# A Review Study: The Effect of Activator Type on the Compressive Basic Creep of Alkali-Activated Materials under Autogenous Condition

Lei Li, Naima Boumediene, Camille Follet, Boyuan Yang, Geoffrey Promis

► To cite this version:

Lei Li, Naima Boumediene, Camille Follet, Boyuan Yang, Geoffrey Promis. A Review Study: The Effect of Activator Type on the Compressive Basic Creep of Alkali-Activated Materials under Autogenous Condition. Results in engineering, 2025, pp.108831. 10.1016/j.rineng.2025.108831. hal-05447172

HAL Id: hal-05447172

https://u-picardie.hal.science/hal-05447172v1

Submitted on 10 Jan 2026

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

<figure><img src="image_1.png" /></figure>

Distributed under a Creative Commons CC BY 4.0 - Attribution - International License
Journal Pre-proof

A Review Study: The Effect of Activator Type on the Compressive
Basic Creep of Alkali-Activated Materials under Autogenous
Condition

Lei Li, Naima Boumediene, Camille Follet, Boyuan Yang, Geoffrey Promis

PII: S2590-1230(25)04874-1

DOI: https://doi.org/10.1016/j.rineng.2025.108831

Reference: RINENG 108831

To appear in: *Results in Engineering*

Received date: 17 October 2025

Revised date: 16 December 2025

Accepted date: 19 December 2025

Please cite this article as: Lei Li, Naima Boumediene, Camille Follet, Boyuan Yang, Geoffrey Promis, A Review Study: The Effect of Activator Type on the Compressive Basic Creep of Alkali-Activated Materials under Autogenous Condition, *Results in Engineering* (2025), doi: https://doi.org/10.1016/j.rineng.2025.108831

This is a PDF of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability. This version will undergo additional copyediting, typesetting and review before it is published in its final form. As such, this version is no longer the Accepted Manuscript, but it is not yet the definitive Version of Record; we are providing this early version to give early visibility of the article. Please note that Elsevier's sharing policy for the Published Journal Article applies to this version, see: https://www.elsevier.com/about/policies-and-standards/sharing#4-published-journal-article. Please also note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

© 2025 Published by Elsevier B.V.

This is an open access article under the CC BY-NC-ND license
(http://creativecommons.org/licenses/by-nc-nd/4.0/)

<figure><img src="image_2.png" /></figure>
1 Highlights

2 • The compressive basic creep at the macro scale is determined by the microphases of hydration products, anhydrous particles and pores.

3 • The burger model is available for the prediction of the compressive basic creep evolution.

4 • The creep mechanisms in AAMs could interpret more theories besides the water movement at early age and the C-S-H type gel rearrangement at long term.

7
# A Review Study: The Effect of Activator Type on the Compressive Basic Creep of Alkali-Activated Materials under Autogenous Condition

Lei Li¹,²*, Naima Boumediene³, Camille Follet¹, Boyuan Yang¹,², Geoffrey Promis³

1. Civil Engineering and Civil Mechanical Laboratory (LGCGM), INSA Rennes, University of Rennes, 20 avenue des Buttes de Coësmes, Cedex, 35007, Rennes, France

2. Civil Engineering Mechanics and Materials Laboratory (L2MGC), CY Cergy-Paris University, 5 mail Gay-Lussac, 95031 Neuville-sur-Oise, France

3. Innovative Technologies Laboratory (LTI), University of Picardie Jules Verne, Avenue des Facultés—Le Bailly, 80025 Amiens, France

* Correspondence: lei.li1@cyu.fr;

**Abstract:** To activate the hydration process and obtain the proper mechanical performance of mineral additions, like Blast-Furnace Slag (BFS), the chemical activation is a commonly used and friendly solution. As the nature of chemical activators affects the microstructure of materials, this paper provides a review of the activator effect on the compressive basic creep of alkali-activated materials under autogenous conditions. The impact of a partial replacement of BFS by Fly Ash (FA) is also reviewed. Alkali-activated BFS mortar and concrete generally show a larger creep than traditional Portland cement systems. The combined activation of sodium silicate and sodium hydroxide is reported to lead the higher mechanical performance and lower creep. It can be explained by its denser matrix characterized by a lower porosity and lower pore connectivity, possibly limiting the water movement, one of the basic creep driving mechanisms at early age. Another reason is related to the C-S-H type gels' mobility at long term. Among the sole use of activator types, the sodium silicate performs better, followed by the sodium hydroxide, sodium carbonate, and sodium sulfate. A small content of FA leads to a lower mechanical performance and higher creep magnitude at long term.

**Keywords:** alkali activator; blast-furnace slag; compressive basic creep; fly ash, microstructure

## 1 Introduction

For several decades, binders incorporating high volumes of mineral additions have been widely adopted as a strategy to partially replace conventional Portland Cement (PC), thereby mitigating its environmental impact. This approach primarily targets the reduction of clinker content, the main constituent of PC, whose production is highly energy-intensive and responsible for significant CO₂ emissions. By lowering the clinker factor, the overall carbon footprint of concrete can be substantially reduced [1–3]. One of the typical and commonly used mineral additions is ground granulated Blast-Furnace Slag (BFS) [1–6], an industrial by-product derived from the iron and steel manufacturing process, to produce the blended cements, such as CEM II and CEM III, used in Europe [4,7–9]. Its incorporation not only contributes to sustainability but also enhances the durability performance of concrete, including improved resistance to chemical attack and reduced permeability [4-6]. Another widely used supplementary cementitious material is Fly Ash (FA), a fine powder obtained from coal combustion in thermal power plants. In particular, the low-calcium variety-classified as Class F-FA (F-FA), is frequently utilized due to its high content of reactive aluminosilicates, with SiO₂+Al₂O₃+Fe₂O₃ exceeding 70% as specified in EN 450-1 [10]. Both BFS and FA serve as primary aluminosilicate precursors in the development of Alkali-Activated Materials (AAMs), a class of alternative binders considered to be environmentally friendly substitutes for traditional PC. These binders rely on the chemical activation of industrial by-products through alkaline solutions to form hardened matrices, offering a promising route toward sustainable construction materials with a significantly reduced carbon footprint.

To achieve appropriate hydration kinetics and desirable mechanical performance, both BFS and FA can be activated through various techniques aimed at enhancing their reactivity. These include mechanical grinding [11–13], thermal treatment [14–17], and, most notably, chemical activation [3,18–22]. Among these, chemical activation has received the most
attention, particularly the use of alkaline activators, such as sodium silicate ($Na_2SiO_3$), sodium hydroxide (NaOH), sodium carbonate ($Na_2CO_3$), and sodium sulfate ($Na_2SO_4$) [1,21,23-26]. Alkali silicates and hydroxides are known to significantly accelerate the BFS hydration kinetics [27-29], sometimes surpassing that of PC. However, these activators also present some limitations, such as a rapid setting times [4,10,27,30], a significant cost [20,23]. Moreover, they are corrosive and unhealthy [1,6]. In contrast, alkali carbonates and sulfates are more environmentally friendly [1,5,6], less expensive [20], and non-corrosive [31-33]. Thanks to these advantages, alkali sulfate-activated BFS materials have recently attracted significant attention [22,24,34-39]. However, as the nature of activators affects the C-(N)-A-S-H gel formation, the choice of activators still focuses on their impacts on hydration products, hydration kinetics, and mechanical properties [19,40,41]. The study of compressive creep behaviour in AAMs remains limited, particularly on the creep mechanisms.

Understanding these delayed deformations is essential for accurately assessing the durability of AAMs, particularly the cracking risk when their delayed deformations are restrained by internal or external elements. This risk is influenced by several material properties, including Young's modulus [8,42-45], tensile strength [8,44-46], shrinkage [42,44,47], and creep [8,42,44,47,48].

The present study consolidates and critically analyzes the current body of literature to investigate the influence of the activator type on the basic compressive creep behavior of AAMs. This analysis is conducted in relation to both microstructural features and macroscopic mechanical properties, with additional consideration given to the effects of partial FA substitution. Proposed creep mechanisms reported in the literature are examined, and existing research gaps are identified to guide future studies in this emerging field.

## 2 Remarks on testing methods

In previous studies [8,42,47-55], the compressive creep of mortars and concrete is most commonly measured using a creep rig apparatus. This method uses a manually adjustable loading frame that allows long-term testing under controlled conditions for most of the creep tests. Specimens are placed vertically (Fig. 1), and precise centering is essential to ensure uniform stress distribution and minimize eccentric loading errors.

The applied load is typically maintained by a hydraulic or gravimetric jack (i.e., Fig. 1a manufactured by 3r-labo.fr) [52], corresponding to 10-40% of the specimen's maximum compressive strength at the loading age (see Appendix A). This range provides a stable creep response, preventing crack damage and non-linear deformation. Deformation measurement is conducted using two stainless steel balls, attached along the specimen's axis with a fixed gauge length of 20 cm (Fig. 1a) [52,55]. Alternatively, strain gauges are bonded directly to the specimen surface (Fig. 1b) [50,54]. The test duration begins immediately after loading, and data are recorded manually or automatically at preset intervals. This setup allows for simultaneous testing of multiple specimens under identical curing conditions, and several advantages, including low cost and flexibility for testing different sizes and loading levels.

To monitor the creep evolution immediately after mixing, the Temperature Stress Testing Machine (TSTM) is also used in several studies, to evaluate the thermomechanical deformation behaviour of cementitious materials at a very early age [47,56,57] (Fig. 2). This apparatus allows simultaneous measurement of the internal temperature, delayed deformation, and stress evolution during the early hydration, making it suitable for studying early-age creep and stress relaxations [57]. However, this TSTM is expensive and complicated, leading to variable displacement readings during the early setting phase.

