than in the TS1 sample cured at 70°C. Sample TS2 cured at 1,000°C has 70% in the same pore diameter range and sample TS3 almost 88% volume in the pore diameter range from 100–850 µm. Sample TT1 cured at 1,000°C has more than 60% of the volume of pores in the pore diameter range from 100 to 500 µm (Figure 12I), less than in the same sample cured at 70°C. A high volume of pores in this pore diameter range is also seen in the TT2 sample, reaching 70% (less than the sample cured at 70°C) while sample TT3 has poor porosity development.

The results for RI samples thermally treated at 1,000°C show that the majority of pores were close to spherical in shape, as their sphericity values are above 0.50 (between 0.55 and 0.95) with no huge difference between mixtures (Figure 8D). The same trend is evident as in RI samples thermally treated at 70°C. The same trend also occurs in TS samples thermally treated at 1,000°C, where TS1 samples have a sphericity smaller than 0.6, while more spherical pores occur in TS2 and TS3 (up to 0.8; Figure 8E). The sphericity of the pores in TT samples thermally treated at 1,000°C is relatively low, similar to as in TT samples thermally treated at 70°C, confirming their irregular shape, especially for sample TT3, which has the lowest sphericity due to poor pore shape (Figure 8F). The results for RI samples thermally treated at 1,000°C also show that sphericity is equal between mixtures, whereas it can be seen that sphericity slightly changes according to the mixture in TS samples (mixtures with higher addition of foaming agent have higher sphericity values), and in TT samples huge changes are observed between the mixtures, similar to in samples thermally treated at 70°C.

## Phase Changes by X-Ray Powder Diffraction

The XRD patterns of the studied ashes and AA foams cured at 1,000°C are shown in Figure 9. The results show multiple peaks specific to quartz and hematite, and in some cases mullite in the RI samples, and gehlenite in TS and TT samples. The aforementioned crystalline phases detected in the initial material remain apparently unaltered with heating, as it was not expected that the mullite would melt or dissolve during firing to 1,000°C (given their melting point of approximately 1,830°C). It is believed that the presence of free sodium in the matrix causes the dissolution/melting of the mullite into the geopolymer at
<figure><img src="image_469.png" /><figcaption>2 Theta (°)</figcaption></figure>

<figure><img src="image_470.png" /><figcaption>2 Theta (°)</figcaption></figure>

<figure><img src="image_471.png" /><figcaption>2 Theta (°)</figcaption></figure>

**FIGURE 9** | Phase composition of the analyzed alkali-activated foams cured at 70 and 1,000°C (RI, TS, and TT fly ash used, 1, 2, and 3 represent amount of foaming and stabilizing agents used), as determined by means of X-ray powder diffraction.

lower temperature, as alkali's are known to be excellent sintering agents (Rickard et al., 2012). Hematite, however, exhibited an increased peak intensity, especially in the TT samples (due to their higher concentration of iron). Rickard et al. (2012) reported that hematite is likely to have crystallized from the amorphous iron in the fly ash.

When the temperature is increased to 1,000°C, the diffraction peaks of the aforementioned phases become weak, and new phases start to crystalize into pyroxene, nepheline, tridymite, sodalite, brownmillerite, and gehlenite as response to the thermal treatment. The reduction of quartz peak intensity is likely due to a partial phase change to tridymite. In the presence of alkali oxides, quartz converted to tridymite, depending on the temperature. It has been reported that tridymite starts to form above 872°C with Na₂O and above 883°C with K₂O (Holmquist, 1961). These high temperature phases, tridymite (SiO₂) and nepheline (NaAlSiO₄), are reported to improve the thermal resistance of the geopolymers due to their high melting

points (tridymite 1,670°C and nepheline 1,257°C) (Rickard et al., 2012). In general, if silicon dioxide is available at such a high temperature, nepheline is no longer stable and feldspar will form. The amount of calcium available causes different types of feldspar to form (Dombrowski et al., 2007). It seems that the matrix decomposes, then the aluminosilicate species react with the released alkalis and calcium to form crystalline phases (nepheline and gehlenite, respectively) (Gyekenyesi et al., 2011). One reason to explain this could be that the mullite and quartz were needed to form the large amount of gehlenite (Ca₂(Al(AlSi)O₇)), (Dombrowski et al., 2007), as seen in the samples (especially TS and TT) treated at 1,000°C. In addition, at 1,000°C, a phase belonging to the pyroxene group (diopside) is additionally formed due to sufficient Al₂O₃ content, which has been reported to have a positive role in its crystallization (Qian et al., 2006). Yang et al. (2017) also reported that the Mg-silicate crystalline phases of diopside may work as compensation for the thermal shrinkage of the geopolymer gel phase, in order to
<figure><img src="image_472.png" /><figcaption>**FIGURE 10** | Representative SEM image of alkali-activated foams thermally treated at 1,000°C (RI, TS, and TT fly ash used, 1 and 3 represent the amount of foaming and stabilizing agent added). The bars on all SEM photos represent 500 µm (×35 magnification).</figcaption></figure>

