# 3D crystalline phase and pore structure evolution upon CO2 exposure in sodium sulfate-activated cement pastes

Zengliang Yue, Zixian Su, Partha P. Paul, Alastair T.M. Marsh, Alice Macente, Marco Di Michiel, John L. Provis, Philip J. Withers, Susan A. Bernal

► To cite this version:

Zengliang Yue, Zixian Su, Partha P. Paul, Alastair T.M. Marsh, Alice Macente, et al., 3D crystalline phase and pore structure evolution upon CO2 exposure in sodium sulfate-activated cement pastes. Cement and Concrete Research, 2025, 187, pp.107716. 10.1016/j.cemconres.2024.107716. hal-05099723

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# 3D crystalline phase and pore structure evolution upon CO₂ exposure in sodium sulfate-activated cement pastes

Zengliang Yue<sup>a</sup>, Zixian Su<sup>b</sup>, Partha P. Paul<sup>b,c</sup>, Alastair T.M. Marsh<sup>a</sup>, Alice Macente<sup>a</sup>, Marco Di Michiel<sup>c</sup>, John L. Provis<sup>d</sup>, Philip J. Withers<sup>b</sup>, Susan A. Bernal<sup>a,e,*</sup>

## ARTICLE INFO

**Keywords:**

X-ray computed tomography

X-ray diffraction computed tomography

Alkali-activated cements

Sodium sulfate

Carbon capture

Pore network

## ABSTRACT

The effects of CO₂ exposure on sodium sulfate-activated blast furnace slag cement paste have been characterised by X-ray (attenuation) computed tomography revealing changes in micron-scale pore structure, and X-ray diffraction computed tomography (XRD-CT) elucidating changes in the spatial distribution of crystalline and semi-crystalline phases. Accelerated carbonation reduced ettringite volumes and induced formation of hydro-talcite, demonstrating the critical role of Mg-Al-SO₄-layered double hydroxide phases in the CO₂ uptake of these cements. These changes yield a refinement of small pores and increase the overall porosity, reaching values comparable to those of blended Portland cements. Formation factor values were determined considering the pore solution electrical resistivity, calculated from thermodynamic modelling, and the porosity. A correlation between simulated tortuosity and porosity is proposed to estimate the diffusion tortuosity and formation factor of sodium sulfate-activated slag pastes. This approach represents a significant step forward for assessing carbonation resistance and CO₂ uptake capacity of cementitious pastes.

## 1. Introduction

Alkali-activated slag cements (AASCs) are clinker-free binders manufactured through the chemical reaction of ground granulated blast furnace slag (GGBFS) with alkaline substances called activators, typically sodium hydroxide or silicates [1]. AASCs are among the most promising low-carbon cements that could be used in place of Portland cement in concrete production, and thus can contribute to achieving net zero targets by the construction industry [2]. These cements also develop unique properties making them attractive for niche but important applications for tackling climate change and other environmental issues, such as nuclear waste immobilisation [3] and well-cementing for carbon storage [4,5].

Sodium sulfate (Na₂SO₄) is gaining the attention of industry and academia as a potential alternative to sodium hydroxide or sodium silicate in activating AASCs. This is mainly because cements produced with this activator develop pore solution pH values comparable to those of conventional Portland cement [6], reducing concerns regarding

corrosion of steel reinforcement. This activator can be mined or sourced as a by-product of hydrogen chloride production and silica pigments at potentially competitive costs and emissions footprints [7–9], further increasing the sustainability credentials of these cements. However, to date, sodium sulfate-activated slag cements have received limited attention, with research, restricted mainly to identifying the phase assemblage (type and amount of reaction products forming) in these materials using a single source of slag [10–12]. Its durability performance, particularly when exposed to CO₂, remains largely unknown.

Carbonation is one of the most common threats influencing the durability of cement-based structures. The ingress of CO₂ present in the air decreases the pH of the cement pore solution via the formation of carbonic acid, which then reacts with the cement reaction products, inducing microstructural alterations that can lead to deterioration of the strength of these materials [13,14]. The reduced pH also increases the risk of corrosion of mild steel reinforcement. It is well known that the carbonation mechanism for alkali-activated slag cements is different from that of Portland cements due to the differences in the types of

<footnote><sup>a</sup> School of Civil Engineering, University of Leeds, Leeds, United Kingdom</footnote>

<footnote><sup>b</sup> Henry Royce Institute, Department of Materials, University of Manchester, Manchester M13 9PL, United Kingdom</footnote>

<footnote><sup>c</sup> ESRF-The European Synchrotron, 71 Avenue des Martyrs, 38000 Grenoble, France</footnote>

<footnote><sup>d</sup> Paul Scherrer Institute, Laboratory for Waste Management, Forschungsstrasse 111, 5232 Villigen PSI, Switzerland</footnote>

<footnote><sup>e</sup> Institute for Building Materials (IfB), ETH Zürich, CH-8093 Zürich, Switzerland</footnote>

<footnote>* Corresponding author at: School of Civil Engineering, University of Leeds, Leeds, United Kingdom.
E-mail address: s.a.bernallopez@leeds.ac.uk (S.A. Bernal).</footnote>
reaction products that dominate each binder. The absence of portlandite in AASC means that carbonation causes the decalcification of the main strength-providing phase, an aluminium-substituted calcium silicate hydrate (C-A-S-H), favouring the formation of cross-linked aluminosilicate gels and CaCO₃ polymorphs [15,16]. This leads to changes in the pore structure and consequently in the transport properties (movement of fluids or ions) of the carbonated binders [16,17], facilitating the diffusion of gaseous CO₂ [18,19]. The complexity of the physical-chemical interaction of CO₂ in cementitious materials means that elucidating the change in transport properties upon carbonation to predict their carbonation performance is non-trivial.

Hardened cements are multi-phase porous materials whose pore size varies from mm to nm scale [20–22]. The long-term performance of the cementitious materials is determined by the ionic transport of species via advection and/or diffusion through the pore structure [23–25], which in turn affects the rate of degradation mechanisms such as carbonation [18,26], sulfate attack [27,28], or chloride-induced corrosion [29,30]. Assessing the durability of a novel AASC is important for its widespread uptake in different applications. This is strongly dependent on its microstructure features including the pore network [31].

Techniques currently employed to investigate porosity in cementitious systems include nitrogen adsorption/desorption (NAD) [32], mercury intrusion porosimetry (MIP) [33], ¹H NMR (nuclear magnetic resonance) relaxometry [34], small-angle X-ray/neutron scattering (SAX/SANS) [35] and micro-CT (micron-scale computed tomography) [36]. X-ray micro-CT is the reconstruction of a 3D image of the object on the basis of many projections (radiographs) collected as the specimen is rotated, usually through 180° or 360°; it generates grayscale images based on variations in the density and composition of the specimen [37]. This technique enables 3D visualisation of the pore structure, a capability that bulk measurement techniques like MIP, NAD, SAXS/SANS, and ¹H NMR relaxometry lack, as they provide only averaged measurements across the entire sample and require drying the sample prior to testing. In this way, the three-dimensional (3D) pore network can be characterised non-destructively. From this, the connectivity, tortuosity, and porosity can be quantified [31,38,39]. However, there is a trade-off between spatial resolution and specimen size.

Previously, microtomography (µCT) has been conducted on AASCs with a voxel size of around 1.5 µm [40] and in Portland cement as fine as 0.5 µm [41]. These voxel sizes are characteristic of µCT and mean that identification and quantification of the nanometre-size pores is not possible by µCT. Pixel size is calculated by dividing the field of view (FOV) by the matrix size, while spatial resolution refers to the system's ability to distinguish between closely spaced structures, typically assessed in terms of spatial or contrast resolution [42]. In addition to pixel size and matrix size, spatial resolution is heavily influenced by various technical factors of the CT system, such as detector technology [43], X-ray characteristics [44] and reconstruction algorithms [45]. Nevertheless, the pore size larger than 10 µm [46] can be observed, which are the most relevant types of pore in terms of ionic movement and hence the durability of cementitious materials [47]. In addition, nano-sized pores predominantly appear in the calcium silicate hydrate (C-S-H) or aluminium-substituted calcium silicate hydrate (C-A-S-H), which although responsible for the strength properties of the cement, might be less critical in defining the durability of these materials, depending on the main mechanism of attack [48]. Therefore, the resolution achievable by µCT is deemed sufficient to characterise the pore structure and provide meaningful insights into the transport properties of cementitious systems.

Carbonation causes changes in the phase assemblage and pore structure of cementitious materials, but there is very limited understanding of how phase assemblage alterations modify the pore structure, or vice versa. Although scanning electron microscopy coupled with energy-dispersive X-ray spectroscopy provides elemental distribution information and enables visualisation of the phase assemblage (including porosity), it is limited to two dimensions [49] and is

destructive, meaning that it is difficult to conduct repeat analysis of the same point(s) on a specimen. Phase contrast tomography (PCT) [50] has been used to provide more contrast between the constituents, while nanotomography has been used to achieve greater spatial resolutions [51,52].

There are, however, very few techniques able to collect 3D structural and mineralogy data simultaneously, and most of them are rarely applied to the study of cementitious materials. One such technique is X-ray diffraction computed tomography (XRD-CT) [53] which combines X-ray diffraction with computed tomography to deliver spatially resolved information on crystalline/semi-crystalline phases. There have been a few studies using XRD-CT to understand the phase changes that occur in cement and concrete [54,55]. Seminal studies have focused on evaluating the hydration of ordinary Portland cement [56,57] and blended Portland cements [58], mainly to identify the distribution of portlandite or ettringite and C-S-H gel in the material. However, to date, the technique has not been used to study alkali-activated cements, or for delineating the effects of carbonation in such systems.