<figure><img src="image_3.png" /><figcaption>(a)</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>(b)</figcaption></figure>

90

91
92

<figure><img src="image_5.png" /><figcaption>Figure 1 The picture of compressive creep test for the dimension - diameter × height: 15×30 cm [52] (a) and 10×30 cm [45,50,54] (b)</figcaption></figure>

93

94

<caption>Figure 2 A scheme of TSTM (a), the disposition system with insulation hidden layers (b), and its picture (c) [56]</caption>

Despite the widespread use of these methods, no specific standardized protocol has been established for compressive creep testing. This lack of standardization often leads to differences in the specimens' shape [45,50,58] and dimensions [50,55]. However, a larger size (diameter×height: 15×32 cm) seems likely to be useful for the alkali sulfate-activated materials with a slower hydration kinetics at early age [24,52,59]. The strain gauge measurement is more applicable for the sodium silicate-activated materials with a fast early-age hydration kinetics [46,47,54], on the prismatic specimens (7×7×28 cm). To define the basic creep behaviour in the Alkali-Activated Materials (AAMs) under autogenous condition, a hypothesis used for the traditional Portland cement has been proposed [42]: the creep parameter (J-specific creep) is calculated using Equation 1, and the basic creep deformation is shown in Figure 3. This approach assumes that the total deformation measured can be superposed by the elastic, autogenous shrinkage, and basic creep deformations. Two stages are observed for the classic evolution of compressive basic creep:

* Loading stage (from $\tau_c$ to $\tau_d$): the autogenous shrinkage and basic creep deformations progressively increase over time, due to the internal hydration of materials and the microstructural development of matrix [42,46,47]. The mass monitoring is carefully controlled for the same batch specimen for the autogenous shrinkage and basic creep tests, making sure no effect of drying under the sealed condition.

* Unloading stage (after $\tau_d$): This stage is related to the relaxation behaviour of materials, characterized by the elastic modulus after unloading and the residual permanent deformation. Two processes (an initial recovery and a delayed creep recovery, respectively) are presented in Fig.3, which is not further discussed in this study.

$$J(t)(\mu\text{m/MPa}) = \frac{\varepsilon_{cp}(t)}{\sigma} = \frac{\varepsilon_{tl}(t) - \varepsilon_{f}(t) - \varepsilon_{i}(\tau_{c})}{\sigma} \quad \#(1)$$

Where $\varepsilon_{cp}(t)$ is the basic creep deformation $\sigma$ (MPa) the applied stress, $\varepsilon_{i}(t)$ the total deformation and $\varepsilon_{f}(t)$ the free deformation including the thermal and autogenous deformation, $\varepsilon_{i}(t)$ is the elastic deformation at the loading age $\tau_c$ (days).

<figure><img src="image_6.png" /></figure>
<caption>**Figure 3** The evolution of compressive creep deformation at the loading age $\tau_c$ and unloading age $\tau_d$, adapted from [42]</caption>

The compressive basic creep test is conducted under the autogenous condition, using a double-layer aluminium sheet sealed at 20±1 °C, preventing any moisture exchange with the external environment. Under this condition, the measured deformation indicates $\epsilon_{il}$, related to the pure basic creep ($\epsilon_{cp}$), as the constant temperature and relative humidity eliminate the effects of thermal and drying variations. Therefore, the basic creep deformation is primarily governed by two internal factors of materials: (1) the evolution of mechanical properties (e.g., elastic modulus and strength development) [42,55]; (2) the advancement of the hydration process [8,42,52,60].

### 3 Microstructure

The effect of activator type and the partial substitution of fly ash on the microstructure is characterized in terms of hydration products, anhydrous particles and porosity.

#### 3.1 Hydration products

Hydration products in Alkali-Activated Materials (AAMs) are determined by the distinct activation process, differing from the hydration reaction in Portland Cement (PC). Their composition and evolution have been extensively studied using several techniques [3,12,61,62]: nanoindentation, Thermogravimetric Analysis (TGA), X-Ray Diffraction (XRD), Nuclear Magnetic Resonance (NMR), Fourier-Transform Infrared spectroscopy (FTIR), and Scanning Electron Microscopy (SEM)..

The type of hydration products as a function of activator type is characterized in Table 1. Previous studies [30,40,63–67] have correlated activator composition and critical cationic ratios (Ca/(Si+Al) and Al/Si) with the formation of C-S-H type gel, analogous to hydrated PC, and C-A-S-H gel in AAMs. These gels exhibit a similar thermal decomposition pattern between 110°C and 250°C in TGA, a higher aluminium and alkali content (Na, K) with a consistent Ca/Si ratio of 0.9-1.2 [59,66]. XRD, $^{27}$Al NMR, FTIR and SEM presented an amorphous phase at the peak 2θ=29-32°[40,59,64–66], a chemical shift of 5–10 ppm $^{27}$Al spectra [61,68,69], Si-O-Al stretching bands at ~950 cm$^{-1}$ [61,68] and dense needle-like morphologies [59,68], respectively, corresponding to the presence of C-(N)-A-S-H gel. Based on the XRD and TGA analysis, C-(N)-A-S-H type gels are the main hydration product in BFS-based binders [24,59]. Filippis et al. [64] found that the content of C-A-S-H was approximately twice that of secondary hydrotalcite phases in NaOH-activated BFS binders with activator-to-binder mass ratios of 0.1–0.15 (Fig.4). Similar results were reported for Na$_2$SO$_4$- and Na$_2$SiO$_3$-activated systems [24]. The combined Na$_2$SiO$_3$-NaOH activation leads to a constant increase in C-A-S-H type gel formation, comparable to the C-S-H content observed in the hydrated PC system [30].

The creep of traditional Portland cement is dominated by the capillary porosity phase rather than by the inter-particle spacing and sliding of C-S-H nanoparticles under a sustained load (more details are discussed in Section 5.4.2). In AAMs, the two main binding gels differ structurally: C-A-S-H gels have a more compliant and stiffer structure than N-A-S-H. This is consistent with nano-indentation results (Table 1) [67,70,71], showing that C-A-S-H presents a higher elastic modulus (26-35 GPa), resulting in less viscoelastic deformation than N-A-S-H. However, these differences alone do not fully explain the macroscopic creep behavior. Several additional factors must be considered in AAMs systems, including the gel transformations with the coexistence of C-A-S-H and N-A-S-H [62], the stress redistribution onto the unreacted or partially reacted phases [70,72], the role of capillary porosity [62] and the influence of other potential reaction products.

Systems with high magnesium concentrations facilitate the formation of M-S-H type gels and Layered Double Hydroxides (LDH) with the partial substitution of Mg$^{2+}$ by Al$^{3+}$ [73,74]. These phases exhibit variable Mg/Al ratios. M-S-H gels show a high elastic modulus (111 GPa) compared to hydrotalcite-like phases (Table 1). Filippis et al. [64] reported the progressive formation of hydrotalcite with curing age, correlating with the improved long-term strength in NaOH-activated BFS mortars: 55 GPa at 90 days. In addition, the presence of sulfate ions in the pore solution further promotes ettringite and potential monosulfoaluminate formation, particularly in BFS system activated with sulfate activators [24,35,75].

The partial substitution of BFS with FA induces the formation of N-(C)-A-S-H gels through partial replacement of Ca$^{2+}$ by Na$^{+}$ or K$^{+}$ [70], characterized by reduced elastic modulus and hardness in FA or metakaolin-based systems activated by NaOH-Na$_2$SiO$_3$ (Table 1). These geopolymer gels, formed through polymerization processes (detailed in section 5.4.3), coexist with C-(N)-A-S-H phases in the hydrated matrix, characterized by the XRD peak 2θ=33° [61,68], and stretching
FTIR bands H-OH around 3500 cm⁻¹ [61,68]. Notably, systems activated solely with NaOH exhibit higher proportions of

N-(C)-A-S-H gels than C-(N)-A-S-H gels in BFS/FA blended systems activated with NaOH-Na₂SiO₃ [76].

<figure><img src="image_7.png" /></figure>
<table><caption>**Table 1** Hydration products of Alkali-Activated Materials (AAMs)</caption><thead><tr><th rowspan="2">Hydration products</th><th rowspan="2">PC systems</th><th colspan="2">AAMs systems</th><th rowspan="2">Structure</th><th colspan="2">Micro-properties</th><th rowspan="2">Capabilities</th><th rowspan="2">Reference</th></tr><tr><th>Precursors</th><th>Activator</th><th>Modulus (GPa)</th><th>Hardness (GPa)</th></tr></thead><tbody><tr><td>C-S-H</td><td>C<sub>3</sub>S/C<sub>2</sub>S</td><td>-</td><td>-</td><td>Tobermorite, Jennite</td><td>10-36</td><td>0.6-1.3</td><td>Stronger stiffness</td><td>[77]</td></tr><tr><td>C-A-S-H</td><td>-</td><td>BFS</td><td>NaOH, Na<sub>2</sub>SiO<sub>3</sub>, Na<sub>2</sub>CO<sub>3</sub>, Na<sub>2</sub>SO<sub>4</sub>,</td><td>Tobermorite like</td><td>26-35</td><td>-</td><td>Pore-filling capability</td><td>[67,71,78]</td></tr><tr><td>C-S-H type-gel</td><td>-</td><td>FA, Metakaolin</td><td>NaOH+ Na<sub>2</sub>SiO<sub>3</sub>,</td><td>Highly crosslinked</td><td>13-16</td><td>0.6-0.7</td><td>Excellent chemical resistance</td><td>[3,39]</td></tr><tr><td>N-A-S-H</td><td>-</td><td>BFS</td><td>MgO, (Na<sub>2</sub>SO<sub>4</sub>)</td><td>Talc-like</td><td>111</td><td>-</td><td>Lower energy demand</td><td>[73,74,79]</td></tr><tr><td>M-S-H</td><td>MgSO<sub>4</sub> attack</td><td>BFS</td><td>Na<sub>2</sub>SiO<sub>3</sub>,</td><td>Talc-like</td><td>-</td><td>-</td><td>-</td><td>[72]</td></tr><tr><td>C-A-H</td><td>-</td><td>BFS</td><td>MgO+NaOH, Na<sub>2</sub>CO<sub>3</sub></td><td>Staked layer network</td><td>28-38</td><td>0.07-0.12</td><td>Carbonation resistance, capillary pore formation</td><td>[64,80]</td></tr><tr><td>Hydrotalcite</td><td>Magnesia PC</td><td>BFS</td><td>Na<sub>2</sub>SO<sub>4</sub></td><td>Needle-like</td><td>15-25</td><td>0.5-0.6</td><td>A high-water content Pore-filling capability 1.8 g/cm<sup>3</sup></td><td>[81,82]</td></tr><tr><td>AFt phases</td><td>SO<sub>4</sub><sup>2-</sup> in PC</td><td>BFS</td><td>Na<sub>2</sub>SO<sub>4</sub>, Na<sub>2</sub>CO<sub>3</sub></td><td>Lamellar double hydroxide</td><td>-</td><td>-</td><td>-</td><td>[59]</td></tr></tbody></table>