improve the volume stability. The intensities of peaks and small amounts of already crystallized aluminosilicate detected corresponded to sodalite (Na₈(Al₆Si₆O₂₄)Cl₂), which has been observed to increase in fly ash samples activated by NaOH solution when exposed to temperatures of 1,000°C. Sodalite crystallizes in the cubic system as a mineral complex, formed by the reaction of sodium and chlorine with the main elements of the raw material (aluminum, silicon, and oxygen). The natural sodalite consists of an Al-O-Si network that encompasses Cl⁺ cations, but the one resulting from geopolymerization shows

inter-structural Na⁺ cations, similar to zeolites (Nergis et al., 2020).

Based on the properties of all three precursors, it can be seen that the fly ash which contains the highest amounts of SiO₂ and Al₂O₃, and the highest amorphous phase content, i.e., RI, is also the most reactive, both in terms of pozzolanic activity (Kramar and Ducman, 2018), and the alkali activation process. By adding foaming and stabilizing agents, the density can be tailored to one's needs, but the initial precursor should provide sufficient reactivity and consequently result in good mechanical properties. In our

<figure><img src="image_473.png" /><figcaption>**FIGURE 11** | Region of interest of the analyzed alkali-activated foams cured at 1,000°C (RI, TS, and TT fly ash used, 1, 2, and 3 represent the amount of foaming and stabilizing agent added) presenting grayscale image and the biggest main pores. Colors correspond to a range of size classes with respect to EqDiameter, purple (0.85–1 mm), green (1–1.6 mm), and yellow (more than 1.6 mm).</figcaption></figure>
<figure><img src="image_474.png" /><figcaption>**FIGURE 12** | Graphic representation of the pore size distribution of the analyzed alkali-activated foams cured at 1,000°C (RI, TS, and TT fly ash used, 1, 2, and 3 represent the amount of foaming and stabilizing agents added; (A)–(C) RI fly ash, (D)–(F) TS fly ash, (G)–(I) TT fly ash), determined by means of X-ray micro-tomography. The results show the number, percentage, and volume of pores with regard to pore size diameter (EqDiameter).</figcaption></figure>

study the RI fly ash exhibited the lowest density and when densities were comparable to those obtained from the other two fly ashes the compressive strength was significantly higher, which can be ascribed to the higher reactivity. Contrary to the expectation that BET specific surface would play an important role in the alkali activation process, by providing a surface available for reaction, in our case the BET specific surface area was not recognized as an important parameter.

## CONCLUSIONS

Three types of fly ash were tested as precursors for the preparation of AA foams. SPM was used as the foaming agent and SDS as the stabilizing agent. Foam properties were determined after curing at 70°C and after exposure to 1,000°C. In all three systems the

foaming process was successfully implemented, resulting in foams of density down to 0.3 g/cm³. Compressive strength is highly dependent on density, and for the densities from 1.2 down to 0.3 g/cm³, compressive strength ranged from 1 to 6 MPa. Comparing all three fly ashes the favorable results in terms of performance (density vs. mechanical properties) were achieved from the fly ash with the highest SiO₂ and Al₂O₃ content, and the highest amorphous phase content, i.e., RI fly ash. Also compressive strength is influenced by the type of fly ash; the sample from RI fly ash exhibited the highest compressive strength at a comparable (or even lower) density than the other two fly ashes. With the high amount of SiO₂ and Al₂O₃, and high amorphous phase content, the alkali activation reaction is completed to a greater extent, resulting in better mechanical properties. In the case of the fly ash with the highest SiO₂ and Al₂O₃ content, i.e., RI fly ash, the density remained virtually the
same after firing to 1,000°C, but the compressive strength increased on average by 50%. In the case of the other two types of fly ash the density increased slightly after firing, but significant shrinkage occurred. Density did not increase, as would be expected due to the high shrinkage. Even though density did not increase, compressive strength increased hugely, by as much as 800% in the case of the TT fly ash, which contains the lowest amount of SiO₂ and Al₂O₃.