Here, the effects of CO₂ exposure on the microstructure of sodium sulfate-activated slag cements produced with three types of blast furnace slag are investigated. In order to simultaneously elucidate the changes in phase assemblage and changes in pore structure resulting from carbonation, synchrotron XRD-CT imaging and higher resolution synchrotron µCT results are presented and discussed. This is the first study to visualise the reaction products forming in sodium sulfate-activated slag cements and to determine how they change as a consequence of accelerated carbonation. Changes in pore structure, diffusion tortuosity and formation factor of the assessed material are also examined by using three commercial European blast furnace slag precursors. More broadly, this study also demonstrates the potential of combining XRD-CT and µCT to study complex multi-component composite materials such as alkali-activated cements, providing unique and novel insight into their microstructural features as a function of mix design and chemical interaction with CO₂.

## 2. Experimental section

### 2.1. Material characterisation and sample preparation

Three commercial blast furnace slags were used as the main precursors for producing AASC. The oxide compositions of these slags were determined by X-ray fluorescence (XRF) using a fused bead preparation method in a Rigaku ZSX Primus II, Table 1. The three slags were selected with varying Mg/Al ratios of 0.8, 0.7, and 0.5 and are denoted as S1, S2,

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

Chemical oxide composition and physical properties of the anhydrous slags.</caption><thead><tr><th rowspan="2">Oxides (wt%)</th><th colspan="3">Precursors</th></tr><tr><th>Slag 1 (S1)</th><th>Slag 2 (S2)</th><th>Slag 3 (S3)</th></tr></thead><tbody><tr><td>CaO</td><td>42.7</td><td>40.3</td><td>42.2</td></tr><tr><td>SiO<sub>2</sub></td><td>35.8</td><td>35.6</td><td>34.4</td></tr><tr><td>Al<sub>2</sub>O<sub>3</sub></td><td>9.6</td><td>11.5</td><td>12.0</td></tr><tr><td>MgO</td><td>7.2</td><td>8.0</td><td>5.6</td></tr><tr><td>SO<sub>3</sub></td><td>2.0</td><td>2.4</td><td>2.9</td></tr><tr><td>Fe<sub>2</sub>O<sub>3</sub></td><td>0.7</td><td>0.3</td><td>0.5</td></tr><tr><td>TiO<sub>2</sub></td><td>0.7</td><td>0.6</td><td>1.1</td></tr><tr><td>Na<sub>2</sub>O</td><td>0.3</td><td>0.3</td><td>0.3</td></tr><tr><td>K<sub>2</sub>O</td><td>0.5</td><td>0.5</td><td>0.6</td></tr><tr><td>Others</td><td>0.5</td><td>0.5</td><td>0.4</td></tr><tr><td>Loss on ignition (LOI)<sup>a</sup></td><td>0.7</td><td>1.1</td><td>1.1</td></tr><tr><td>Physical properties</td><td></td><td></td><td></td></tr><tr><td>Particle size (µm) d<sub>10</sub>, d<sub>50</sub>, d<sub>90</sub></td><td>3.30/12.5/32.3</td><td>2.2/13.4/41.8</td><td>3.4/11.7/31.3</td></tr><tr><td>BET specific surface area (m<sup>2</sup>/g)</td><td>1.48 ± 0.17</td><td>1.38 ± 0.11</td><td>1.53 ± 0.18</td></tr></tbody></table>

<footnote>a Loss on ignition (LOI) determined by thermogravimetric analysis at 700 °C, under 60 mL/min of N₂ flow.</footnote>
and S3 respectively. The particle size distributions were determined using a Malvern Mastersizer 2000 laser diffractometer, employing iso-propanol as the dispersing medium. Nitrogen adsorption measurements were conducted at 77 K using a Micromeritics Tristar 3000 instrument to ascertain the Brunauer-Emmett-Teller (BET) surface area [59]. The three slag samples were stored in a desiccator until use.

The activator used in this study is an anhydrous sodium sulfate powder (Acros Organics, Na₂SO₄ > 99 %), which was dissolved in deionised water at 40 °C and then cooled down to room temperature (20 °C) prior to use. The Na₂SO₄-activated slags were prepared with an activator dose of 6 g Na₂SO₄ per 100 g slag, and a water-to-solids (slag + sodium sulfate) mass ratio of 0.40. To generate homogeneous mixtures, a high-shear mixer (Heidolph) was used to mix the pastes for 10 min. A mixing speed of 200 rpm was used for the first 3 min, increased to 400 rpm for the next 3 min, and further increased to 800 rpm for the final 3 min. Between speed changes, the mixer was stopped for 30 s to remove the paste adhered to the container wall, which was placed back into the middle of the container to ensure homogeneous mixing. The pastes were then cast in homemade moulds (inner diameter: 2 mm and height: 4 mm) and then sealed for curing at 20 °C for 28 days.

## 2.2. CO₂ exposure

Each specimen was mounted on a rotation stage and scanned by XRD-CT (described in Section 2.3) in the pristine state. Subsequently, the specimen was placed into a desiccator under vacuum. A 5 % CO₂/95 % N₂ (v/v) gas mix was then injected until atmospheric pressure was reached. One cycle of CO₂ exposure corresponded to 18 min of exposure. After completing the first cycle of CO₂ exposure (cycle 1), the mounted specimen was placed back on the stage. Similar steps were repeated for cycle 2 of carbonation (i.e., an additional 18 min of CO₂ exposure). The CO₂ concentration used is quite high for evaluating the carbonation of alkali-activated slag cements for predicting performance under environmental conditions [60], mainly because exposure beyond 1 % CO₂ induces changes in the types of sodium and calcium carbonates forming. However, this study aimed to validate the applicability of the XRD-CT technique to determine different microstructural features induced by accelerated carbonation of these materials, and therefore this CO₂ concentration was selected to induce severe and rapid changes in the paste within the experimental synchrotron beamtime constraints.

## 2.3. Synchrotron µCT and XRD-CT measurement

The µCT and XRD-CT experiments were conducted at the European Synchrotron Radiation Facility (ESRF) on Beamline ID15A [61]. A monochromatic 50 keV X-ray beam was used for the scan. After precise positioning and alignment, the specimen was rotated through 360° during data acquisition. A total of 7201 projections were captured with an exposure time of 0.04 s per projection, resulting in an overall scan time of 7201 projection × 0.04 s/projection = 288 s (approximately 5 min) per sample. A CMOS detector (PCO-edgeHS 5.5, 2560 × 2160 px) with 10× optical magnification was used to obtain the pixel size of 0.7 µm. The Nabu pipeline [62] was used to perform filter-back projections and reconstruct the volume of the sample.

For the XRD-CT, a monochromatic X-ray beam with an incident energy of 50 keV (wavelength λ = 0.2480 Å), with a Dectris Pilatus3 X CdTe 2 M area detector at a distance of 2 m from the sample, was used to collect diffractograms. The measured *d*-spacings ranged from 3.7 Å to 27.3 Å. A silver behenate standard was used for the detector geometry corrections. Compound refractive lenses were used to focus the beam to a ~10 × 10 µm² spot size on the sample. To obtain spatially resolved phase information, XRD patterns were collected at 300 different angular orientations, with an angular step size of 1.2° and a step size of 10 µm in both the horizontal and vertical directions. Each vertical stack, with a height of 10 µm, required approximately 15 min for the line scan, and a total of 31 line scans were performed for each sample. The 2D

diffractograms were azimuthally integrated using the PyFAI module [63]. Subsequently, the integrated diffractions were tomographically reconstructed with the filter back projections using in-house Matlab scripts, to obtain 3D spatial distributions of diffractograms. All diffractograms were converted to equivalent 2θ based on Cu-Kα radiation (λ = 1.5406 Å) for comparison with results collected from lab-based XRD. Further details on the data analysis pipeline are provided in Sections 1 and 2 of the Supporting Information (SI) file.

## 2.4. Data analysis

For µCT, an unsupervised machine learning approach, Trainable WeKa Segmentation (TWS), was used to segment the phases in the volume images (tomograms) [64]. Avizo 2022.1® (ThermoFisher Scientific) was used to characterise the pore network and for 3D rendering. A simulated formation factor analysis was conducted in Avizo XLab, by inputting the electrical resistivity of pore solution results collected from thermodynamic modelling using GEM-Selektor v.3 (https://gems.web.psi.ch) [65]. Diffusion tortuosity was calculated using the pyrox Python script [66]. The detailed workflow used in processing µCT, XRD-CT datasets, as well as the modelling approaches adopted to simulate the formation factor and pore solution resistivity, are reported in the Supporting Information (SI) file, Sections S2 and S3.

# 3. Results and discussion

## 3.1. Effect of secondary reaction products on carbonation performance

The reaction products of sodium sulfate slag cements are well-known, but their exact relationship with the evolving porosity under carbonation is still not fully understood. The XRD-CT measurements were limited to angles <26° Cu-Kα equivalent 2θ, where the main diffraction peaks of ettringite and Mg-Al layered double hydroxide (LDH) type phases are identifiable in alkali-activated cements [10]. These are the main crystalline reaction products forming in sodium sulfate-activated slag cements, and consequently, this section will mainly focus on analysing the effect of carbonation exposure on changes in their crystallinity.