<figure><img src="image_8.png" /><figcaption>**Figure 4** The evolution of mass loss content assigned to hydrated phases in NaOH-activated BFS binders, activator to binder mass ratio (a) 0.1, and (b) 0.15 [64]</figcaption></figure>
## 3.2 Anhydrous particles

The anhydrous particles in AAMs are mainly the anhydrous BFS and FA. A similar elastic modulus is observed for these two typical precursors, as the nanoindentation analysis conducted on the anhydrous areas indicates a similar value (90~100 GPa) for BFS and FA (Table 2) [75,80,81], which is quite close to the modulus of anhydrous particles in PC [83]. Despite quite similar micromechanical properties of anhydrous particles in AAMs and PC, two main characteristics are more important for BFS and FA anhydrous particles in the microstructural analysis: morphology and the volume fraction.

<table><caption>**Table 2** The micromechanical modulus of anhydrous particles</caption><thead><tr><th>Single phase</th><th>Anhydrous PC [67,71,83]</th><th>Anhydrous BFS [67,71]</th><th>Ca-rich anhydrous FA [70]</th></tr></thead><tbody><tr><td>Indentation modulus (GPa)</td><td>85±20</td><td>96±6</td><td>90±10</td></tr></tbody></table>

According to the SEM-BSE technique, the anhydrous BFS particles are characterized by a white irregular shape surrounded by a hydrated rim with a thin layer inferior to 1 µm (Fig.5). Anhydrous FA particles show small spheres (>2 µm), with a part of the spheres containing black hollows (Fig. 5). Their spherical shape provides more cohesion between particles, and they are less reactive, possibly due to the crystalline quartz phases. The morphology of hydrated gel is identified with a foil-like phase, observed for C-A-S-H type gels of alkali-activated BFS [59,84,85] and hydrated CEM III type cement at 28 days [7,9]. Until now, it is still challenging to identify the other hydration products in the hydrated gels, such as ettringite, hydrotalcite and portlandite, the last resulting from the PC hydration.

As hydration progresses over time in AAM binders, the volume of hydrated gel gradually increases. However, the volume of unreacted (anhydrous) phases remains significantly larger than that of the hydrated phases. In alkali sulfate-activated BFS binders, the anhydrous phase accounts for approximately 50-55% more than the combined volume of pores and hydrated products [59]. This suggests that the volume fraction of unreacted particles in various activated systems is likely influenced by the type of activator, which affects both the dissolution rate of BFS or FA particles and the precipitation rate of hydration products, particularly C-A-S-H type gels.

The higher volume fraction of anhydrous BFS phases has a stronger effect on the microscale creep behavior than the hydrated phases. These unreacted particles are stiffer and less deformable (Table 1&2), which limits the movement of the surrounding gel and affects local stress distribution. As a result, the overall creep response is mainly controlled by the interaction between the anhydrous and hydrated phases.

Partially replacing BFS with FA generally leads to a reduction in the amount of unhydrated BFS, resulting in a matrix with lower hydrated gel volume fractions. In sodium silicate-activated BFS systems [86], increasing the FA content (from 30% to 100% of the total binder mass) results in a notable increase in the volume fraction of pore phases (as shown by the black areas - in Fig.5), Additionally, the spherical morphology of FA particles at the microscale improves the workability of the fresh mix and promotes better dispersion of unreacted particles [2], which may contribute to enhanced pore connectivity over the long term.

<figure><img src="image_9.png" /><figcaption>**Figure 5** Microstructural features of Alkali-Activated Fly ash (AAF), Alkali-Activated Slag (AAS) and Alkali-Activated blended Fly ash and Slag (AAFS) [86]</figcaption></figure>

205

206

207
208

## 3.3 Porosity at 28 days

209

210

211

212

213

214

215

216

217

218

219

220

221

222

223

224

225

226

227

228

229

230

The pore structure is typically assessed using various techniques, such as oven drying, mercury intrusion porosimetry, nitrogen adsorption, water vapour sorption, and image analysis [87–92], with a primary focus on total porosity and pore size distribution. A denser pore network is generally associated with lower overall porosity and smaller pores, often filled with hydration products, which possibly decreases more volume fractions of unreacted phases, improving the deformation at the macroscale.

Previous studies [89–92] have shown that alkali-activated BFS systems using a combination of Na₂SiO₃ and NaOH exhibit lower porosity compared to those activated with single-component activators like NaOH, Na₂CO₃, or Na₂SO₄. In particular, vacuum drying at 45 °C applied to Na₂SO₄-activated mortars revealed a total porosity of approximately 17% at 28 days [52], due to a slow hydration process after mixing, and a lower formation of C-A-S-H type gels. A higher permeability was presented using the Klingenberg coefficient (<0.1, indicating large pores) with Na₂SO₄-activated systems [52]. As hydration continues beyond 28 days, the gel formation further helps fill the pore space, reducing the porosity and pore size, forming a denser microstructure. The denser microstructure is indeed characterized by a lower sorptivity value (3.56) with Na₂SiO₃-NaOH activation [93].

The incorporation of FA as a partial replacement for BFS generally increases porosity. For example, Na₂SiO₃-NaOH activated mortars show a higher proportion of pores in the 10-10⁴ nm range, rising from 13% to 20% at 28 days when 40% of the binder mass is substituted with FA [91]. Under comparable activator content (4% equivalent Na₂O), a 20% FA replacement in BFS mortars leads to an approximate 3% increase in total porosity when using Na₂SO₄ as the activator [52]. Overall, the higher porosity observed with partial FA replacement leads to a coarser matrix characterised by a lower content of hydration products, regardless of activator type.

231

## 4 Macroscopic properties under autogenous condition

232

233

234

235

236

The effect of activator type and partial fly ash substitution is reviewed on the three mechanical parameters under the autogenous condition: mechanical strength, Young's modulus and autogenous shrinkage. These properties strongly determine the evolution of compressive basic creep. The compressive strength is linked to the applied loading level, Young's modulus determines the initial elastic deformation, and autogenous shrinkage is used to subtract the basic creep deformation.

237

### 4.1 Mechanical strength

238

239

240

241

242

243

244

245

246

247

248

249

250

251

252

253

254

255

256

257

258

259

Several studies have reported that the Water-to-Binder ratio (W/B), the activator content (as equivalent Na₂O content of the total binder mass), and the silicate modulus affect the compressive strength of the alkali-activated system. The optimal ranges [28,94] are typically 0.4-0.45 W/B, the 4-6% equivalent Na₂O content, and the 0.8-1.6 silicate modulus. Mohamed et al. [61] indicated that the combined use of Na₂SiO₃ and NaOH produces better strength at 28 days under autogenous conditions, compared to Na₂CO₃ and NaOH. This is a main advantage of using Na₂SiO₃ with NaOH as reported in the literature [1,61,68,95]. Furthermore, the single use of Na₂SiO₃ shows superior performance in optimized paste, mortar, and concrete compositions under sealed curing conditions, where the effects of carbonation and temperature are eliminated, followed by the NaOH, Na₂CO₃ and Na₂SO₄. This rank is also consistent with the loading strength results used for the creep tests in Appendix A. For instance, the compressive strength of Na₂SiO₃-activated mortars presents a superior strength of 50-70 MPa at 28 days [45,52,54], and a sorptivity below 0.05 mm·√min⁻¹. Its activation promotes proton-balanced, cross-linked silicate chains, evidenced by ²⁹Si NMR Q³(1Al) signals at approximately -70 ppm, resulting in reduced ionic ingress [96]. The high concentration of silicate ions promotes the denser formation of C-(N)-A-S-H type gels [97]. These gels enhance mechanical strength by reducing pore volume and strengthening the chemical bonds between hydration products.