XRD analysis also confirmed that the lime contained in the TS and TT fly ash disappeared after firing to 1,000°C, which could have contributed to the loss in mass and consequent reduction in density (in spite of significant shrinkage on firing). The crystallization process also took place during firing, resulting in newly formed crystal phases, including nepheline, sodalite, tridymite, and gehlenite, which could contribute to the increased compressive strength. SEM and microtomography analysis showed differences in the microstructure of foams, whereby pores are more uniform and spherical in foams based on RI fly ash compared to the other two fly ashes, which may also influence properties of thermal insulation and compressive strength. Improvements in the performance of foams in terms of density and mechanical properties, especially after exposure to elevated temperature when their structure remains stable after

firing up to 1,000°C, could expand the potential use of inorganic AAM foams into areas where resistance to elevated temperature is required.

## DATA AVAILABILITY STATEMENT

All datasets presented in this study are included in the article.

## AUTHOR CONTRIBUTIONS

Conceptualization, VD; data curation, VD and LK; formal analysis, VD and LK; funding cquisition, VD and LK; methodology, VD, LK; supervision, VD; writing—original draft, LK; writing—review and editing, VD and LK.

## ACKNOWLEDGMENTS

We acknowledge financial support from the Slovenian Research Agency through the project No. J2-9197 (VD) and post-doctoral project No. Z2-1861 (LK).

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**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2020 Korat and Ducman. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
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# Alkali-Activated Binders From Waste Incinerator Bottom Ashes and Metakaolin Reinforced by Recycled Carbon Fiber Composites

Stefania Manzi¹*, Isabella Lancellotti², Giulia Masi¹ and Andrea Saccani¹

¹Department of Civil, Chemical, Environmental and Materials Engineering, University of Bologna, Bologna, Italy, ²Department of Engineering "Enzo Ferrari", University of Modena and Reggio Emilia, Modena, Italy

In view of creating low-impact materials for the building industry, the fostering of alkali-activated binder gains high importance. Metakaolin can successfully be activated with alkalis at room temperature, but the contemporary use of wastes to create mixed binders can further increase the environmental benefits. Bottom ashes obtained from the incineration of municipal solid wastes have been tentatively mixed in different amounts to develop matrix with acceptable mechanical properties, which still can be cured at room temperature. Moreover, scraps obtained from the production of epoxy/carbon fiber composites are employed as a reinforcing phase. No chemical or physical treatments have been used to modify the epoxy/carbon fiber wastes, apart from size reduction, thus minimizing the overall economic and energy impact of the process. The workability, physical and mechanical properties, microstructure, and porosity of the obtained materials are investigated. Up to a 50 weight percent of bottom ashes from municipal solid waste incineration can be mixed with metakaolin. Fibers still embedded in the epoxy matrix disclose a fair interaction with the matrix, thus managing to increase flexural strength, toughness, and dimensional stability without decreasing the compressive strength.

**Keywords:** building materials, carbon fibers composites wastes, bottom ashes, recycling, alkali-activated materials

## INTRODUCTION

Municipal waste incineration can be used as a waste handling solution. This process can recover energy, exploiting the heat of combustion of the organic fraction of the waste and allows the overall reduction of the volume of wastes, consequently decreasing landfilling depletion. As a drawback, the process itself generates solid wastes, defined as fly ash and bottom ash. Bottom ashes from municipal solid waste incineration have some remarkable characteristics. They contain negligible amounts of toxic elements, have high amorphous fraction, and are largely composed of silica and aluminium oxides. These characteristics have already attracted interest for their use in producing new glass compositions (Saccani et al., 2001) or building materials, especially their pozzolanic activity in traditional Portland cement composites (Saccani et al., 2005; Li et al., 2012; Kuo et al., 2013; Tang et al., 2016; Chen and Yang, 2017; Chen et al., 2019). In recent years, a new class of binders has been investigated, characterized by a lower carbon dioxide footprint than Portland cement, that is, alkali-activated materials (De Vargas et al., 2011; Sun and Vollpracht, 2019). There are a set of inorganic

**OPEN ACCESS**

**Edited by:**
Patricia Krawczak,
IMT Lille Douai, France

**Reviewed by:**
Nassim Sebaibi,
École Supérieure d'Ingénieurs des
Travaux de la Construction, France

Girts Bumanis,
Riga Technical University, Latvia

*Correspondence:
Stefania Manzi
stefania.manzi4@unibo.it

**Specialty section:**
This article was submitted to *Structural Materials*, a section of the journal *Frontiers in Materials*