The primary limitation preventing the use of QXRD (quantitative X-ray diffraction) in this study is the restricted diffraction range, which extends only up to 26° 2θ (Cu-Kα equivalent). During measurements, troubleshooting was necessary due to the presence of unreacted crystalline quartz and calcite. Notably, the quartz diffraction peak (~28° 2θ, Cu-Kα equivalent) is an order of magnitude higher than the detector's intensity threshold when using 50 keV X-rays, prompting the detector to shut down as a protective measure. To address this, the detector was re-engineered and partially masked, allowing data collection up to the point just before the self-protection mechanism was triggered, as shown in Figs. S11 and S12. Consequently, the primary peaks of disordered C-(A)-S-H and the amorphous slag (~30° 2θ Cu-Kα equivalent) were excluded from the collected XRD patterns. Since the volume fraction of unreacted slag and C-(A)-S-H gel exceeds 50 %, performing Rietveld quantitative analysis within the current diffraction range could result in an overestimation of the ettringite and LDH phases. To prevent providing potentially misleading results, QXRD was therefore not conducted. A detailed discussion of the experimental limitations associated with specific X-ray diffraction ranges, along with future recommendations for applying this technique to evaluation of cement pastes, can be found in Section 4 of the SI file.

Fig. 1a, b and c present the XRD patterns (recalculated to give angles equivalent to Cu-Kα radiation) shown layer-by-layer (i.e. data integrated over the volume at each specific depth) from the surface to the core of a specimen, before and after exposure to carbonation in the S2-based pastes. The main crystalline product identified is ettringite (powder diffraction file (PDF)# 01-072-0646). As CO₂ tends to diffuse from the surface to the core, creating a concentration gradient, the bound amount
<figure><img src="image_4.png" /></figure>

<figure><img src="image_5.png" /><figcaption>**Fig. 1.** Cu-Kα equivalent XRD patterns of each layer (depths as marked) from the surface to the core of the specimen, for (a) uncarbonated; (b) after cycle one (18 min); and (c) after cycle two (36 min) of carbonation of sodium sulfate activated S2-slag cements. (d) Cumulative XRD patterns for entire specimens before and after carbonation. C-(A)-S-H: aluminium-substituted calcium silicate hydrate, E: ettringite; LDH: Mg-Al-CO₃-LDH; G: gypsum; V: vaterite.</figcaption></figure>

of CO₂ decreases as the distance from the surface increases.

As carbonation proceeds from the specimen surface to its core, the amount of ettringite (proportional to the integrated area of the peak at 2θ = 9.138°) diffracted gradually decreases. A subtle decrease was also observed in the gypsum (PDF#00-001-0385) and vaterite (PDF#00-024-0030) phases, consistent with the progress of carbonation of ettringite in the carbonated cements [67]. Fig. 1d shows the XRD patterns of the entire scanned specimen before and after carbonation. The primary reaction product, a tobermorite-type C-(A)-S-H (PDF# 00-019-0052), is observed. Upon carbonation, the C-(A)-S-H gel exhibits increased crystallinity, as indicated by the shift of the peak position towards a higher diffraction angle, corresponding to a reduced *d*-spacing. This reduction in *d*-spacing may be attributed to the removal of interlayer calcium [15], leading to the formation of a more crystalline C-(A)-S-H gel [17], and/or the collapse of gel pores during the drying process, which facilitates the polymerisation of Si-O-Si bonds between C-S-H sheets [68], resulting in a decreased internal surface area of the C-S-H [69]. Fig. 2 illustrates the spatial distribution of C-(A)-S-H gel before and after carbonation for S2-based pastes. As carbonation progresses, crystalline C-(A)-S-H gel progressively disappears from the exterior towards the interior of the specimen, likely due to the decalcification of this primary binding phase [17,67]. The carbonation of C-(A)-S-H type gel and the changes induced at the nanostructure scale are not well

<figure><img src="image_6.png" /></figure>

<figure><img src="image_7.png" /></figure>

understood and require further investigation.

Fig. 3 shows the 2D (left) and 3D (right) spatial distribution of the ettringite from the surface of the sample to the centre after 2 cycles of carbonation. In general, this phase is no longer detected in the samples exposed to increasing cycles of carbonation. Locally, a lower amount of ettringite at the sample's edges is observed, where there is more CO₂ compared to the core of the specimen. This indicates that ettringite is one of the phases governing the carbonation of sodium sulfate-activated slag cements. This observation is consistent with bulk lab-based XRD results where it has been identified that carbonation leads to a decreased amount of ettringite and consequent degradation of C-A-S-H gel in sodium sulfate-activated slag cement pastes [67].

XRD-CT enables the role of secondary reaction products in mitigating carbonation to be further investigated; layered double hydroxides are known to play an important role in defining the durability of alkali-activated slag cements, although the exact mechanisms are not yet fully understood. A crystalline Mg-Al-CO₃-LDH phase was identified in the diffractograms in Fig. 1b and c, especially in the layers close to the surface (0.1 mm) upon carbonation, which forms from Mg-Al-SO₄-LDH. Interlayer anion interchange between sulfate and carbonate ions happens under accelerated carbonation conditions, forming the observed Mg-Al-CO₃-LDH phase. This phase is more crystalline than its sulfate analogue [70]. A similar phenomenon is observed in the carbonation of
<figure><img src="image_8.png" /><figcaption>**Fig. 2.** Cross-section (left) and corresponding 3D spatial distribution (right) images obtained by XRD-CT showing the distribution of C-(A)-S-H gel in sodium sulfate-activated S2-slag pastes: (a) uncarbonated; (b) after 1 cycle of carbonation (18 min); (c) after 2 cycles of carbonation (36 min). The white boundary (1-pixel size) delineates the specimen edge.</figcaption></figure>

Na₂CO₃-activated slag cements, where OH⁻ groups in meixnerite interchange for CO₃²⁻ upon CO₂ exposure [16]. Fig. 4 shows the spatial distribution of Mg-Al-CO₃-LDH after 1 and 2 cycles of carbonation. This phase preferentially precipitates at the surface of the specimen and increases with increased CO₂ uptake. The formation of an Mg-Al-SO₄-LDH phase was previously hypothesised to take place in Na₂SO₄-activated slag cements based on ²⁷Al MAS NMR results, but this phase has not been detected by lab-based XRD due to its poorly crystalline structure [10,70], and the spatial heterogeneity of such phases has been unexplored in previous works. The formation of a SO₄²⁻ intercalated Mg-Al-LDH indicates that sulfate anions are consumed in the formation of both ettringite and Mg-Al-SO₄-LDH during sodium sulfate-activation of slags. When the sulfate is consumed from the pore solution, the chemical reaction is likely to proceed similarly to a NaOH-activated slag cement system; a similar reaction mechanism has also been hypothesised for Na₂CO₃-activated slag cements [71–73].

The variations in the phase assemblage under carbonation, as shown in Fig. 1b and c at each layer and structural alteration of C-(A)-S-H gel (Fig. 1d) at the entire specimen, are expected to modify the pore network. The changes in the porosity as a function of the radial distance from the surface to the core of the specimen are presented in Fig. 5. As CO₂ penetrates the specimen, the decalcification of ettringite (Fig. 3) and C-(A)-S-H (Fig. 2) results in an increase in porosity, accompanied by the more polymerised Si tetrahedra and the formation of gypsum, CaCO₃ polymorphs, and aluminosilicates [15,67,74,75]. These effects are particularly noticeable in the evaluated cement specimens, especially within the first 0.3 mm from the surface. A decrease in porosity from the CO₂-exposed surface to the core of the sample is consistent with a CO₂ concentration gradient across the specimens, making less CO₂ available to react at its core, consequently preserving its original microstructural features. Interestingly, although ettringite is almost fully consumed on the exposed surface, this has not led to a significant increase in the

porosity compared to the values recorded in the uncarbonated sample centre. This may indicate that the formation of Mg-Al-CO₃-LDH, due to the exchange of carbonate for sulfate in the interlayer space of the LDH phase, compensates for some of the volume change associated with ettringite decomposition. It has been reported that the density of Mg-Al-SO₄-LDH (2.21 g/cm³) is higher than that of Mg-Al-CO₃-LDH (2.02 g/cm³) [76], so the tetrahedral anion SO₄²⁻ in the interlayer causes an increased c lattice constant and unit cell volume compared to the trigonal planar CO₃²⁻ analogue [70,77]. This finding further confirms the role of Mg-Al-LDH phases in resisting carbonation by acting as a CO₂ sink in alkali-activated slag cements, which delays damage to Ca-bearing phases without significantly affecting the microstructure [16,73,78,79].