NaOH-activated systems (at the silica modulus equal to 8 molarity) show a lower strength (about 47 MPa) [11]. The Na₂CO₃/Na₂SO₄ activation system leads to a less than 28 days compressive strength around 20-30 MPa [34,38,75], due to the other possible product formations, such as hydrotalcite and ettringite (as they are presented in Table 1). So their microstructure [52] is quite coarser, due to the excitation of carbonate and sulfate, compared to that of Na₂SiO₃ or NaOH activated systems. The lower compressive strength decreases the stiffness of materials to promote more deformation under a constant load.
260

261

262

263

264

265

266

267

268

269

270

271

272

273

274

275

276

277

278

279

280

281

282

283

284

285

286

287

288

289

290

291

292

293

294

295

296

297

298

299

300

301

302

303

304

305

306

307

308

309

310

311

312

The partial substitution of BFS with FA in alkali-activated systems reduces the amount of available anhydrous BFS, thereby limiting the dissolution of reactive particles and slowing down the hydration kinetics at early age. The activator effect is also adaptable for BFS/FA blends. The Na₂SiO₃ leads to better 28-day compressive strengths (50-70 MPa), compared to 47 MPa with NaOH activator in BFS/FA binders (the activator content is defined by the equivalent Na₂O content: 5-7%) [61,93,98].

Additionally, the change in the pore solution alkalinity, depending on the type of activator used [35,40], may further influence the precipitation rate of hydration products. This slower rate could be associated with the formation of alternative gel phases, such as those arising from geopolymerisation processes [99]. The coexistence of N-A-S-H and C-A-S-H gels may hinder further dissolution of aluminosilicate sources [2,99], resulting in increased pore volume and overall porosity, as illustrated in Fig.5.

## 4.2 Young's modulus

The elastic modulus is a critical mechanical property of cementitious materials, reflecting their stiffness and load-bearing capacity. Unlike purely chemical analyses, the elastic modulus is a volumetric integral quantity that captures the interactions between gel particles, their organization, spatial arrangement, porosity and bonding mechanisms. A higher modulus leads to more stiffness to resist deformation. In alkali-activated BFS concretes, the elastic modulus is largely determined by the formation of distinct reaction products (C-A-S-H type gels) [87], and the Interfacial Transition Zones (ITZ). The use of a combination NaOH+Na₂SiO₃ leads to a higher modulus (35-50 GPa) [100–102], due to a significant formation of C-A-S-H type gels and their densely packed calcium-rich layers. The porosity of ITZ also decreases, due to a higher stiffness and strong interparticle bonding with the presence of aggregate. However, the sole use of NaOH activator leads to a 43% reduction of modulus (17.1 GPa vs. 12 GPa) at a ratio W/B of 0.55 [103]. The rank of modulus as a function of activator is likely similar to the rank of compressive strength, as a larger number of AAMs studies [101,104–106] have indicated an almost linear relationship between the elastic modulus (20-40 GPa) and compressive strength (30-60 MPa). It is related to the effect of the activator on the hydration degree and the matrix characterized by the content of hydration products and porosity.

The partial replacement of BFS with FA leads to a reduction in Young's modulus at 28 days, primarily due to two factors: (1) The decrease in BFS content reduces its overall contribution to stiffness. For instance, a low-density BFS (63 kg/m³) can exhibit a higher modulus (21 GPa) compared to high-density FA (344 kg/m³), which shows a lower modulus of approximately 15 GPa under a water-to-binder ratio of 0.42 and 4% activator content [105]. (2) The formation of N-(C)-A-S-H-type gels in FA-rich systems contributes to lower stiffness. N-A-S-H gels are characterized by a more disordered and porous structure, resulting in a lower elastic modulus (typically 10–20 GPa - Table 1).

To achieve a higher elastic modulus that meets ACI and Eurocode requirements, an optimized mix design is essential. This includes a higher BFS content, a moderate activator modulus ranging between 0.7 and 1.5 [106], and a reduced water-to-binder ratio of around 0.4. Under these optimized conditions, elastic moduli as high as 34.2 GPa have been reported, comparable to those of conventional concrete, demonstrating the potential of such alkali-activated systems as sustainable alternatives for structural applications.

## 4.3 Autogenous shrinkage

The autogenous shrinkage behavior of alkali-activated materials (AAMs) has been widely investigated [26,107], with several studies examining the influence of key parameters on shrinkage magnitude. These factors include the type of precursor, the nature and concentration of the alkaline activator (expressed in terms of molarity or equivalent Na₂O content), the aggregate content, and the W/B ratio. A comprehensive review focuses primarily on systems activated with sodium silicate. It has been consistently observed that sodium silicate-activated pastes exhibit significantly higher autogenous shrinkage values -ranging from 3000 to 8500 µm/m - when prepared with an optimal activator modulus (SiO₂/Na₂O = 1) and a w/b ratio of 0.4. These values are substantially greater than those observed for other activator systems. However, when aggregates are introduced (i.e., in mortars or concretes), autogenous shrinkage is notably reduced due to the restraining effect of the aggregates - an effect that is comparable in both AAMs and PC systems [26]. In contrast, pastes and mortars activated with sodium carbonate or sodium hydroxide generally show lower autogenous shrinkage than those activated with sodium silicate. This reduction is attributed to the lower alkalinity (i.e., lower pH) of these systems [24], which slows the dissolution rate of anhydrous BFS particles and limits the availability of silicate species in the pore solu-
tion. As a result, the formation of hydration products is reduced, thereby minimizing shrinkage. Among the different activators studied, sodium sulfate appears to induce the lowest autogenous shrinkage in both paste and mortar systems [24,59], making it a promising alternative for applications where shrinkage control is critical. The autogenous shrinkage of alkali sulfate-activated BFS is inferior to the activation of Na₂SiO₃ and NaOH [26], but quite similar to that of PC [22,26]. The use of Na₂SO₄ activator leads to a larger autogenous shrinkage than that of K₂SO₄ activator (Fig.6b). It is related to their higher content of hydrates (e.g., ettringite & C-A-S-H), leading to a denser matrix and a faster self-desiccation process [52]. Mohamed et al. [61] have indicated that the formation of ettringite in AAMs leads to a minor early-age shrinkage (<0.1%), followed by negligible long-term degradation, compared to the PC systems.

The capillary pressure mechanism [26] is the principal theory to explain the autogenous shrinkage of AAMs as a function of activator type. That seems to be also available for the activation of BFS/FA. The partial substitution of BFS by FA mitigates the autogenous shrinkage of Na₂SiO₃-activated pastes and Na₂SO₄-activated mortars (Fig. 6), related to higher contents of anhydrous particles and possible formation of N-(C)-A-S-H type gels. These geopolymer gels observed in the previous studies [26,99], coexist with C-(N)-A-S-H gels in the hydrated matrix, leading to a modified pore structure and a slowing-down evolution of autogenous shrinkage.

<figure><img src="image_10.png" /><figcaption>**Figure 6** Autogenous shrinkage of (a) alkali silicate and hydroxide activated BFS paste [26] and (b) alkali sulfate activated BFS mortars [52]</figcaption></figure>

328
329
330
331

## 5 Basic creep

332

The influence of activator type and a small FA content on the compressive basic creep evolution is discussed for BFS-activated materials, as well as the driving forces for compressive creep.

333
334

335

### 5.1 Blast-furnace slag

336
337
338
339
340
341
342
343
344
345
346
347
348
349

Basic creep is a critical parameter in evaluating the long-term durability of cementitious materials, as it is influenced by a range of interrelated factors. These influencing factors are summarized in Table A (Appendix), based on findings from multiple studies. In both AAMs and PC systems, a similar trend is observed: a higher aggregate volume fraction, a lower water-to-binder (w/b) ratio, and a later loading age all contribute to reduced creep deformation, primarily due to the increased stiffness and rigidity of the matrix [50,108]. Figure 7 presents a compilation of available experimental data on the evolution of specific creep as a function of activator type, with loading applied at 28 days, corresponding to the conditions outlined in Table A. Under comparable test parameters, including a loading level of 30–40% of maximum compressive strength, a w/b ratio between 0.42 and 0.5, and a limited activator dosage (≤10% of binder mass), BFS-based mortars and concretes activated with various chemical activators generally exhibit greater creep deformation than their PC counterparts [45,52,109,110]. Ma et al. [110] verified the higher creep evolution with the use of sodium silicate and sodium hydroxide (Na₂SiO₃+NaOH), at a 28-day loading age compared to PC concrete. However, the effect of Na₂SiO₃ on the creep evolution is less significant, compared to the use of other activators, possibly related to the enhanced refinement of C-(N)-A-S-H gel structure, facilitated by the available silicate species and forming sufficient formation products. The higher elastic modulus after loading and compressive strength also indicates its denser microstructure.

350

351

352

353

354

355

356

357

358

359

360

361

362

363

364

365

366

367

368

The effect of Na₂CO₃ on the compressive creep is more significant compared to Na₂SiO₃, as Fig.7a shows a final value of creep equal to 150µm/m/MPa at 300 days. It is possibly due to the coarser matrix, characterized by a lesser formation of C-(N)-A-S-H. Indeed, a lower compressive strength (38 MPa) and modulus (28.7GPa) at 28 days is lower than 40 MPa
and 30 GPa for Na₂SiO₃ activation [60], due to a more porous binder matrix. In the NaOH activation, the larger creep deformation is due to the strong action of alkaline ions on the surface tension within the porous network and the instability of C-S-H, due to the incorporation and adsorption of alkaline ions or long-term polymerization [58]. The use of Na₂SO₄ activator leads to a faster kinetics of specific creep evolution during the first week after loading (an average value equal to 10 µm/m/MPa at 29 days), but a quite similar value (20~30 µm/m/MP) at long term with the Na₂SiO₃-activated concrete [52]. For the sulfate activation, the Na⁺ in the sulfate activation is more favourable to form more contents of C-(N)-A-S-H than K⁺ [75], leading to a denser pore structure. The lower porosity (about 17%) was measured for Na₂SO₄-activated mortars using the water porosity tests in the study of Li [52]. It could also explain the higher modulus and lower permanent deformation (Appendix Table A). Additionally, the lower creep deformation with Na₂SO₄ is possibly related to the formation of space-filling ettringite (Table 1).