**Received:** 15 July 2020
**Accepted:** 06 October 2020
**Published:** 12 November 2020

**Citation:**
Manzi S, Lancellotti I, Masi G and Saccani A (2020) Alkali-Activated Binders From Waste Incinerator Bottom Ashes and Metakaolin Reinforced by Recycled Carbon Fiber Composites. *Front. Mater.* 7:583400.
doi: 10.3389/fmats.2020.583400
phases showing amorphous or semicrystalline structure formed by a three-dimensional network of AlO₄ and SiO₄ tetrahedra. They are obtained from a chemical reaction between a binding component (the source of aluminosilicates) and an alkaline activator, usually a blend of caustic alkalis or alkaline salts. For their synthesis, raw materials such as metakaolin, calcined clays, or industrial wastes, such as fly ash or slags, have been used (Boca Santa et al., 2016; Borges et al., 2016; Logesh Kumar and Revathi, 2016; Boca Santa et al., 2017; Bai et al., 2019). Bottom ashes from municipal solid waste incineration can also have an application in these materials because they can be alkali activated when mixed with metakaolin (Lancellotti et al., 2013; Lancellotti et al., 2014; Lancellotti et al., 2015; Silva et al., 2017; Wongsa et al., 2017; Huang et al., 2018). They cannot be used alone as precursor because the amount of Al₂O₃ is too low and the higher amount of calcium can change the chemistry of the polymerized network that is formed from alkali activation with the creation of C-A-S-H gels (Borges et al., 2016; Zhu et al., 2018). The alkali activation of this class of materials can take place at different temperatures (Rovnaník, 2010). The most interesting results are obtained from the activation at room temperature (Aredes et al., 2015). Other wastes can be activated only at high temperatures (Sha et al., 2020). It is thus important to verify the mechanical properties of the mixed binders when cured at room temperature.

However, the main problems concerning alkali-activated materials are their complete brittleness and their low dimensional stability. Both defects can be mitigated by modifying the composites through fibers addition, as reported in the literature (Uddin and Shaikh, 2013; Vilaplana et al., 2016; Al-Mashhadani et al., 2018; Bhutta et al., 2018; Luna-Galiano et al., 2018; Akono et al., 2019; Farooq et al., 2019). Waste carbon fibers (CF), embedded in a thermosetting polymer matrix, virtually have high mechanical properties. Either they can be derived from post-user dismissed items or they can be the pre-preg offcuts derived from the manufacturing process. Bare fibers can be recovered, eliminating the thermosetting matrix by means of high thermal treatments or selective chemical dissolution. Both the processes imply energy consumption and waste production. The direct use of scraps has been proposed in different matrixes, such as Portland cement, alkali-activated fly ashes, and pure metakaolin. Positive results have been obtained in terms of increased flexural strength and toughness, implying that the presence of a polymeric interphase (epoxy resin) between the fiber and the ceramic matrix still enables a positive interaction between fiber and matrix (Saccani et al., 2019). It seems thus interesting to evaluate the effect of CF/epoxy scraps also in the alkali-activated systems where metakaolin is partially substituted by bottom ashes from municipal solid waste incineration. The mix design has been investigated focusing on the workability of fresh mixtures, microstructure (i.e., fibers dispersion and porosity), flexural strength, and toughness. A room temperature curing has been studied, being the lowest energy-consuming treatment. In case these materials should prove to have acceptable mechanical properties, which is the main scope of the research, they could provide potentially high environmental benefit because they could enable the simultaneous recycling of

<figure><img src="image_476.png" /><figcaption>**FIGURE 1** | (A) Full view of composite pre-pregs scraps. (B) Micrograph of the shredded multi layers.</figcaption></figure>

<table><caption>**TABLE 1** | Composition of the raw materials (wt%).</caption><thead><tr><th>Material</th><th>SiO<sub>2</sub></th><th>Al<sub>2</sub>O<sub>3</sub></th><th>Na<sub>2</sub>O</th><th>K<sub>2</sub>O</th><th>CaO</th><th>MgO</th><th>Fe<sub>2</sub>O<sub>3</sub></th><th>TiO<sub>2</sub></th><th>MnO</th><th>LOI</th></tr></thead><tbody><tr><td>Metakaolin</td><td>53.5</td><td>41.9</td><td>0.0</td><td>2.6</td><td>0.0</td><td>0.0</td><td>1.4</td><td>0.6</td><td>0.0</td><td>0.0</td></tr><tr><td>Bottom ash</td><td>49.5</td><td>9.4</td><td>5.7</td><td>1.4</td><td>17.5</td><td>2.9</td><td>4.9</td><td>0.9</td><td>0.2</td><td>4.9</td></tr></tbody></table>

different wastes, reducing the amount of metakaolin and virgin carbon fibers to be used. The first issue eliminates the thermal treatments of aluminosilicate precursor that are necessary, thus avoiding CO₂ emissions, and the second one would prevent scrap damping and avoid virgin fiber production. Indeed, the complete life cycle assessment of the derived material should be carried out, a process that is however beyond the aims of this study that merely faces the technical problems related to the material performance upon modification.