## 3.2. Tomographic approach to determine changes in microscale pore structure, and predicting transport properties

The precursors and reaction products can be distinguished by their contrast and morphology in µCT images. Fig. 6 shows representative ortho slices, 3D renderings of reconstructed images, segmented phases and pore size distributions in S2-based pastes before and after 1 and 2 carbonation cycles. The unreacted blast furnace slag has an irregular jagged shape [80]. The presence of tiny bright specks originates from the presence of very small particles of metallic iron that are entrained in the production process of blast furnace slag [81,82]. During activation with a sodium sulfate solution, the unreacted slag maintains its original geometry and particle shape, exhibiting a relatively higher level of brightness [31,83], whereas the hydrous (and/or intrinsically porous) reaction products such as C-A-S-H type gel, ettringite and layered double hydroxides (LDH) have lower brightness due to their lower X-ray attenuation coefficient compared to the slag [8,10,84]. Ettringite intermixes with C-A-S-H type gel in the outer reaction product regions,
<figure><img src="image_9.png" /><figcaption>Fig. 3. Cross-section (left) and corresponding 3D spatial distribution (right) images obtained by XRD-CT showing the distribution of ettringite in sodium sulfate-activated S2-slag pastes: (a) uncarbonated; (b) after 1 cycle of carbonation (18 min); (c) after 2 cycles of carbonation (36 min). The white boundary (1-pixel size) delineates the specimen edge.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>Fig. 4. 2D virtual cross-sections (left) and 3D renderings (right) obtained by XRD-CT showing the spatial distribution of Mg-Al-CO₃-LDH in Na₂SO₄-activated S2-slag cements after (a) 1 cycle (18 min) and (b) 2 cycles of carbonation (36 min).</figcaption></figure>

where space was originally filled with the activating solution [85]. Simultaneously, LDH intermixes with C-A-S-H gel within the inner reaction product areas, corresponding to the regions where the slag has reacted [86]. All the reaction products are intermixed on the nanoscale

and C-A-S-H type gel also has nano-sized pores [10,87]. Nevertheless, those reaction products appear homogeneous at the moderate spatial resolution achieved in this study.

The black regions correspond to the pores, which have a range of
<figure><img src="image_11.png" /><figcaption>Fig. 5. Porosity (derived from µCT segmentation and analysis) as a function of distance from the sample surface, for the uncarbonated sample, 1-cycle carbonated sample and 2-cycle carbonated sample.</figcaption></figure>

different sizes. Considering the effect of noise in the tomography data, only the pore size range larger than 2.1 µm was considered, which is 3 times the voxel size. Deleterious ions are mainly transported in pores larger than 10 µm [31,47], therefore the resolution in this study is deemed sufficient to evaluate the pore network of the sodium sulfate-activated slag cements to infer their effect on the durability performance of these cements. For the paste sample produced with slag S2, the observed increase in porosity after 1 cycle of carbonation is small and within the estimated uncertainty for the uncarbonated sample; after 2 cycles of carbonation, the increase of porosity is significant.

Fig. 7 shows the pore size distribution of the 3D segmented pore volume in sodium sulfate-activated slag cements before and after accelerated carbonation for three slags (Fig. 7a, b and c), as well as a summary of porosity, as determined by 3D µCT (Fig. 7d). For all 3 slags, accelerated carbonation led to an increase in overall porosity, which is consistent with previous studies where accelerated carbonated alkali-activated slag cements showed a highly porous microstructure [16,38,87]. This is attributed to the fact that exposure to an elevated CO₂ concentration leads directly to the decalcification of C-(A)-S-H gel (Fig. 2) and other secondary reaction products forming in these cements [79]. Irrespective of the slag used, carbonation gave a reduced pore volume for the pore size range between 2.1 µm and 10 µm. This is associated with the precipitation of CaCO₃ polymorphs in the pore space due to the carbonation of alkali-activated slag cements [16,87,88]. For pores larger than 10 µm, the pore volume increases, contributing to increasing the overall porosity of the materials, which is especially pronounced for the pores larger than 20 µm (Fig. 5). For cements produced with slag S2, 2 cycles of accelerated carbonation resulted in further reduction of pore volume in pores smaller than 10 µm. Extended carbonation leads to the formation of more pores with sizes larger than 10 µm, and a consequent increase in average porosity. The highest porosity values were observed for the S3-based pastes throughout CO₂ exposure. It is related to the lowest Mg/Al ratio (0.5) of S3 slag, with the lowest degree of reaction and amount of reaction products (see the SI file, Section S3). The 3D pore structures will be used in calculating the diffusion tortuosity and formation factor in the following sections.

### 3.2.1. Diffusion tortuosity

Fig. 8a presents the simulated diffusion tortuosity of sodium sulfate-activated slag cements as a function of the slag type, before and after accelerated carbonation, as calculated using a random walker-based method in the reconstructed pore network. Detailed descriptions of

diffusion tortuosity and the random walker method are presented in the SI file, Section 2.3. For the three slag cement pastes, accelerated carbonation leads to a decrease in the calculated tortuosity, which means that more connected pores formed in the pore structure resulted in a shorter pathway for the simulated walkers to move. This is in agreement with previous studies in alkali-activated slag/metakaolin cements carbonated at 3 % CO₂, which showed an increased pore volume after carbonation [89]. The cement paste produced with slag S1, with the highest Mg/Al ratio (0.8), records the highest tortuosity values at all stages.

Fig. 8b shows the relationship between segmented porosity from the µCT images and diffusion tortuosity. The dashed (fitted) curve presents an inverse power law relationship, and a horizontal asymptote to a tortuosity at very high porosity, which means that walkers move in free space without any resistance. A similar trend has also been reported in µCT studies of sodium silicate-activated slag cements [31,38]. The vertical asymptote tends to infinity when the porosity is 0.15, which generally agrees with the porosity value of the non-percolated pores in Portland cement by modelling of cement hydration (0.18) [90] and sodium hydroxide-activated slag cement characterised by ¹H NMR (~0.2) [87]. The similarity of our results with other reported studies further supports the accuracy of the segmentation approach of the µCT results followed in this study.

### 3.2.2. Formation factor

Fig. 9 shows the simulated electrical charge transport and formation factor in sodium sulfate-activated slag cements along the X (Fig. 9a), Y (Fig. 9b) and Z (Fig. 9c) axes, before and after carbonation. The formation factor (also called the electrical tortuosity in the field of battery materials research) is the ratio of the material resistivity to its pore solution resistivity, and it is directly related to the electrically resistive effects of porous materials [91,92]. It is related in principle to transport properties and may also be used to analyse diffusivity within porous cementitious materials [24,93]. This factor can be estimated by simulating electrical conductivity for the segmented pore 3D reconstruction using a finite volume method [94]. Fig. 10a shows the formation factor simulations for sodium sulfate-activated slag cements before and after carbonation, as a function of the slag type used. Accelerated carbonation leads to more connected pores, and consequently, more electrical charge can be transported in the pore structure. A higher formation factor indicates a higher electrical resistivity of the material due to a higher amount of non-conductive solid formed compared to the conductive component (pore space), consistent with a more compact and tortuous microstructure. Before inducing carbonation, large deviations of the formation factor along the three axes were observed, irrespective of the slag used. This is consistent with the partial settlement of solids in the direction of the cement casting along the Z axis before it hardens [95]. The deviation in the X and Y axes is possibly related to some streak and ring artefacts during scanning and reconstruction [96], and slightly affects the calculated results of connectivity of the pore network [31]. After CO₂ exposure, the differences in results along the three axes significantly decrease, meaning more connected pores are forming in the 3D space.

The paste produced with slag S1, with the highest Mg/Al ratio (0.8), has the highest formation factor (Fig. 10a), whereas the S3-based cement shows the lowest value. This result agrees well with experimental results when exposing similar cement pastes to a lower CO₂ concentration (1 % CO₂), where it was identified that sodium sulfate-activated S1 cements possessed the highest carbonation resistance compared to S2 and S3 [67]. With 1 and 2 cycles of accelerated carbonation, increased porosity (Fig. 7d) and large capillary pores (Fig. 7b) led to a reduced formation factor. The simulated formation factors presented here (Fig. 10a) generally align with the published range of previous experimental results for various cementitious materials. These include fresh Portland cement pastes and blended fly ash or slag Portland cements at 90 min of curing (FF ranging from ~2.1 to
<figure><img src="image_12.png" /><figcaption>b</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>c</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>**Fig. 6.** From left to right; representative ortho slices, 3D rendering of reconstructed images, segmented phases and pore structure using WeKa machine learning method, uncarbonated (a) and after cycle 1 (b) and cycle 2 (c) of carbonation in S2-based pastes. The uncertainties are calculated from a probability map in the WeKa segmentation, based on the 1σ value of a Gaussian distribution.</figcaption></figure>
<figure><img src="image_15.png" /><figcaption>Fig. 7. The pore size distribution of 3D segmented pore volumes identified in sodium sulfate-activated slag cements produced with: (a) slag S1; (b) S2; and (c) S3 before and after carbonation. (d) Summary of µCT porosity before and after carbonation.</figcaption></figure>

<figure><img src="image_16.png" /><figcaption>Fig. 8. (a) Diffusion tortuosity of sodium sulfate-activated slag cements before and after accelerated carbonation using slag S1, S2 and S3. (b) Segmented porosity vs. diffusion tortuosity; σ and τ represent the porosity and diffusion tortuosity in the fitted curve, respectively.</figcaption></figure>

~3.1) [91], 29-year-old Portland cement pastes with the same water/solid ratio (FF of ~49.1) [97], sodium silicate-activated slag paste after 28 days of curing using two different water/solid ratios (FF ranging from ~20 to ~60) [93] and Portland cement pastes with added sodium hydroxide at different water/solid ratios, measured from 3 days to 90 days (FF ranging from ~10 to ~60) [98]. It should be noted that the

formation factors of mortar and concrete are two orders of magnitude higher than that of pastes due to the addition of sand and aggregate, as the more insulating components increase the resistivity of the bulk specimens [99,100].