Moreover, it appears that the impact of two combined activators on the creep behaviour is more complicated, rather than the superposed effect of only one activator used. Indeed, the creep value with the use of Na₂CO₃ and Na₂SiO₃ activators is larger (50 µm/m/MPa) than the value with the sole Na₂CO₃ activator [54]. The refinement of this matrix seems to be hindered, and the authors contributed to the synergistic effects of silicate and carbonate ions, leading to the calcium carbonate precipitation and restricted availability of reactive silicate species. However, it requires more strong evidence to explain the effect of two combined activators, such as the experimental results on the pore structure, hydration kinetics and formation kinetics.

<figure><img src="image_11.png" /><figcaption>**Figure 7** Evolution of specific creep with alkali-activated BFS concrete (a) [54] and mortars (b) at 28-day loading age [52,58]</figcaption></figure>

## 5.2 Fly ash effect

The partial replacement of BFS with FA results in a reduction in creep development under autogenous curing conditions (100% RH, 23 ± 0.5 °C) [58]. This behavior is observed in the Na₂SiO₃-NaOH activation system, where loading began at 76 days (Fig.8a). However, a higher creep was reported for mixtures with 80% FA replacement during a 10-month loading period, possibly due to the incorrect loading age (76 days instead of the standard 28 days). The aging effect and curing history were not properly normalized or controlled in the experiment. Therefore, the results reported by Hojat *et al.* [58] should be re-evaluated to verify the true influence of FA in Na₂SiO₃-NaOH-activated BFS mortars.

In addition, the influence of FA on creep also varies as a function of activator nature under the same experimental conditions [52] (Fig.8b). During the first week after loading, a faster specific creep evolution is observed in the K₂SO₄-activated system with the 20% FA replacement. Beyond, the evolution slows down with a lower value at 125 days (41 µm/m/MPa), compared to 52 µm/m/MPa for BFS mortars without using FA. In contrast, a large creep was recorded for the Na₂SO₄-activated BFS mortars containing FA. The partial replacement of BFS by FA reduces BFS content, compressive strength, and autogenous shrinkage, leading to a coarser microstructure. Although several studies [10,26] have reported that the incorporation of FA into BFS does not alter the nature of C-A-S-H gels, the observed difference in creep behaviour suggests that additional microstructural mechanisms may be involved. This aspect will be further discussed in Section 5.4.3.
<figure><img src="image_12.png" /><figcaption>Figure 8 Specific of the evolution of (a) Na₂SiO₃-NaOH-activated BFS mortars loaded at 40% of the 28-day compressive strength [58] and (b) sulfate-activated BFS mortars at 30% of the 28-day compressive strength [52]</figcaption></figure>

## 5.3 Numerical models

The compressive basic creep deformation is a viscoelastic response comprising instantaneous elastic ($\epsilon_i$) and basic creep components ($\epsilon_{cp}$) (Fig.3). This type of deformation is measured under sealed conditions, thereby excluding the influence of external moisture exchange. The modelling work is performed using the Burger model, which combines the Kelvin-Voigt and Maxwell model chains [42,45,47,77,111]. The corresponding rheological representation is shown in Figure 9. Rifai et al. [45] predicted the evolution of basic creep at 3, 7, and 28 days of loading for NaOH-activated mortars using a single Kelvin-Voigt chain. But this simplified model shows limitations in accurately reproducing long-term creep behavior (after 10 days of their loading ages), compared to the Burger model, including Kelvin-Voigt-Maxwell elements. This observation has been further verified in Na₂SiO₃-NaOH-activated geopolymers [55] and Na₂SO₄-activated BFS mortars [52], where the burger model provides a better fit to the experimental creep data, using the Isqnonlin function to solve the nonlinear least-squares problem. The Kelvin-Voigt model is described with Equations (5)-(7), while Equations (4) and (8) characterize the Maxwell model [42,45,47,111]. The equations (4) and (5) are conditioned at 20°C, where the ambient temperature (t) is initialized to the equivalent time ($t_0$) and the initial setting time (s) [55].

<figure><img src="image_13.png" /><figcaption>Figure 9 Viscoelastic behaviour of the hydraulic binder via the Burger model [55]</figcaption></figure>

$$ \epsilon_{bc} = \epsilon_{KV} + \epsilon_M^{\#(2)} $$

$$ \dot{\sigma} = \left( E_{KV}(t) + \frac{\eta_{KV}(t) \dot{\epsilon}_{KV}}{E_{KV}} \right) \dot{\epsilon}_{KV} + \eta_{KV}(t) \dot{\epsilon}_{KV}^{\#(3)} $$

$$ \sigma = \eta_M(t) \dot{\epsilon}_M^{\#(4)} $$

$$ E_{KV}(t) = E_{KV0} \exp \left( s \left( 1 - \sqrt{\frac{28 - t_0}{t - t_0}} \right) \right)^{\eta_{EKV}} \quad \#(5) $$

$$ \eta_{KV}(t) = \eta_{KV0} \exp \left( s \left( 1 - \sqrt{\frac{28 - t_0}{t - t_0}} \right) \right)^{\eta_{EKV}} \quad \#(6) $$
$$ \eta_{KVt} = E_{KV0} \times t_c \#(7) $$

$$ \eta_M(t) = \eta_{M0} \frac{t}{t_0} \#(8) $$

Where $\varepsilon_{KV}$ and $\varepsilon_M$ are the creep strains linked to the Kelvin-Voigt and the Maxwell chain, respectively. $E_{KV}$ the apparent stiffness, $\eta_{KV}$ the apparent viscosity. $E_{KV0}$, $\eta_{E KV}$, $\eta_{KV0}$, $\eta_{M0}$ and $t_c$ (characteristic time) are the model parameters. $E_{KV0}$ and $\eta_{M0}$ both determine the creep amplitude at long term after loading at $t_0$. $E_{KV0}$ at the 28 days loading age in a traditional Portland mortar is 55 GPa, superior to 22 GPa with BFS mortars activated by NaOH. The $\eta_{E KV}$ and $\eta_{\eta KV}$ is more related to the evolution rate of creep deformation, defining the creep simulation at early age.

The Burger model is a mathematical approach based on macroscopic creep measurements and provides an effective means of predicting creep behavior as a function of activator type in AAMs. The combination of this model with artificial intelligence and neural networks offers a promising pathway to accelerate research and optimize the design of low-carbon AAMs.

## 5.4 Mechanisms

In the absence of thermal and drying effects, understanding where creep occurs and how it develops is crucial for mitigating and predicting creep evolution in different alkali-activation systems. In Portland Cement (PC)-based materials, the main creep mechanisms are associated with water movement at early ages and C-S-H network rearrangement at later ages. The possible additional mechanisms may exist, due to differences in gel chemistry and microstructure in Alkali-Activated Materials (AAMs).

### 5.4.1 At short term

The water movement theory is commonly employed to explain the short-term evolution of basic creep, particularly within the first 10 days following load application. Similar to hydrated PC systems, alkali-activated BFS materials exhibit a porous structure composed of a solid skeleton of anhydrous particles interspersed with a network of capillary and gel pores (Fig. 10). Upon the application of a sustained load ($\sigma_{sph}$), liquid water, primarily located within the pore structure, partially migrates from the hydration products into the surrounding pore spaces.

This internal water redistribution facilitates the transfer of the applied load to the solid skeleton of the material, resulting in time-dependent deformation at the macroscale. The effect is particularly pronounced in systems characterized by a coarser pore structure and higher initial water content, where capillary pores provide a more accessible pathway for water movement. In such cases, the applied stress promotes enhanced water flow toward the capillary network, which in turn contributes to microstructural rearrangement and progressive densification of the matrix over time.

<figure><img src="image_14.png" /><figcaption>Figure 10 The schematic representation of water movement [42]</figcaption></figure>

To better understand the creep mechanisms in AAMs, the measurable characteristics at the microscale are linked to the compressive basic creep at 28-day loading age 28 with the loading level at 30-40% of compressive strength, as presented in Table 3. A higher kinetic and magnitude of compressive specific creep is presented with NaOH, $Na_2CO_3$, and $K_2SO_4$. They are both characterized by a higher porosity and larger pore size. In contrast, the lower creep is for
Na₂SiO₃/Na₂SO₄-activated mortars, due to a lower porosity and smaller pore size [88,90], forming a denser microstructure. Their higher autogenous shrinkage evolution validates the faster water consumption [24] and the lower content of water limits the water movement, leading to the smaller creep evolution. So, the water movement mechanism is true for the AAMs activated by NaOH, Na₂SiO₃, Na₂CO₃ and Na₂SO₄. It seems to be adaptable for BFS/FA-based materials activated by sulfates in Fig. 8b [52].