## EXPERIMENTAL

### Composites Waste Scraps

Pre-preg offcuts derived from the production of CF/epoxy (hereafter defined as W_CF/EP) composites have been used. The as-received scraps had an almost rectangular shape with side dimension ranging from 10 to 40 cm (Figure 1A shows an example). They were made of multiple CF/epoxy impregnated layers with an average thickness ranging from 150 to 450 µm.
<table><caption>**TABLE 2** | Formulation of composites.</caption><thead><tr><th>Sample</th><th>Metakaolin (g)</th><th>BA (g)</th><th>NaOH 8M (ml)</th><th>Na<sub>2</sub>SiO<sub>3</sub> (ml)</th><th>W_CF/EP (g)</th><th>w/b ratio</th></tr></thead><tbody><tr><td>BA25-W_CF/EP0</td><td>75</td><td>25</td><td>25</td><td>60</td><td>0</td><td>0.36</td></tr><tr><td>BA25-W_CF/EP3</td><td>75</td><td>25</td><td>25</td><td>60</td><td>3</td><td>0.36</td></tr><tr><td>BA25-W_CF/EP6</td><td>75</td><td>25</td><td>25</td><td>60</td><td>6</td><td>0.36</td></tr><tr><td>BA25-W_CF/EP9</td><td>75</td><td>25</td><td>25</td><td>60</td><td>9</td><td>0.36</td></tr><tr><td>BA25-W_CF/EP12</td><td>75</td><td>25</td><td>25</td><td>60</td><td>12</td><td>0.36</td></tr><tr><td>BA50-W_CF/EP0</td><td>50</td><td>50</td><td>16</td><td>40</td><td>0</td><td>0.28</td></tr><tr><td>BA50-W_CF/EP3</td><td>50</td><td>50</td><td>16</td><td>40</td><td>3</td><td>0.28</td></tr><tr><td>BA50-W_CF/EP6</td><td>50</td><td>50</td><td>16</td><td>40</td><td>6</td><td>0.28</td></tr><tr><td>BA50-W_CF/EP9</td><td>50</td><td>50</td><td>16</td><td>40</td><td>9</td><td>0.28</td></tr><tr><td>BA50-W_CF/EP12</td><td>50</td><td>50</td><td>16</td><td>40</td><td>12</td><td>0.28</td></tr></tbody></table>

BA, bottom ash; CF, waste carbon fibers; W_CF/EP, pre-preg offcuts derived from the production of CF/epoxy; w/b ratio: water/binder ratio.

<figure><img src="image_477.png" /><figcaption>**FIGURE 2** | Relative workability of the investigated composites.</figcaption></figure>

<figure><img src="image_478.png" /><figcaption>**FIGURE 3** | Cumulative intruded volume in the plain matrices.</figcaption></figure>

The original multiple layer sheets were treated for 1 h at 130°C to complete the cross-linking reactions and then cut to obtain fibers bundles with the length of 13 ± 2 mm and lateral size ranging from 5 to 10 mm (Figure 1B). The amount of resin that surrounds the carbon fibers was evaluated by weight loss through the use of a TGA instrument (Thermogravimetric Analysis, Q50 Model TA Instruments) and was estimated as a 30 ± 8 wt%.

## Binders

Metakaolin with a dimension lower than 75 µm was used as binder. Bottom ashes from municipal waste incineration, hereafter simply defined as BA, once collected by the company managing the disposal of ashes from different incinerator plants are submitted to beneficiation. The process involves ageing, grinding, and Fe and Al separation by means of magnetic and eddy current systems. After these treatments, BA are no more classified as waste according to the European Waste Code, but as secondary raw materials (end of waste). Final product was ball-milled and sieved through a 100 µm sieve to favor their homogeneous mixing with metakaolin and subsequent reactions. In details, metakaolin powders show the following parameters: D₁₀ = 1.40, D₅₀ = 11.08, and D₉₀ = 50.98. For BA the values are as follows: D₁₀ = 1.81, D₅₀ = 21.53, and D₉₀ = 61.71. Data have been derived by laser particle sizer Fritsch, Analysette 22. Table 1 reports the chemical composition of both metakaolin and BA as derived by X-ray fluorescence (Philips PW 2004). Other detected elements, such as sulphur, phosphor, and chlorine, are not reported. Loss on ignition was determined after a 2 h treatment at 1,100°C. X-ray analysis was performed in order to evaluate the amorphous or crystalline nature of BA and to identify the crystalline phases present. From the analysis (not reported for briefness sake), the typical amorphous/
<table><caption>**TABLE 3** | Classification of porosity in the investigated samples according to IUPAC.</caption><thead><tr><th>Designation</th><th>Porosity range (nm)</th><th>BA 25% (mm<sup>3</sup>Hg/g) (%)</th><th>BA 50% (mm<sup>3</sup>Hg/g) (%)</th></tr></thead><tbody><tr><td>Micropores</td><td>&lt;1.25</td><td>n.d.</td><td>n.d.</td></tr><tr><td>Mesopores</td><td>1.25-25</td><td>156.0 (68.4)</td><td>67.0 (42.4)</td></tr><tr><td>Macropores</td><td>25-5,000</td><td>71.9 (31.5)</td><td>88.0 (55.7)</td></tr><tr><td>Directly accessible large pores</td><td>5,000-50,000</td><td>0.1 (0.0)</td><td>3.0 (1.9)</td></tr><tr><td>Total porosity</td><td>—</td><td>228.0 (100)</td><td>158.0 (100)</td></tr></tbody></table>