Fig. 10b presents the segmented porosity vs. formation factor of the cement pastes evaluated here. Increased porosity causes a decrease in
<figure><img src="image_17.png" /><figcaption>Fig. 9. Simulated electrical charge transport in sodium sulfate activated S2 slag cements along the a) X, b) Y and c) Z directions, before (left) and after (middle and right) accelerated carbonation. An electrical potential of 1 V is applied along X, Y and Z axis, respectively. FF stands for formation factor, and error estimates are based on the calculated FF in the three directions shown.</figcaption></figure>

<figure><img src="image_18.png" /><figcaption>Fig. 10. (a) Simulated formation factor of sodium sulfate-activated slag cements before and after accelerated carbonation, and (b) relationship between porosity and formation factor; the fitting curve is based on Archie's law. F and σ represent formation factor and porosity, respectively.</figcaption></figure>

formation factor. This trend agrees with experimental observation in other cementitious materials [91,97,98]. The fitted curve is based on Archie's law and gives a cementation factor of ~2.4, which agrees well with the published value of ~2.5 for Portland-based materials having lower porosity [97,101], the relationship of formation factor and

porosity for reservoir rocks [102] and highly cemented rocks (1.3 to 2.2) [91,103]. The cementation factor is related to the type of porous material and is normally obtained by fitting a large number of datasets from experiments [91,102]. It seems that independent of the chemical composition of the slag used or exposed CO₂ concentration, there is a
fixed relationship between the formation factor and porosity value in sodium sulfate-activated slag cements. However, more data for AASCs are needed to establish an empirical equation for assessing the durability of these materials.

It must be borne in mind that the calculation of diffusion tortuosity and formation factor is influenced by both the pore size distribution and connectivity of the segmented porous space from µCT. Here the ~2 µm pore resolution limit of the µCT arrangement means that gel pores and finer capillary pores are not recorded. From a gas permeability viewpoint, the largest pores are probably the most important, whereas for water permeability the capillary action may mean pores finer than can be measured by µCT may become important [104].

## 4. Significance and implications

This study elucidates the distinct microstructural features of sodium sulfate-activated slag cements. Before exposure to CO₂, sodium sulfate-activated slag cements, regardless of slag composition, exhibit total porosities (above the spatial resolution) of approximately 20 % after 28 days of curing. This finding is consistent with reported porosities for sodium hydroxide-activated slag cements (~17 % using scanning electron microscopy) [105], sodium silicate-activated slag cements (~18 % using µCT) [38], and sodium carbonate-activated slag cements (~16 % using µCT) [106], and sodium sulfate-activated slag cements at 28 days (~22 % using MIP) [10]. Additionally, the porosity of alkali-activated slag cements is comparable to, or even lower than, that of metakaolin-blended Portland cements cured for 90 days (~30 % to ~35 % using MIP) [107], fly ash blended slag cements with 50 % Portland cement substitution for 28 days (~33 % using MIP) [108], and slag blended cements with 70 % clinker replacement after 1 year of curing (~18 % using MIP) [109]. Exposure to an elevated CO₂ concentration causes an increase in percolated porosity in sodium sulfate-activated slag cements, consistent with what has been observed in other alkali-activated slag cements [16,38,87,88]. Similar results have also been found in blended Portland cements with over 50 % of the Portland cement replaced by slag [109], fly ash [109,110] or calcined clay [111]. This then demonstrates that the behaviour of the porosity developed in alkali sulfate-activated slag cements is comparable to that of other AASCs, and similar to blended Portland cements.

The chemical durability of cementitious materials is strongly dependent on the ability of aggressive ions to ingress the material, which is influenced significantly by the tortuosity of its pore structure and consequently its effect on the diffusivity of ions [23]. Accurate estimation of tortuosity based on non-destructive 3D imaging enables the determination of the diffusivity potential of a given material. This includes considering factors like porosity (the volume available for ionic transport), pore size distribution and pore morphology, which significantly affect tortuosity values, and are listed in approximately decreasing order of impact [92]. There are few studies determining the formation factor values of alkali-activated slag cements in the literature. The simulated formation factors of sodium sulfate-activated slag cements in this study show similar trends, and comparable values, to those reported for commercial blended Portland slag cements [91,97]. Laboratory-based rapid bulk electrical resistance tests reported in the literature are restricted to testing the specimen for a few hours [112], under non-steady state ionic transport conditions that might lead to an increased formation factor. Modelling the formation factor based on VOI from µCT images enables the calculation of the resistivity in steady-state transport, minimising errors in the calculations. The simulation workflow developed in this study is a significant step forward in developing a rapid and reliable estimation of the formation factor in cementitious materials.

This study demonstrates the capability of Mg-Al-LDH phases in sodium sulfate-activated slag cements in entrapping carbonate ions, acting as in-situ CO₂ sinks. By controlling the mix design, especially the Mg and Al content present in the slag, the amount of Mg-Al-LDH phases can be

controlled to optimise the reaction kinetics [73], carbonation [16,113-115] and chloride uptake [116-118] in alkali-activated materials. Lower porosity and higher tortuosity values after carbonation in S1-based pastes compared to S2 and S3 indicate that using slags with higher Mg/Al ratio leads to an increased carbonation resistance. This is mainly because, upon alkali-activation of slags with elevated Mg contents, the formation of more LDH phases is favoured [70]. Intercalation of sulfate ions within the LDH phase causes an increased pH in the pore solution, as in Portland cement-based materials [6]. Further, this study proves the potential of Mg-Al-LDH phases in enhancing carbonation resistance and increasing the CO₂ uptake potential of cementitious materials, and offers insights into tailoring compositions to absorb atmospheric carbon, thereby contributing to sustainable construction practices.

## 5. Conclusion

This study pioneers the application of synchrotron-based X-ray micron scale computed tomography (µCT) and X-ray diffraction computed tomography (XRD-CT) to assess the effects of CO₂ exposure in Na₂SO₄-activated slag cement pastes. The carbonation of Na₂SO₄-activated slag leads to the decalcification of C-(A)-S-H gel, ettringite decomposition, and the formation of crystalline carbonate-intercalated Mg-Al-LDH phases. The spatially resolved ettringite distribution confirms that the amount of ettringite decreases upon carbonation. Synchrotron XRD in this study has, for the first time, provided irrefutable evidence of the formation of the Mg-Al-SO₄-LDH phase in sodium sulfate-activated slag cements, and its role as a CO₂ sink, by anion exchange of sulfate by carbonate in the interlayer space of the LDH phase.

Carbonation induced an increase in porosity and a decrease in tortuosity in all of the alkali-activated slag pastes evaluated. Alkali sulfate-activated slag cements achieve comparable pore networks with lower porosity and higher tortuosity compared to blended Portland cements. This was more noticeable when using a slag with the highest Mg/Al ratio (0.8), which also presented the highest carbonation resistance. A correlation between the simulated tortuosity and 3D porosity, approximately following Archie's law was identified. A fitted empirical equation for calculating the diffusion tortuosity and formation factor of these cements, as a function of the total porosity of the materials assessed, is proposed.

This study underscores the important role of XRD-CT and µCT in providing comprehensive physicochemical, non-destructive, three-dimensional insights into the carbonation processes of cements. As automated (machine learning) data analysis tools improve, this should simplify and expedite XRD-CT and µCT phase and pore analyses, making it easier to obtain robust quantitative information. Our future research will aim to integrate XRD-CT and µCT data into predictive reactive-transport life service models for the cementitious systems evaluated in this study.

### CRediT authorship contribution statement

**Zengliang Yue:** Writing – review & editing, Writing – original draft, Investigation, Formal analysis, Data curation, Conceptualization.
**Zixian Su:** Writing – review & editing, Writing – original draft, Investigation, Formal analysis, Data curation.
**Partha P. Paul:** Writing – review & editing, Writing – original draft, Supervision, Formal analysis, Data curation.
**Alastair T.M. Marsh:** Writing – review & editing, Writing – original draft, Supervision, Investigation, Data curation.
**Alice Macente:** Writing – review & editing, Supervision, Resources.
**Marco Di Michiel:** Writing – review & editing, Resources, Investigation.
**John L. Provis:** Writing – review & editing, Resources, Methodology, Investigation, Funding acquisition, Conceptualization.
**Philip J. Withers:** Writing – review & editing, Supervision, Resources, Methodology, Funding acquisition, Conceptualization.
**Susan A. Bernal:** Writing – review & editing, Supervision, Resources, Project administration, Methodology,
## Funding acquisition, Formal analysis, Conceptualization.

## Declaration of competing interest

The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: Susan A. Bernal reports financial support was provided by Engineering and Physical Sciences Research Council. Philip J. Withers reports financial support was provided by Engineering and Physical Sciences Research Council. John L. Provis reports financial support was provided by Engineering and Physical Sciences Research Council. The co-authors Susan A. Bernal and John L. Provis serve as Editors of Cement and Concrete Research. If there are other authors, they declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Acknowledgements

This research was sponsored by the Engineering and Physical Sciences Research Council (EPSRC) Early Career Fellowship EP/R001642/1. Z. Yue is grateful to the University of Leeds for his PhD studentship. The authors acknowledge Beamline ID15A at The European Synchrotron (ESRF) for the experiments (MA-5424, https://doi.esrf.fr/10.1515/ESRF-ES-895887564). PJW and PPP are grateful for the support for the Henry Royce Institute for Advanced Materials, established through EPSRC Grants EP/R00661X/1, EP/P025498/1 and EP/P025021/1 and an EPSRC centre-to-centre collaboration with the ESRF through grants EP/W003333/1 and EP/T02593X/1 as well as access to the NXCT funded through EP/X034941/1 and EP/T02593X/1. The participation of JLP in this work was supported by in part by EPSRC under grant EP/T008407/1. The data associated with this paper are openly available at https://doi.esrf.fr/10.1515/ESRF-ES-895887564.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.cemconres.2024.107716.