<table><caption>**Table 3** Basic creep and measurable parameters at the microscale [24,45,52,110]</caption><thead><tr><th rowspan="2">Activators</th><th colspan="2">Basic creep</th><th colspan="3">Pore structure</th><th rowspan="2">C-S-H/C-A-S-H</th></tr><tr><th>Kinetics</th><th>Magnitude</th><th>Water content</th><th>Porosity</th><th>Pore size</th></tr></thead><tbody><tr><td>NaOH</td><td>+</td><td>+</td><td>-</td><td>+</td><td>+</td><td>+</td></tr><tr><td>Na₂SiO₃</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>+</td></tr><tr><td>Na₂CO₃</td><td>+</td><td>+</td><td>+</td><td>+</td><td>+</td><td>-</td></tr><tr><td>Na₂SO₄</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>+</td></tr><tr><td>K₂SO₄</td><td>+</td><td>+</td><td>+</td><td>+</td><td>+</td><td>-</td></tr></tbody></table>

Note: +: high value/fast kinetics, -: lower value/slow kinetics;

### 5.4.2 At long term

The long-term creep is related to the viscous phenomenon occurring in the C-S-H interlayers during loading, which is an irreversible part of creep for the traditional PC materials. This viscous process involves different ways under a load, leading to a reorganization of the C-S-H nanostructure, such as the sliding of C-S-H [112], the consolidation of the C-S-H [113], the breaking and remaking of C-S-H bonds [42] or microcracking [51]. However, the difficulty is to have an appropriate theory supported by experimental evidence, based on the literature. One of the main theories widely adopted is the micro-prestress theory, as proposed by P. Bazant *et al.* [113]. The viscosity depends on a tensile micro-prestress carried by the bonds and bridge crossing the micropores (gel pores-traverse bonds on Fig. 10) in the hardened cement gel. This micro prestress characterizes self-equilibrated stress and is initially produced by incompatible volume change in the microstructure during hydration. These stresses stretch and break the interatomic bonds of C-S-H layers, allowing the viscous shear slips of C-S-H sheets (Fig. 11). The new hydration products form between a reorganized capillary space. This theory is based on the hypothesis: C-S-H is the only phase responsible for the macroscopic creep without its pores, supported by short-term nanoindentation data [77]. The inverse analysis using the fast Fourier transform method [77] has revealed that the C-S-H presents a slower creep evolution rate than the creep strain of paste or concrete, where a partial load is transferred to the non-creeping phases, such as portlandite or unreacted clinker minerals. The aging and hydration effects have not been discussed.

Table 3 only presents a brief link between C-(N)-A-S-H content and creep magnitude in OH/SO₄²⁻/CO₃²⁻ based activated BFS mortars, but its mechanisms governing this behavior in AAMs remain insufficiently understood. The creep response and gel relaxation in C-A-S-H and N-A-S-H systems remain unclear (micro-elastic values in Table 1). The formation of hydrotalcite, ettringite, and AFm phases, due to carbonate and sulfate activation (Table 1), may introduce additional uncertainties in stress distribution among the different phases. Accurate microscale measurements are required to clarify these and the porosity phase, and the effects of aging and ongoing hydration in AAMs will also be carefully considered.

<figure><img src="image_15.png" /><figcaption>**Figure 11** The scheme of sliding of C-S-H gel sheets for explaining long-term basic creep with the disjoining pressure (Pd) and micro-prestress (S) [42]</figcaption></figure>
478

### 5.4.3 Potential mechanisms

479

480

The other mechanisms may exist for the AAMs systems. FA could lead to the geopolymerization process [114] with the formation of N-A-S-H gels, presenting the four stages shown in Fig.12:

481

482

* The 1st stage is related to the dissolution process of anhydrous FA particles characterized by a dimension greater than 5 µm. The dissolution of cations and anions resulted from the activators, could also occur.

483

484

* During the 2nd stage, the aluminosilicate gel is formed because the ionic monomers accumulate to form dimers, which then interact with another monomer to form trimers.

485

486

487

* The 3rd stage is related to two types of gels: Gel 1, with a higher content of aluminum, is due to a quicker dissolution of aluminum than Si-O bonds and Al-O bonds. Gel 2 is due to more dissolution of Si species with a higher Si concentration (Si/Al ratio is around 2).

488

489

490

491

* The 4th stage begins when the gels start to accumulate and form a larger molecular structural unit, named N-A-S-H gel, leading to a three-dimensional change of geopolymer networks. This process is also linked to the polymerization reaction happening during the hardening process. Within a hardened and stable binder structure, the positions of the gel particles are “fixed” by adjacent gel, hence no longer mobile.

<figure><img src="image_16.png" /></figure>

492

493

<caption>**Figure 12** A general model of geopolymerization mechanisms for the sodium silicate activated FA systems [114]</caption>

494

495

496

497

498

499

500

501

502

503

504

505

506

507

508

509

510

511

512

513

514

515

516

517

518

519

520

521

522

523

The process of geopolymerization has been an important theory to explain the autogenous shrinkage of AAMs systems. The dimension change for autogenous shrinkage is not only explained by the chemical reaction, like capillary pressure, but also from a gel or even molecular level, particularly for the alkali-activated FA [26]. Based on the current review [26,114] on the geopolymerization theory, it has been difficult to quantify the corresponding deformation. Zheng *et al.* [115] explained the higher autogenous shrinkage of alkali silicate-activated slag by an increased polymerization degree of Si between C-A-S-H gels, which is also observed with alkali silicate-activated FA and metakaolin-based geopolymers [116,117]. It is attributed to the formation of N-A-S-H gel phases (Figure 12, stage 4). The finer pore network slows down the micro-diffusion of pore solution, resulting in a lower creep rate in geopolymer concrete compared with Portland cement concrete. Sounean *et al.* [55] reported that a higher content of N-A-S-H gels reduces porosity and produces a more compact microstructure, thereby decreasing creep. In contrast, Chi Hou Un [118] observed the opposite trend, attributing larger creep to increased slippage within the amorphous N-A-S-H gels rather than the formation of more stable crystalline phases such as Mg-rich diopside. Therefore, the influence of FA on the basic creep mechanisms remains uncertain, as several competing mechanisms may simultaneously affect the deformation behavior.

Understanding the effect of FA on the creep behaviour of the BFS system is also useful to explain the FA impact on the PC blended systems. Several controversial results were found in the PC/FA blended binders. Bamforth *et al.* [119] demonstrated that the substitution of 30% of the binder with FA leads to lower creep values compared to those of PC based concrete. Zhao *et al.* [120] reported that an increase in the FA content from 20% to 40% replacing PC can reduce the amount of C-S-H type gels, increase the porosity and the specific creep of concrete loaded at 10 days, which is also consistent with the early-age creep results of Klausen *et al.* [121].

Apart from the delayed deformation, the basic creep of AAMs also seems to have the potential to decrease, due to the formation of crystalline reaction products, such as ettringite in the sulfate-activated materials [59]. The expansion associated with their formation can only compensate for the part of autogenous shrinkage [26], decreasing the total deformation value in a certain stage under the compressive load. Currently, no method is available to quantify this effect in AAMs. The effect of hydrotalcite formation on the delayed deformation is insignificant, due to the limited amount. The micro-crack theory has been proposed [51], based on the effect of loading type on the creep deformation. Micro-cacks may occur, when aligning parallel to the loading direction in compression. The autogenous shrinkage increases, leading to an increase in the total deformation and a higher microcrack density. This theory is possibly related to the water movement mechanism.
524

525

526

527

528

529

530

531

532

533

534

535

536

537

538

539

540

541

542

543

544

545

546

547

548

549

550

551

552

553

554

555

556

557

558

559

560

561

562

563

564

565

566

567

568

569

570

571

572

573

574

575

576

In summary, a widely accepted consensus in the literature is that water movement, driven by ongoing hydration and internal water consumption, is a key mechanism underlying the basic compressive creep of AAMs. Due to the high surface tension of the pore solution [26], the typically denser pore structure, and a high degree of saturation, water mobility in AAMs is more pronounced than in PC-based systems. This promotes greater deformation under sustained load, particularly in the short term. The increased short-term deformability of AAMs is also attributed to the pronounced viscoelastic behavior of C-(N)-A-S-H. However, over the long term, this viscoelasticity tends to reduce creep deformation as the structure stabilizes.

Despite the usefulness of the water movement theory and the gel viscoelasticity hypothesis, these models alone do not fully explain the basic compressive creep behavior observed in alkali-activated BFS systems. In Portland cement, capillary pressure becomes dominant only after the microstructure has sufficiently stabilized. In contrast, the chemical complexity of AAMs, particularly the early-stage reactions involved in geopolymerization, suggests that additional mechanisms may be at play. These include inter-particle forces within the evolving gel network shortly after loading, which may influence early-age creep behavior.

Such alternative explanations, especially those rooted in geopolymer chemistry, are promising for AAMs, given the distinct nature of their reaction products. However, in the absence of comprehensive experimental validation, these mechanisms remain largely hypothetical and require further investigation through advanced characterization and modelling.

## 6 Conclusions

In this paper, we present a systematic literature review on the compressive basic creep of Alkali-Activated Materials (AAMs), including the common testing methods, existing creep data, the current numerical modelling, and key factors affecting the creep mechanisms in different alkali activation systems. Based on this review, the following conclusions can be drawn:

* The experimental methods used for the traditional Portland mortar and concrete are also applicable to AAMs. The choice of specimen size should be carefully considered, as it is affected by the early-age hydration process and the development of mechanical properties. The applied load should be standardized as a fixed ratio of the compressive strength at the same loading age to ensure test consistency.

* At the macroscale, the compressive basic creep is shaped by the interplay among hydrated gels, unreacted BFS phases, and pore characteristics. The influence of activator nature on basic creep can be understood as an extension of its effects on mechanical strength, Young's modulus, and autogenous shrinkage under autogenous conditions.

* C-S-H nanostructure models from Portland cement do not adequately capture compressive creep behavior at the mortar or concrete scale. The Burger model is currently the preferred approach for predicting creep. More studies are required to validate creep mechanisms linked to geopolymerization and other gel structures.

Future research on the creep behaviour of AAMs is summarized as below:

* The experimental protocols for the creep behaviour of AAMs at a very early age need to be developed.

* The potential microstructural factors affecting compressive basic creep still require systematic verification. These include the influence of the relaxation behavior of hydrated gels under a sustained load, and the characteristics of the interfacial transition zone. Understanding these aspects would provide deeper insight into the creep mechanisms of AAMs and their long-term deformation behavior.