Pore size classification (percentage of porosity over the total specific volume of Hg is reported in brackets).

<figure><img src="image_479.png" /><figcaption>**FIGURE 4** | Water absorption of all the investigated samples.</figcaption></figure>

crystalline nature of the material is determined. Among the crystalline fraction, the main detected phase is α-quartz (α-SiO₂, JCPDF file 33-1161) followed by calcite (CaCO₃ JCPDF file 5-586), and aluminosilicates as albite (NaAlSi₃O₈ JCPDF file 10-393) and ghelenite (Ca₂Al(Al,Si)O₇ JCPDF file 35-755), thus reflecting the chemical analysis of a typical ash rich in calcium and sodium.

## Activators

Sodium hydroxide reagent grade and sodium silicate solution, a viscous liquid produced for the cement industry with a water content of 57 wt% and a SiO₂/Na₂O ratio of 3, were used as alkaline activators.

## Mixing Procedures and Compositions

Sodium hydroxide 8M solution and sodium silicate have been first mixed according to the quantities determined in previous experiments (Lancellotti et al., 2013; Lancellotti et al., 2014). The quantities added aimed primarily to obtain equal workability. Moreover, due to the higher content of Na in BA, less amount of sodium hydroxide and silicate are added. On account of the higher ratio of SiO₂/Al₂O₃ in BA, less silicate is added (Table 1). Afterward, metakaolin mixed with the different quantity of BA (25 or 50 wt%) was added. The workability was determined through the minislump test proposed by EN-ETC. No additional water has been added to the fiber-modified materials and consequently workability decreased as the amount of W\_CF/EP increased. However, up

to the highest investigated quantity, it was still possible to effectively fill the moulds without creating macro-voids. W\_CF/EP amounts were 3, 6, 9, and 12 wt% on the powders. Table 2 reports the mix design of the investigated materials. The dimension of the prepared samples during casting procedures was different on the basis of further characterizations. Prisms of 100 mm × 20 mm × 20 mm size were cast (nine for each composite composition) and cured for 28 days at 25 ± 1°C and 98 ± 1% relative humidity. Prisms of 250 mm × 25 mm × 25 mm were used to investigate the free-drying shrinkage vs time when samples were cured at 25°C and 35 ± 5% relative humidity.

## Tests and Procedures

### Water Absorption and Porosity

Water absorption has been evaluated according to the UNI 7699 Standard (UNI 7699, 2005) on three samples (100 mm × 20 mm × 20 mm) cured for 28 days. The cumulative open porosity and the size distribution of the porosities in samples without fibers (i.e., the plain matrixes formulated with 25 and 50 wt% of BA but without the fibers) have been evaluated by mercury intrusion porosimetry (MIP) equipped with a macropore unit (Thermo Scientific, Pascal 140 and 240). Measurements were conducted at 28 days of curing. MIP measurements performed on composite samples were not reliable or reproducible because of the small dimension of the tested specimens allowed and the unavoidable damage induced by specimen size reduction. Accordingly, they will not be reported.
<figure><img src="image_480.png" /><figcaption>**FIGURE 5** | Optical micrograph of BA25-W_CF/EP samples cross sections at 3 (A), 6 (B), 9 (C), and 12 (D) wt. amount of fibers at 20x.</figcaption></figure>

<figure><img src="image_481.png" /><figcaption>**FIGURE 6** | Flexural strength of the investigated materials.</figcaption></figure>

## Mechanical Tests

Mechanical tests (three-point bending) were performed at room temperature with a relative humidity of 50 ± 10% by means of 100 kN Amsler Wolpert equipment, with a 5 mm/min displacement rate. Flexural strength values are reported as the average of six measurements. Compressive strength measurement was performed on both the prisms derived from the previous flexural test (12 measurements).

## SEM Analysis

Morphological investigations were carried out on fractured surfaces of mortar samples sputtered by graphite by means of

<figure><img src="image_482.png" /><figcaption>**FIGURE 7** | Force vs deflection plot of specimens containing 50% of BA at different amounts of W_CF/EP.</figcaption></figure>

electron scanning microscopy (ESEM, Quanta-200, FEI Co.), exploiting the signal collected by secondary electrons detectors. An accelerating voltage of 20 kV was applied.