## Data availability

The data associated with this paper are openly available from the University of Leeds Research Data Repository at https://doi.org/10.5518/1569, and ESRF database at https://doi.esrf.fr/10.1515/ESRF-ES-895887564.

## References

[1] R.J. Myers, S.A. Bernal, J.L. Provis, A thermodynamic model for C-(N-)A-S-H gel: CNASH ss. Derivation and validation, Cem. Concr. Res. 66 (2014) 27-47.

[2] J.L. Provis, Alkali-activated materials, Cem. Concr. Res. 114 (2018) 40-48.

[3] Y. Zhu, Z. Zheng, Y. Deng, C. Shi, Z. Zhang, Advances in immobilization of radionuclide wastes by alkali activated cement and related materials, Cem. Concr. Comp. 126 (2022) 104377.

[4] D. Jiapei, B. Yuhuan, C. Xuechao, S. Zhonghou, S. Baojiang, Utilization of alkali-activated slag based composite in deepwater oil well cementing, Construct. Build Mater. 186 (2018) 114-122.

[5] M. Samarakoon, P. Ranjith, W.H. Duan, B.K. Chen, M. Khandelwal, Early age properties of alkali-activated cement and class G cement under different saturation conditions in oil well applications, Construct. Build Mater. 271 (2021) 121543.

[6] J. Fu, A.M. Jones, M.W. Bligh, C. Holt, L.M. Keyte, F. Moghaddam, S.J. Foster, T. D. Waite, Mechanisms of enhancement in early hydration by sodium sulfate in a slag-cement blend-insights from pore solution chemistry, Cem. Concr. Res. 135 (2020) 106110.

[7] N. Mobasher, S.A. Bernal, O.H. Hussain, D.C. Apperley, H. Kinoshita, J.L. Provis, Characterisation of Ba(OH)₂-Na₂SO₄-blast furnace slag cement-like composites for the immobilisation of sulfate bearing nuclear wastes, Cem. Concr. Res. 66 (2014) 64-74.

[8] N. Mobasher, S.A. Bernal, J.L. Provis, Structural evolution of an alkali sulfate activated slag cement, J. Nucl. Mater. 468 (2016) 97-104.

[9] B. Akturk, S. Nayak, S. Das, A.B. Kizilkanat, Microstructure and strength development of sodium carbonate-activated blast furnace slags, J. Mater. Civ. Eng. 31 (2019) 04019283.

[10] A.T.M. Marsh, Z. Yue, Y. Dhandapani, K. Button, S. Adu-Amankwah, S.A. Bernal, Influence of limestone addition on sodium sulphate activated blast furnace slag cements, Construct. Build Mater. 360 (2022) 129527.

[11] A.M. Rashad, Influence of different additives on the properties of sodium sulfate activated slag, Construct. Build Mater. 79 (2015) 379-389.

[12] A. Rashad, Y. Bai, P. Basheer, N. Milestone, N. Collier, Hydration and properties of sodium sulfate activated slag, Cem. Concr. Comp. 37 (2013) 20-29.

[13] C.-F. Lu, W. Wang, Q.-T. Li, M. Hao, Y. Xu, Effects of micro-environmental climate on the carbonation depth and the pH value in fly ash concrete, J. Clean. Prod. 181 (2018) 309-317.

[14] K. De Weerdt, G. Plusquellec, A.B. Revert, M. Geiker, B. Lothenbach, Effect of carbonation on the pore solution of mortar, Cem. Concr. Res. 118 (2019) 38-56.

[15] T.F. Sevelsted, J. Skibsted, Carbonation of C-S-H and C-A-S-H samples studied by ¹³C, ²⁷Al and ²⁹Si MAS NMR spectroscopy, Cem. Concr. Res. 71 (2015) 56-65.

[16] X. Ke, M. Criado, J.L. Provis, S.A. Bernal, Slag-based cements that resist damage induced by carbon dioxide, ACS Sustain. Chem. Eng. 6 (2018) 5067-5075.

[17] S. Park, B. Lothenbach, J.G. Jang, H.-K. Kim, N. Lee, Thermodynamic modeling and experimental study of carbonation of alkali-activated slag cements, ACS Sustain. Chem. Eng. 11 (2023) 4049-4063.

[18] S. von Greve-Dierfeld, B. Lothenbach, A. Vollpracht, B. Wu, B. Huet, C. Andrade, C. Medina, C. Thiel, E. Gruyaert, H. Vanoutrive, I.F. Saéz del Bosque, I. Ignjatovic, J. Elsen, J.L. Provis, K. Scrivener, K.-C. Thielen, K. Sideris, M. Zajac, N. Alderete, Ö. Cizer, P. Van den Heede, R.D. Hooton, S. Kamali-Bernard, S.A. Bernal, Z. Zhao, Z. Shi, N. De Belie, Understanding the carbonation of concrete with supplementary cementitious materials: a critical review by RILEM TC 281-CCC, Mater. Struct. 53 (2020) 136.

[19] G.J. Gluth, X. Ke, A. Vollpracht, L. Weiler, S.A. Bernal, M. Cyr, K. Dombrowski-Daube, D.A. Geddes, C. Grengg, C. Le Galliard, M. Nedeljkovic, J.L. Provis, L. Valentini, B. Walkley, Carbonation rate of alkali-activated concrete and high-volume SCM concrete: a literature data analysis by RILEM TC 281-CCC, Mater. Struct. 55 (2022) 225.

[20] Y. Song, G. Dai, J. Zhou, Z. Bian, L. Zhao, L. Song, Characterizing porous volume of cement-based concrete by multiscale image analysis, J. Mater. Civil Eng. 32 (2020) 04020267.

[21] C. Liu, M. Zhang, Multiscale modelling of ionic diffusivity in unsaturated concrete accounting for its hierarchical microstructure, Cem. Concr. Res. 156 (2022) 106766.

[22] P.J. McDonald, O. Istok, M. Janota, A.M. Gajewicz-Jaromin, D.A. Faux, Sorption, anomalous water transport and dynamic porosity in cement paste: a spatially localised ¹H NMR relaxation study and a proposed mechanism, Cem. Concr. Res. 133 (2020) 106045.

[23] E. Samson, J. Marchand, Modeling the transport of ions in unsaturated cement-based materials, Comput. Struct. 85 (2007) 1740-1756.

[24] O.B. Isgor, W.J. Weiss, A nearly self-sufficient framework for modelling reactive-transport processes in concrete, Mater. Struct. 52 (2019) 3.

[25] E. Samson, J. Marchand, K.A. Snyder, J. Beaudoin, Modeling ion and fluid transport in unsaturated cement systems in isothermal conditions, Cem. Concr. Res. 35 (2005) 141-153.

[26] M. Auroy, S. Poyet, P. Le Bescop, J.-M. Torrenti, T. Charpentier, M. Moskura, X. Bourbon, Impact of carbonation on unsaturated water transport properties of cement-based materials, Cem. Concr. Res. 74 (2015) 44-58.

[27] T. Bakharev, J.G. Sanjayan, Y.-B. Cheng, Sulfate attack on alkali-activated slag concrete, Cem. Concr. Res. 32 (2002) 211-216.

[28] H. Ye, Z. Chen, L. Huang, Mechanism of sulfate attack on alkali-activated slag: the role of activator composition, Cem. Concr. Res. 125 (2019) 105868.

[29] Z. Chen, H. Ye, Understanding the impact of main seawater ions and leaching on the chloride transport in alkali-activated slag and Portland cement, Cem. Concr. Res. 164 (2023) 107063.

[30] J. Shi, M. Wu, J. Ming, Long-term corrosion resistance of reinforcing steel in alkali-activated slag mortar after exposure to marine environments, Corros. Sci. 179 (2021) 109175.

[31] J.L. Provis, R.J. Myers, C.E. White, V. Rose, J.S.J. van Deventer, X-ray microtomography shows pore structure and tortuosity in alkali-activated binders, Cem. Concr. Res. 42 (2012) 855-864.

[32] M.C.G. Juenger, H.M. Jennings, The use of nitrogen adsorption to assess the microstructure of cement paste, Cem. Concr. Res. 31 (2001) 883-892.

[33] R. Kumar, B. Bhattacharjee, Study on some factors affecting the results in the use of MIP method in concrete research, Cem. Concr. Res. 33 (2003) 417-424.

[34] A.C.A. Muller, Characterization of Porosity and C-S-H in Cement Pastes by ¹H NMR, Doctoral Thesis, EPFL, Switzerland, 2014, p. 208.

[35] A. Baral, C. Pesce, A.S. Yorkshire, Z. Zhakiyeva, R. Snellings, T. Hanein, J. L. Provis, A. Peys, Characterisation of iron-rich cementitious materials, Cem. Concr. Res. 177 (2024) 107419.

[36] S. Brisard, M. Serdar, P.J.M. Monteiro, Multiscale X-ray tomography of cementitious materials: a review, Cem. Concr. Res. 128 (2020) 105824.

[37] P.J. Withers, C. Bouman, S. Carmignato, V. Cnudde, D. Grimaldi, C.K. Hagen, E. Maire, M. Manley, A. Du Plessis, S.R. Stock, X-ray computed tomography, Nature Reviews Methods Primers 1 (2021) 18.

[38] S.Y. Wang, E. McCaslin, C.E. White, Effects of magnesium content and carbonation on the multiscale pore structure of alkali-activated slags, Cem. Concr. Res. 130 (2020) 105979.