* Multi-scale creep models, linking nano-, micro-, and macro-level processes, need to be further refined by developing more comprehensive numerical methods. The thermal and drying effect on the compressive basic creep can first be explored by the numerical method.

* AI-driven models combined with existing AAM databases can be used to optimize material composition and loading conditions. These databases include creep results and microstructural parameters. Such integration offers a promising direction for improving the design and performance of AAMs. The low-carbon AAMs with designed microstructures, through fibers, nanoparticles, or slag/fly ash blends, will present better control of viscoelastic deformation in structural and/or high-temperature applications.

**Data availability:** The data of this study are available from the author, Dr. Lei LI, upon reasonable request.

**Acknowledgments:** The author acknowledges the financial support provided by the China Scholarship Council Grant (No. 201806160296), as well as the Laboratory of Civil Engineering and Mechanical Engineering (LGCGM).
## Appendix A

<table><caption>**Table A** Main parameters related to the compressive basic creep tests of alkali-activated Blast-Furnace Slag (BFS) systems loading at 28 days in the literature. The table is categorized based on the type of Alkaline Activator (AA). w/b is the water-to-binder ratio and Activator Content (AC). T (°C) the curing temperature and S<sub>sp</sub> (cm<sup>2</sup>/g) the specific surface of BFS, Ø the diameter. The content of the activator is defined relative to the weight of total binder mass, and the molarity (M) is linked to the silicate modulus.</caption><thead><tr><th colspan="2">Specimen</th><th colspan="5">Composition</th><th colspan="3">Main parameters and results from creep tests</th></tr><tr><th>System<sup>1</sup></th><th>Size<br>(cm)</th><th>T (°C)</th><th>S<sub>sp</sub><br>(cm<sup>2</sup>/g)</th><th>AA</th><th>w/b</th><th>AC</th><th>σ<sub>c-applied</sub>/f<sub>c</sub><sup>28 max</sup><br>(MPa)</th><th>f<sub>c</sub><sup>28 max</sup><br>(MPa)</th><th>Elastic modulus after<br>loading (GPa)</th><th>Elastic modulus after<br>unloading (GPa)</th><th>Permanent deformation<br>(μm/m)</th></tr></thead><tbody><tr><td>Mortar [52]</td><td>Ø × h/16 × 32</td><td>20±1</td><td>4200</td><td>Na<sub>2</sub>SO<sub>4</sub></td><td>0.42</td><td>8%</td><td>0.3</td><td>16.8±0.5</td><td>29.9±1.4</td><td>41.5±0.7</td><td>296±5</td></tr><tr><td>Mortar [52]</td><td>Ø × h/16 × 32</td><td>20±1</td><td>4200</td><td>K<sub>2</sub>SO<sub>4</sub></td><td>0.42</td><td>10%</td><td>0.3</td><td>14.7±0.5</td><td>22.7±2.6</td><td>36.2±0.1</td><td>448±5</td></tr><tr><td>Mortar [45]</td><td>7 × 7 × 28</td><td>20±3</td><td>4400</td><td>NaOH</td><td>0.5</td><td>2.5M</td><td>0.3</td><td>18.1±0.8</td><td>18.2±1</td><td>20.9±0.3</td><td>490<sup>2</sup></td></tr><tr><td>Mortar [109]</td><td>Ø × h/10 × 20</td><td>23±1</td><td>-</td><td>Na<sub>2</sub>SiO<sub>3</sub><br>+NaOH</td><td>0.38</td><td>1.6M</td><td>0.4</td><td>34</td><td>20.26</td><td>-</td><td>-</td></tr><tr><td>Concrete<br>[110]</td><td>Ø × h/10 × 20</td><td>20±2</td><td>4500</td><td>Na<sub>2</sub>CO<sub>3</sub></td><td>0.45</td><td>10%</td><td>0.4</td><td>38.4</td><td>28.7</td><td>-</td><td>-</td></tr><tr><td>Concrete<br>[110]</td><td>Ø × h/10 × 20</td><td>20±2</td><td>4500</td><td>Na<sub>2</sub>SiO<sub>3</sub></td><td>0.45</td><td>10%</td><td>0.4</td><td>40</td><td>29.8</td><td>-</td><td>-</td></tr><tr><td>Concrete<br>[110]</td><td>Ø × h/10 × 20</td><td>20±2</td><td>4500</td><td>Na<sub>2</sub>SiO<sub>3</sub><br>+Na<sub>3</sub>CO<sub>3</sub></td><td>0.45</td><td>5%+5%</td><td>0.4</td><td>37.6</td><td>28.1</td><td>-</td><td>-</td></tr></tbody></table>

1 Stainless steel balls are used in the measurement system.

2 is deduced by the compliance creep evolution, the specific evolution and the applied load.
584

References

585

586

[1] J.L. Provis, Alkali-activated materials, Cement and Concrete Research 114 (2018) 40-48.
https://doi.org/10.1016/j.cemconres.2017.02.009.

587

588

589

[2] B. Sun, G. Ye, G. De Schutter, A review: Reaction mechanism and strength of slag and fly ash-based alkali-activated materials, Construction and Building Materials 326 (2022) 126843.
https://doi.org/10.1016/j.conbuildmat.2022.126843.

590

591

592

[3] W. Tu, M. Zhang, Multiscale microstructure and micromechanical properties of alkali-activated concrete: A critical review, Cement and Concrete Composites 152 (2024) 105664.
https://doi.org/10.1016/j.cemconcomp.2024.105664.

593

594

[4] S.A. Bernal, J.L. Provis, Durability of alkali-activated materials: progress and perspectives, J. Am. Ceram. Soc. 97 (2014) 997-1008.
https://doi.org/10.1111/jace.12831.

595

596

[5] L. Coppola, D. Coffetti, E. Crotti, G. Gazzaniga, T. Pastore, The durability of one-part alkali-activated slag-based mortars in different environments, Sustainability 12 (2020) 3561.
https://doi.org/10.3390/su12093561.

597

598

599

[6] J. Zhang, C. Shi, Z. Zhang, Z. Ou, Durability of alkali-activated materials in aggressive environments: A review on recent studies, Construction and Building Materials 152 (2017) 598-613.
https://doi.org/10.1016/j.conbuildmat.2017.07.027.

600

601

[7] A. Darquennes, B. Espion, S. Staquet, How to assess the hydration of slag cement concretes?, Construction and Building Materials 40 (2013) 1012-1020.
https://doi.org/10.1016/j.conbuildmat.2012.09.087.

602

603

604

[8] A. Darquennes, S. Staquet, M.-P. Delplancke-Ogletree, B. Espion, Effect of autogenous deformation on the cracking risk of slag cement concretes, Cement and Concrete Composites 33 (2011) 368-379.
https://doi.org/10.1016/j.cemconcomp.2010.12.003.

605

606

[9] B. Lothenbach, G. Le Saout, M. Ben Haha, R. Figi, E. Wieland, Hydration of a low-alkali CEM III/B-SiO₂ cement (LAC), Cement and Concrete Research 42 (2012) 410-423.
https://doi.org/10.1016/j.cemconres.2011.11.008.

607

608

609

[10] S. Uppalapati, L. Vandewalle, Ö. Cizer, Monitoring the setting process of alkali-activated slag-fly ash cements with ultrasonic P-wave velocity, Construction and Building Materials 271 (2021) 121592.
https://doi.org/10.1016/j.conbuildmat.2020.121592.

610

611

612

[11] Y. Wang, X. He, Y. Su, J. Yang, B. Strnadel, X. Wang, Efficiency of wet-grinding on the mechano-chemical activation of granulated blast furnace slag (GBFS), Construction and Building Materials 199 (2019) 185-193.
https://doi.org/10.1016/j.conbuildmat.2018.11.245.

613

614

615

[12] Y. Li, J. Fang, L. Cheng, X. He, Y. Su, H. Tan, Mechanical performance, hydration characteristics and microstructures of high volume blast furnace ferronickel slag cement mortar by wet grinding activation, Construction and Building Materials 320 (2022) 126148.
https://doi.org/10.1016/j.conbuildmat.2021.126148.

616

617

618

[13] J. Zhang, H. Tan, X. He, W. Yang, X. Deng, Y. Su, J. Yang, Compressive strength and hydration process of ground granulated blast furnace slag-waste gypsum system managed by wet grinding, Construction and Building Materials 228 (2019) 116777.
https://doi.org/10.1016/j.conbuildmat.2019.116777.

619

620

[14] E. Altan, S.T. Erdoğan, Alkali activation of a slag at ambient and elevated temperatures, Cement and Concrete Composites 34 (2012) 131-139.
https://doi.org/10.1016/j.cemconcomp.2011.08.003.

621

622

623

[15] Q. Song, M.-Z. Guo, T.-C. Ling, A review of elevated-temperature properties of alternative binders: Supplementary cementitious materials and alkali-activated materials, Construction and Building Materials 341 (2022) 127894.
https://doi.org/10.1016/j.conbuildmat.2022.127894.

624

625

[16] W. Tu, M. Zhang, Behaviour of alkali-activated concrete at elevated temperatures: A critical review, Cement and Concrete Composites 138 (2023) 104961.
https://doi.org/10.1016/j.cemconcomp.2023.104961.
626

627

628

[17] P. Li, J. Tang, X. Chen, Y. Bai, Q. Li, Effect of temperature and pH on early hydration rate and apparent activation energy of alkali-activated slag, Advances in Materials Science and Engineering 2019 (2019) 1–13. https://doi.org/10.1155/2019/3531543.

629

630

631

[18] F. Puertas, C. Varga, M.M. Alonso, Rheology of alkali-activated slag pastes. Effect of the nature and concentration of the activating solution, Cement and Concrete Composites 53 (2014) 279–288. https://doi.org/10.1016/j.cemconcomp.2014.07.012.