## Dimensional Stability

Dimensional stability has been evaluated according to the ASTM C1012/C1012M standards (ASTM C1012/C1012M, 2015) on three prismatic specimens of 250 mm × 25 mm × 25 mm cured at 25°C and 35 ± 5% relative humidity.
## RESULTS

Figure 2 shows the relative workability of samples containing the different amounts of recycled composites, as well as the different amounts of BA. Standard deviations are not included for the sake of clearness because they almost coincide with the dimension of the symbols. While the effect of the different amounts of BA is negligible as desired, W_CF/EP addition progressively decreases this property. However, as underlined before, all mixtures could be easily cast without the formation of macrovoids. The decrease is partially related to the amount of water requested to wet the reinforcement surfaces. Although the composites scraps are no more flexible after the full-cross-

linking reaction, their small dimension prevents buckling phenomena.

Figure 3 shows the results of MIP analysis: the total intruded volume in the two different plain matrices is reported vs the pore dimension. The sample containing the lower amount of BA (25%) shows the highest overall porosity (228 mm³/g of intruded volume), with almost half of the pores having dimensions lower than 0.01 µm. BA 50% matrix shows a lower overall porosity (158 mm³/g of intruded volume), but the dimension of the pores is larger than that in the previous sample, mainly falling in the range from 0.4 to 0.01 µm. To get a detailed analysis of the results, Table 3 reports the amount of pores divided in different ranges proposed by the IUPAC classification. The main differences are related to the mesopores and macropores that suggest similar observations as those previously proposed. In the BA 25% sample, no large pores (>5,000 nm) are detected, and the amount is quite limited in the BA 50% sample.

This can be explained by assuming that, in the applied curing conditions, the reactions involving the BA fraction are slower than those involving metakaolin, thus producing a lower amount of binding products or polymerized network only partially filling the larger voids, as found in other researches (Guo et al., 2017; Sha et al., 2020). Moreover, the presence of C-S-A-H product can change the porosity of the products. Consequently, although the w/b ratio in the BA 50% sample is lower, the capillary porosity is increased.

Figure 4 reports the water absorption of all the samples after 28 days of curing. The higher amount of BA slightly decreases the amount of absorbed water. Moreover, as the amount of recycled composite increases, values tend to decrease slightly at both the BA amounts. The percentage of absorbed water is usually linked to the overall volume of open porosity in the material. This result is also in accordance with the previous observation on the absence of large voids caused by the workability reduction. Moreover, as the amount of absorbed water progressively decreases, there is absence of a discontinuity

<figure><img src="image_483.png" /><figcaption>**FIGURE 8** | Compressive strength of the investigated materials vs amount of recycled composite.</figcaption></figure>

<figure><img src="image_484.png" /><figcaption>**FIGURE 9** | Microstructure of (A) BA25-W_CF/EP0 and (B) BA50-W_CF/EP0 (fracture surfaces).</figcaption></figure>
<figure><img src="image_485.png" /><figcaption>**FIGURE 10** | Fracture surface of (A) BA25-W_CF/EP6 and (B) BA50-W_CF/EP6.</figcaption></figure>

<figure><img src="image_486.png" /><figcaption>**FIGURE 11** | Dimensional changes for samples cured at 35 ± 5% relative humidity.</figcaption></figure>

between the waste and the matrix, suggesting an efficient adhesion between the different phases and the creation of a continuous interface transition zone between the cured epoxy and the ceramic matrix that can promote the increase in mechanical properties. The positive interaction is related to the presence of dipoles formed by C-O bonds in the resin interacting with ceramic matrix. The slight decrease in BA 50% compared with BA 25% can be explained considering the different total porosity of the two matrixes found by MIP analysis as observed in **Figure 3**. The same observation can explain the progressive reduction in the absorbed water as the amount of W_CF/EP increases. Indeed, a partial substitution of the porous phase, the matrix, with a non-porous fraction, the organic composite, takes place, decreasing the overall fraction of porosities.

The severed surfaces obtained by diamond sawing of the different samples (Figures 5A–D) again underline the absence of large voids (>1 mm) derived from insufficient workability, as was expected by visual inspection during casting operations. No differences at this magnification were found

between the different amounts of BA and consequently only 25 wt% BA loading is reported. No clusters of fibers were also visible from the Figures, implying that an efficient dispersion is obtained.