[39] K. Xu, A.S. Tremsin, J. Li, D.M. Ushizima, C.A. Davy, A. Bouterf, Y.T. Su, M. Marroccoli, A.M. Mauro, M. Osanna, Microstructure and water absorption of
ancient concrete from Pompeii: an integrated synchrotron microtomography and neutron radiography characterization, Cem. Concr. Res. 139 (2021) 106282.

[40] G. Fang, Q. Wang, M. Zhang, In-situ X-ray tomographic imaging of microstructure evolution of fly ash and slag particles in alkali-activated fly ash-slag paste, Compos. Part B Eng. 224 (2021) 109221.

[41] M.A.B. Promentilla, T. Sugiyama, T. Hitomi, N. Takeda, Quantification of tortuosity in hardened cement pastes using synchrotron-based X-ray computed microtomography, Cem. Concr. Res. 39 (2009) 548–557.

[42] J.T. Bushberg, J.M. Boone, The Essential Physics of Medical Imaging, Lippincott Williams & Wilkins, 2011.

[43] A. Gervaise, B. Osemont, S. Lecocq, A. Noel, E. Micard, J. Felblinger, A. Blum, CT image quality improvement using adaptive iterative dose reduction with wide-volume acquisition on 320-detector CT, Eur. Radiol. 22 (2012) 295–301.

[44] D. Seo, F. Tomizato, H. Toda, K. Uesugi, A. Takeuchi, Y. Suzuki, M. Kobayashi, Spatial resolution of synchrotron X-ray microtomography in high energy range: effect of X-ray energy and sample-to-detector distance, Appl. Phys. Lett. 101 (2012).

[45] L.R. Koetzier, D. Mastrodicasa, T.P. Szczykutowicz, N.R. van der Werf, A.S. Wang, V. Sandfort, A.J. van der Molen, D. Fleischmann, M.J. Willemink, Deep learning image reconstruction for CT: technical principles and clinical prospects, Radiology 306 (2023) e221257.

[46] S. Mindess, F. Young, D. Darwin, Concrete, 2nd Edition, Prentice Hall, 2003, p. 585.

[47] R.R. Lloyd, J.L. Provis, K.J. Smeaton, J.S. van Deventer, Spatial distribution of pores in fly ash-based inorganic polymer gels visualised by Wood's metal intrusion, Microporous Mesoporous Mater. 126 (2009) 32–39.

[48] O. Wenzel, M. Schwotzer, E. Müller, V.S.K. Chakradhananu, T. Scherer, A. Gerdes, Investigating the pore structure of the calcium silicate hydrate phase, Mater. Charact. 133 (2017) 133–137.

[49] K. Kjellsen, A. Monsøy, K. Isachsen, R. Detwiler, Preparation of flat-polished specimens for SEM-backscattered electron imaging and X-ray microanalysis—importance of epoxy impregnation, Cem. Concr. Res. 33 (2003) 611–616.

[50] S.C. Mayo, A.W. Stevenson, S.W. Wilkins, In-line phase-contrast X-ray imaging and tomography for materials science, Materials 5 (2012) 937–965.

[51] S. Shirani, A. Cuesta, A. Morales-Cantero, I. Santacruz, A. Diaz, P. Trtik, M. Holler, A. Rack, B. Lukic, E. Brun, 4D nanoimaging of early age cement hydration, Nat. Commun. 14 (2023) 2652.

[52] J.L. Provis, V. Rose, R.P. Winarski, J.S.J. van Deventer, Hard X-ray nanotomography of amorphous aluminosilicate cements, Scr. Mater. 65 (2011) 316–319.

[53] G. Harding, J. Kosanetzky, U. Neitzel, X-ray diffraction computed tomography, Med. Phys. 14 (1987) 515–525.

[54] A. Vamvakeros, S. Jacques, M. Di Michiel, D. Matras, V. Middelkoop, I. Ismagilov, E. Matus, V. Kuznetsov, J. Drnec, P. Senecal, SD operando tomographic diffraction imaging of a catalyst bed, Nat. Commun. 9 (2018) 4751.

[55] G. Artioli, L. Valentini, M. Voltolini, M.C. Dalconi, G. Ferrari, V. Russo, Direct imaging of nucleation mechanisms by synchrotron diffraction micro-tomography: superplasticizer-induced change of C–S–H nucleation in cement, Cryst. Growth Des. 15 (2015) 20–23.

[56] M. Voltolini, M.C. Dalconi, G. Artioli, M. Parisatto, L. Valentini, V. Russo, A. Bonnin, R. Tucoulou, Understanding cement hydration at the microscale: new opportunities from pencil-beam synchrotron X-ray diffraction tomography, J. Appl. Cryst. 46 (2013) 142–152.

[57] G. Artioli, T. Cerulli, G. Cruciani, M. Dalconi, G. Ferrari, M. Parisatto, A. Rack, R. Tucoulou, X-ray diffraction microtomography (XRD-CT), a novel tool for non-invasive mapping of phase development in cement materials, Anal. Bioanal. Chem. 397 (2010) 2131–2136.

[58] F. Claret, S. Grangeon, A. Loschetter, C. Tournassat, W. De Nolf, N. Harker, F. Boulahya, S. Gaboreau, Y. Linard, X. Bourbon, Deciphering mineralogical changes and carbonation development during hydration and ageing of a consolidated ternary blended cement paste, IUCrJ 5 (2018) 150–157.

[59] S. Brunauer, P.H. Emmett, E. Teller, Adsorption of gases in multimolecular layers, J. Am. Chem. Soc. 60 (1938) 309–319.

[60] S.A. Bernal, J.L. Provis, D.G. Brice, A. Kilcullen, P. Duxson, J.S.J. van Deventer, Accelerated carbonation testing of alkali-activated binders significantly underestimates service life: the role of pore solution chemistry, Cem. Concr. Res. 42 (2012) 1317–1326.

[61] G.B.M. Vaughan, R. Baker, R. Barret, J. Bonnefoy, T. Buslaps, S. Checchia, D. Duran, F. Fihman, P. Got, J. Kieffer, S.A.J. Kimber, K. Martel, C. Morawe, D. Mottin, E. Papillon, S. Petitdemange, A. Vamvakeros, J.-P. Vieux, M. Di Michiel, ID15A at the ESRF - a beamline for high speed operando X-ray diffraction, diffraction tomography and total scattering, J. Synchrotron Radiat. 27 (2020) 515–528.

[62] M.P. Erb, N.P. McKay, N. Steiger, S. Dee, C. Hancock, R.F. Ivanovic, L.J. Gregoire, P. Valdes, Reconstructing Holocene temperatures in time and space using paleoclimate data assimilation, Clim. Past 18 (2022) 2599–2629.

[63] J. Kieffer, D. Karkoulis, PyFAI, a versatile library for azimuthal regrouping, J. Phys. Conf. Ser. 425 (2013) 202012.

[64] I. Arganda-Carreras, V. Kaynig, C. Rueden, K.W. Eliceiri, J. Schindelin, A. Cardona, H. Sebastian Seung, Trainable Weka Segmentation: a machine learning tool for microscopy pixel classification, Bioinformatics 33 (2017) 2424–2426.

[65] T. Wagner, D.A. Kulik, F.F. Hingerl, S.V. Dmytrieva, GEM-Selektor geochemical modeling package: TSolMod library and data interface for multicomponent phase models, Can. Miner. 50 (2012) 1173–1195.

[66] T.G. Tranter, M.D. Kok, M. Lam, J.T. Gostick, pytax: a simple and efficient random walk implementation for calculating the directional tortuosity of images, SoftwareX 10 (2019) 100277.

[67] Z. Yue, Y. Dhandapani, S.A. Bernal, Structural alterations in alkali-sulfate-activated slag cement pastes induced by natural and accelerated carbonation, Cem. Concr. Res. 187 (2025) 107713.

[68] J. Yin, W. Li, J. Wang, X. Kong, Irreversible microstructural changes of calcium silicate hydrate during the first drying-resaturation cycle, Cem. Concr. Res. 163 (2023) 107032.

[69] J.J. Thomas, A.J. Allen, H.M. Jennings, Structural changes to the calcium-silicate-hydrate gel phase of hydrated cement with age, drying, and resaturation, J. Am. Ceram. Soc. 91 (2008) 3362–3369.

[70] E. Bernard, W.J. Zucha, B. Lothenbach, U. Mäder, Stability of hydrotalcite (Mg-Al layered double hydroxide) in presence of different anions, Cem. Concr. Res. 152 (2022) 106674.

[71] S.A. Bernal, Advances in near-neutral salts activation of blast furnace slags, RILEM Tech. Letters 1 (2016) 39–44.

[72] S.A. Bernal, J.L. Provis, R.J. Myers, R. San Nicolas, J.S.J. van Deventer, Role of carbonates in the chemical evolution of sodium carbonate-activated slag binders, Mater. Struct. 48 (2015) 517–529.

[73] X. Ke, S.A. Bernal, J.L. Provis, Controlling the reaction kinetics of sodium carbonate-activated slag cements using calcined layered double hydroxides, Cem. Concr. Res. 81 (2016) 24–37.

[74] T. Nishikawa, K. Suzuki, S. Ito, K. Sato, T. Takebe, Decomposition of synthesized ettringite by carbonation, Cem. Concr. Res. 22 (1992) 6–14.

[75] T. Grounds, H. Midgley, D. Novell, Carbonation of ettringite by atmospheric carbon dioxide, Thermochim. Acta 135 (1988) 347–352.

[76] D.P. Prentice, L. Gomez-Zamorano, M. Balonis, B. Erdemli, K. Ellison, N. Neithalath, D. Simonetti, G. Sant, The effects of (di-, tri-valent)-cation partitioning and intercalant anion-type on the solubility of hydrotalcites, J. Am. Ceram. Soc. 103 (2020) 6025–6039.

[77] A. Fahami, G.W. Beall, Mechanosynthesis and characterization of hydrotalcite like Mg-Al-SO₄-LDH, Mater. Lett. 165 (2016) 192–195.

[78] X. Ke, S.A. Bernal, O.H. Hussein, J.L. Provis, Chloride binding and mobility in sodium carbonate-activated slag pastes and mortars, Mater. Struct. 50 (2017) 252.

[79] X. Ke, S.A. Bernal, J.L. Provis, B. Lothenbach, Thermodynamic modelling of phase evolution in alkali-activated slag cements exposed to carbon dioxide, Cem. Concr. Res. 136 (2020) 106158.

[80] P. Duxson, J.L. Provis, Designing precursors for geopolymer cements, J. Am. Ceram. Soc. 91 (2008) 3864–3869.

[81] J. Ahmad, K.J. Kontoleon, A. Majdi, M.T. Naqash, A.F. Deifalla, N. Ben Kahla, H. F. Isleem, S.M. Qaidi, A comprehensive review on the ground granulated blast furnace slag (GGBS) in concrete production, Sustainability 14 (2022) 8783.

[82] S.A. Bernal, V. Rose, J.L. Provis, The fate of iron in blast furnace slag particles during alkali-activation, Mater. Chem. Phys. 146 (2014) 1–5.

[83] R. San Nicolas, S.A. Bernal, R. Mejia de Gutiérrez, J.S.J. van Deventer, J.L. Provis, Distinctive microstructural features of aged sodium silicate-activated slag concretes, Cem. Concr. Res. 65 (2014) 41–51.

[84] R.R. Lloyd, J.L. Provis, J.S.J. Van Deventer, Microscopy and microanalysis of inorganic polymer cements: 2: the gel binder, J. Mater. Sci. 44 (2009) 620–631.

[85] I.G. Richardson, S. Li, Composition and structure of an 18-year-old 5M KOH-activated ground granulated blast-furnace slag paste, Construct. Build Mater. 168 (2018) 404–411.

[86] I.G. Richardson, G.W. Groves, Microstructure and microanalysis of hardened cement pastes involving ground granulated blast-furnace slag, J. Mater. Sci. 27 (1992) 6204–6212.

[87] N. Li, N. Farzadnia, C. Shi, Microstructural changes in alkali-activated slag mortars induced by accelerated carbonation, Cem. Concr. Res. 100 (2017) 214–226.

[88] Z. Shi, C. Shi, S. Wan, N. Li, Z. Zhang, Effect of alkali dosage and silicate modulus on carbonation of alkali-activated slag mortars, Cem. Concr. Res. 113 (2018) 55–64.

[89] S.A. Bernal, R. Mejia de Gutiérrez, J.L. Provis, Engineering and durability properties of concretes based on alkali-activated granulated blast furnace slag/ metakaolin blends, Construct. Build Mater. 33 (2012) 99–108.

[90] D.P. Bentz, E.J. Garboczi, Percolation of phases in a three-dimensional cement paste microstructural model, Cem. Concr. Res. 21 (1991) 325–344.

[91] H. Sallehi, P. Ghods, O.B. Isgor, Formation factor of fresh cementitious pastes, Cem. Concr. Compos. 91 (2018) 174–188.

[92] L. Holzer, P. Marmet, M. Fingerle, A. Wiegmann, M. Neumann, V. Schmidt. Tortuosity and Microstructure Effects in Porous Media: Classical Theories, Empirical Data and Modern Methods, Springer Nature, Cham, 2023.

[93] X. Zhu, Q. Li, K. Yang, S. Mu, Z. Zhang, B. Magee, C. Yang, M. Basheer, Characterisation of temporal variations of alkali-activated slag cement property using microstructure features and electrical responses, Construct. Build Mater. 261 (2020) 119884.

[94] M. Darwish, F. Moukalled. The Finite Volume Method in Computational Fluid Dynamics: An Advanced Introduction With OpenFOAM® and Matlab®, Springer, Cham, 2016.

[95] J.E. Vigor, S.A. Bernal, X. Xiao, J.L. Provis, Automated correction for the movement of suspended particulate in microtomographic data, Chem. Eng. Sci. 223 (2020) 115736.

[96] G.R. Davis, J.C. Elliott, Artefacts in X-ray microtomography of materials, Mater. Sci. Technol. 22 (2006) 1011–1018.
[97] P. Tumidajski, A. Schumacher, S. Perron, P. Gu, J. Beaudoin, On the relationship between porosity and electrical resistivity in cementitious systems, Cem. Concr. Res. 26 (1996) 539–544.

[98] Y. Bu, J. Weiss, The influence of alkali content on the electrical resistivity and transport properties of cementitious materials, Cem. Concr. Comp. 51 (2014) 49–58.

[99] Z. Bajja, W. Dridi, B. Larbi, P. Le Bescop, The validity of the formation factor concept from through-out diffusion tests on Portland cement mortars, Cem. Concr. Comp. 63 (2015) 76–83.

[100] H. Ma, D. Hou, Z. Li, Two-scale modeling of transport properties of cement paste: formation factor, electrical conductivity and chloride diffusivity, Comput. Mater. Sci. 110 (2015) 270–280.

[101] B.J. Christensen, T. Coverdale, R.A. Olson, S.J. Ford, E.J. Garboczi, H. M. Jennings, T.O. Mason, Impedance spectroscopy of hydrating cement-based materials: measurement, interpretation, and application, J. Am. Ceram. Soc. 77 (1994) 2789–2804.

[102] P.F. Worthington, The uses and abuses of the Archie equations, 1: the formation factor-porosity relationship, J. Appl. Geophys. 30 (1993) 215–228.

[103] E. Atkins Jr., G.H. Smith, The significance of particle shape in formation resistivity factor-porosity relationships, J. Petrol. Tech. 13 (1961) 285–291.

[104] Z.-L. Jiang, Y.-J. Pan, J.-F. Lu, Y.-C. Wang, Pore structure characterization of cement paste by different experimental methods and its influence on permeability evaluation, Cem. Concr. Res. 159 (2022) 106892.

[105] M. Ben Haha, B. Lothenbach, G. Le Saout, F. Winnefeld, Influence of slag chemistry on the hydration of alkali-activated blast-furnace slag—part II: effect of Al₂O₃, Cem. Concr. Res. 42 (2012) 74–83.

[106] Z. Yue, Z. Su, A.T.M. Marsh, T.L. Burnett, S.A. Bernal, P.J. Withers, Characterization of pore structure evolutions in alkali activated slag cements using time-lapsed in-situ X-ray computed tomography and scanning electron microscopy, in: 42nd IOM3 Annual Cement & Concrete Science Conference, London, United Kingdom, 11–12 September 2023.

[107] F. Zunino, K. Scrivener, Microstructural developments of limestone calcined clay cement (LC3) pastes after long-term (3 years) hydration, Cem. Concr. Res. 153 (2022) 106693.

[108] Z. Yu, G. Ye, The pore structure of cement paste blended with fly ash, Construct. Build Mater. 45 (2013) 30–35.

[109] B. Wu, G. Ye, Development of porosity of cement paste blended with supplementary cementitious materials after carbonation, Construct. Build Mater. 145 (2017) 52–61.

[110] A. Morandeau, M. Thiéry, P. Dangla, Impact of accelerated carbonation on OPC cement paste blended with fly ash, Cem. Concr. Res. 67 (2015) 226–236.

[111] V. Shah, S. Bishnoi, Analysis of pore structure characteristics of carbonated low-clinker cements, Transp. Porous Media 124 (2018) 861–881.

[112] R. Spragg, C. Villani, K. Snyder, D. Bentz, J.W. Bullard, J. Weiss, Factors that influence electrical resistivity measurements in cementitious systems, Transp. Res. Rec. 2342 (2013) 90–98.

[113] E.R. McCaslin, C.E. White, A parametric study of accelerated carbonation in alkali-activated slag, Cem. Concr. Res. 145 (2021) 106454.

[114] M. Zhang, F. Wang, Y. Long, L. Yu, C. Yang, M. Wen, X. Yu, Y. Tian, M. Zhao, Improving the carbonation resistance of alkali-activated slag mortars by calcined Mg/Al layered double hydroxides, Appl. Clay Sci. 216 (2022) 106379.

[115] R. Xu, H. Wang, Q. Zha, J. Lin, Improving the carbonation resistance of alkali-activated slag mortars with different additives: experimental evaluations, Construct. Build Mater. 366 (2023) 130197.

[116] X. Ke, S.A. Bernal, O.H. Hussein, J.L. Provis, Chloride binding and mobility in sodium carbonate-activated slag pastes and mortars, Mater. Struct. 50 (2017) 252.

[117] Y. Zuo, Effect of chloride salt on the phase evolution in alkali-activated slag cements through thermodynamic modelling, Appl. Geochem. 136 (2022) 105169.

[118] T. Liu, Q. Yu, H. Brouwers, In-situ formation of layered double hydroxides (LDHs) in sodium aluminate activated slag: the role of Al-O tetrahedra, Cem. Concr. Res. 153 (2022) 106697.