**Figure 6** shows the flexural strength of the composites at 28 days of curing. As it can be observed, the increase in the amount of BA, from 25 to 50 wt%, decreases the mechanical properties of the composites. This is probably related to the presence, in the 50% samples, of a higher amount of porosities in the range from 0.02 to 0.4 µm with respect to the 25% BA composition as observed from the MIP measurements. This difference is possibly derived, as explained before, from the different reaction products generated (i.e., from the presence of C-A-S-H products) and also from the different kinetics of the binding reactions involving metakaolin and BA. However, the addition of W_CF/EP almost linearly increases the flexural strength of all the samples, confirming the positive effect of the fibers.

**Figure 7** shows the plots reporting the force vs the deflection of all samples containing 50% of BA. For sake of brevity, only this set of specimens is reported. Along with the increase in flexural strength already discussed, a progressive increase in toughness takes place. In fiber-modified samples instead of one crack leading to the almost instantaneous overall fracture, a web of small cracks is formed leading to a postponed failure. Consequently, as the fully brittle behavior of the unmodified matrix is changed, higher amounts of energy can be absorbed as the amount of recycled composite increases. The mechanism supporting this effect is related to the progressive pull-out of the fibers from the matrix. The toughening effect of the fibers is reflected by the increase in the area beneath the plot.

**Figure 8** reports the values of compressive strength of the investigated materials. This property is almost unaffected by the waste addition, both W_CF/EP and BA. Only at the highest amount of fibers, a slight reduction is found in samples, whereas at the lowest amounts, compressive strength slightly increases. Larger increases in compressive strength were found in other
researches, although with different conditions (Uddin and Shaikh, 2013; Bhutta et al., 2018).

The SEM analysis of mortars without W_CF/EP at different BA amounts is reported in **Figure 9**. Both the matrixes appear as compact without porosities. In the 50% BA matrix, some unreacted BA particles still showing sharp edges are visible. This confirms that at room temperature, the reactions leading to the geopolymerization of at least one fraction of BA are rather slow as previously observed (Uddin and Shaikh, 2013; Bhutta et al., 2018). **Figure 10** reports the morphology of the fractured surfaces derived from the flexural tests of samples modified with W_CF/EP. Some matrix particles adhere to the fibers, and no porosities are present at the interphase between matrix and wastes. This implies a fair interaction of the filler with the matrix, supporting the previous observations related to the amounts of absorbed water.

Indeed, it should also be underlined that no remarkable qualitative differences can be observed on changing the amount of BA in the matrix. Eventually, **Figure 11** reports the dimensional changes at 14 days for samples cured at 35 ± 5% relative humidity. A slight increase in the dimensional shrinkage of sample BA25-W_CF/EP can be observed when compared with BA50-W_CF/EP, probably on account of the overall higher porosity. Samples containing 3% of W_CF/EP show almost the same behavior as the unmodified matrixes, but further additions progressively increase the dimensional stability of the composites, an effect that has been found in conventional Portland cement materials by adding recycled carbon fibers (Wang et al., 2019).

It is important to underline from the results so far obtained that wastes derived from composite manufacturing can be recycled in the production of building materials without previous chemical or thermal treatments. The scraps can improve the mechanical behavior of metakaolin materials geopolymerized at low temperature. In this way, the use of low-energy consuming, low-impact products, where an economical value-added benefit is also achieved, can be successfully promoted.

## CONCLUSIONS

The results of this research can thus be summarized as follows:

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* Metakaolin can be mixed with BA up to an amount of 50 wt %. The binder can still be activated at room temperature still showing even at the highest amount of BA, acceptable mechanical strengths.

* Untreated wastes derived from the production of carbon fiber epoxy composites have been homogeneously dispersed to formulate composites up to 12 weight %. Although the workability of the fresh mixtures decreases as the amount of waste increases, it is possible to cast mortars without high-dimension porosities. Microstructural analysis discloses a fair interaction between the fibers still surrounded by the epoxy coating and the matrix.

* CF/epoxy wastes increase the flexural strength of all the composites and drive their fracture from a brittle behaviour to a semiductile one. The compressive strength is practically unaffected by carbon fibers addition. Moreover, the addition of the fibrous wastes increases the dimensional stability of the composites.

## DATA AVAILABILITY STATEMENT

All datasets presented in this study are included in the article.

## AUTHOR CONTRIBUTIONS

SM: conceptualization, methodology, validation, investigation, resources, data curation, visualization, writing original draft preparation, and editing. IL: methodology, validation, and resources. GM: investigation, data curation, and visualization. AS: conceptualization, methodology, validation, investigation, resources, data curation, visualization, writing original draft preparation and editing, and supervision.

## ACKNOWLEDGMENTS

The authors acknowledge the support of Lorenzo Lipparini (REGLASS) for providing CF/EP composites.

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**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2020 Manzi, Lancellotti, Masi and Saccani. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
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