632

633

634

[19] M. Ben Haha, G. Le Saout, F. Winnefeld, B. Lothenbach, Influence of activator type on hydration kinetics, hydrate assemblage and microstructural development of alkali activated blast-furnace slags, Cement and Concrete Research 41 (2011) 301–310. https://doi.org/10.1016/j.cemconres.2010.11.016.

635

636

637

[20] C.B. Cheah, L.E. Tan, M. Ramli, Recent advances in slag-based binder and chemical activators derived from industrial by-products – A review, Construction and Building Materials 272 (2021) 121657. https://doi.org/10.1016/j.conbuildmat.2020.121657.

638

639

[21] S. Aydın, B. Baradan, Effect of activator type and content on properties of alkali-activated slag mortars, Composites Part B: Engineering 57 (2014) 166–172. https://doi.org/10.1016/j.compositesb.2013.10.001.

640

641

642

[22] C. Yang, J. Liu, L. Liu, L. Kuang, S. Zhang, Z. Chen, J. Li, C. Shi, Effects of different activators on autogenous shrinkage of alkali-activated slag cement, Construction and Building Materials 446 (2024) 138018. https://doi.org/10.1016/j.conbuildmat.2024.138018.

643

644

[23] J.L. Provis, A. Palomo, C. Shi, Advances in understanding alkali-activated materials, Cement and Concrete Research 78 (2015) 110–125. https://doi.org/10.1016/j.cemconres.2015.04.013.

645

646

647

[24] S. Uppalapati, L. Vandewalle, Ö. Cizer, Autogenous shrinkage of slag-fly ash blends activated with hybrid sodium silicate and sodium sulfate at different curing temperatures, Construction and Building Materials 265 (2020) 121276. https://doi.org/10.1016/j.conbuildmat.2020.121276.

648

649

650

[25] R. Navarro, E. Zornoza, I. Sánchez, E.G. Alcocel, Influence of the type and concentration of the activator on the microstructure of alkali activated SiMn slag pastes, Construction and Building Materials 342 (2022) 128067. https://doi.org/10.1016/j.conbuildmat.2022.128067.

651

652

[26] Z. Li, Y. Chen, J.L. Provis, Ö. Cizer, G. Ye, Autogenous shrinkage of alkali-activated slag: A critical review, Cement and Concrete Research 172 (2023) 107244. https://doi.org/10.1016/j.cemconres.2023.107244.

653

654

655

[27] D. Ravikumar, N. Neithalath, Reaction kinetics in sodium silicate powder and liquid activated slag binders evaluated using isothermal calorimetry, Thermochimica Acta 546 (2012) 32–43. https://doi.org/10.1016/j.tca.2012.07.010.

656

657

[28] O. Burciaga-Díaz, J.I. Escalante-García, Structure, Mechanisms of reaction, and strength of an alkali-activated blast-furnace slag, J. Am. Ceram. Soc. 96 (2013) 3939–3948. https://doi.org/10.1111/jace.12620.

658

659

660

661

[29] T. Luukkonen, H. Sreenivasan, Z. Abdollahnejad, J. Yliniemi, A. Kantola, V.-V. Telkki, P. Kinnunen, M. Illikainen, Influence of sodium silicate powder silica modulus for mechanical and chemical properties of dry-mix alkali-activated slag mortar, Construction and Building Materials 233 (2020) 117354. https://doi.org/10.1016/j.conbuildmat.2019.117354.

662

663

664

[30] Z. Yan, Z. Sun, J. Yang, H. Yang, Y. Ji, K. Hu, Mechanical performance and reaction mechanism of copper slag activated with sodium silicate or sodium hydroxide, Construction and Building Materials 266 (2021) 120900. https://doi.org/10.1016/j.conbuildmat.2020.120900.

665

666

667

[31] A. Adesina, Performance and sustainability overview of sodium carbonate activated slag materials cured at ambient temperature, Resources, Environment and Sustainability 3 (2021) 100016. https://doi.org/10.1016/j.resenv.2021.100016.
668

669

670

[32] X. Ke, S.A. Bernal, J.L. Provis, Controlling the reaction kinetics of sodium carbonate-activated slag cements using calcined layered double hydroxides, Cement and Concrete Research 81 (2016) 24-37.
https://doi.org/10.1016/j.cemconres.2015.11.012.

671

672

673

[33] J. Park, S. Kim, J. Moon, Substantially enhanced reaction of steel slag binder stimulated by the addition of sodium sulfate and thermal treatment, Cement and Concrete Composites 152 (2024) 105679.
https://doi.org/10.1016/j.cemconcomp.2024.105679.

674

675

676

[34] A.M. Rashad, Y. Bai, P.A.M. Basheer, N.B. Milestone, N.C. Collier, Hydration and properties of sodium sulfate activated slag, Cement and Concrete Composites 37 (2013) 20-29.
https://doi.org/10.1016/j.cemconcomp.2012.12.010.

677

678

679

[35] 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, Cement and Concrete Research 135 (2020) 106110.
https://doi.org/10.1016/j.cemconres.2020.106110.

680

681

682

[36] Y. Zhao, J. Qiu, S. Zhang, Z. Guo, Z. Ma, X. Sun, J. Xing, Effect of sodium sulfate on the hydration and mechanical properties of lime-slag based eco-friendly binders, Construction and Building Materials 250 (2020) 118603.
https://doi.org/10.1016/j.conbuildmat.2020.118603.

683

684

685

[37] A.S.S.S. Nedunuri, S. Muhammad, Improving the workability and workable time of sodium hydroxide-activated ground granulated blast furnace slag binder-based concrete, CEMENT 16 (2024) 100106.
https://doi.org/10.1016/j.cement.2024.100106.

686

687

688

[38] A. Adesina, C. Rodrigue Kaze, Physico-mechanical and microstructural properties of sodium sulfate activated materials: A review, Construction and Building Materials 295 (2021) 123668.
https://doi.org/10.1016/j.conbuildmat.2021.123668.

689

690

691

[39] J.C. Kuri, M. Nuruzzaman, P.K. Sarker, Sodium sulphate resistance of geopolymers mortar produced using ground ferronickel slag with fly ash, Ceramics International 49 (2023) 2765-2773.
https://doi.org/10.1016/j.ceramint.2022.09.258.

692

693

694

[40] Z. Huang, Q. Wang, J. Lu, The effects of cations and concentration on reaction mechanism of alkali-activated blast furnace ferronickel slag, Composites Part B: Engineering 236 (2022) 109825.
https://doi.org/10.1016/j.compositesb.2022.109825.

695

696

[41] Q. Fu, M. Bu, Z. Zhang, W. Xu, Q. Yuan, D. Niu, Hydration characteristics and microstructure of alkali-activated slag concrete: A review, Engineering 20 (2023) 162-179.
https://doi.org/10.1016/j.eng.2021.07.026.

697

698

699

[42] A. Hilaire, F. Benboudjema, A. Darquennes, Y. Berthaud, G. Nahas, Modeling basic creep in concrete at early-age under compressive and tensile loading, Nuclear Engineering and Design 269 (2014) 222-230.
https://doi.org/10.1016/j.nucengdes.2013.08.034.

700

701

702

[43] X. Hu, Z. Shi, C. Shi, Z. Wu, B. Tong, Z. Ou, G. de Schutter, Drying shrinkage and cracking resistance of concrete made with ternary cementitious components, Construction and Building Materials 149 (2017) 406-415.
https://doi.org/10.1016/j.conbuildmat.2017.05.113.

703

704

705

[44] Z. Li, B. Delsaute, T. Lu, A. Kostiuchenko, S. Staquet, G. Ye, A comparative study on the mechanical properties, autogenous shrinkage and cracking proneness of alkali-activated concrete and ordinary Portland cement concrete, Construction and Building Materials 292 (2021) 123418.
https://doi.org/10.1016/j.conbuildmat.2021.123418.

706

707

708

709

[45] F. Rifai, A. Darquennes, F. Benboudjema, B. Muzeau, L. Stefan, Thermo-chemo-mechanical behavior of alkali-activated slag materials - Focus on early-age, (2017). 2nd International RILEM/COST Conference on Early Age Cracking and Serviceability in Cement-based Materials and Structures - EAC2, Sep 2017, Brussels, Belgium.
https://hal.science/hal-02417342.
710

711

712

[46] A. Naqi, B. Delsaute, M. Königsberger, S. Staquet, Effect of solution-to-binder ratio and alkalinity on setting and rarly-Age properties of alkali-activated slag-fly ash binders, Materials 16 (2022) 373. https://doi.org/10.3390/ma16010373.

713

714

[47] B. Delsaute, J. Torrenti, S. Staquet, Prediction of the basic creep of concrete with high substitution of Portland cement by mineral additions at early age, Structural Concrete 22 (2021). https://doi.org/10.1002/suco.201900313.

715

716

[48] A. Aili, Shrinkage and creep of cement-based materials under multiaxial load : poromechanical modeling for application in nuclear industry, 2017. http://www.theses.fr/2017PESC1014/document.

717

718

719

[49] R. Caron, R.A. Patel, F. Dehn, Experimental study on basic and drying creep for an alkali-activated slag concrete and comparison with existing creep models, Structural Concrete 24 (2023) 6405-6420. https://doi.org/10.1002/suco.202300134.

720

721

[50] N. Ranaivomanana, S. Multon, A. Turatsinze, Basic creep of concrete under compression, tension and bending, Construction and Building Materials 38 (2013) 173-180. https://doi.org/10.1016/j.conbuildmat.2012.08.024.

722

723

724

[51] P. Rossi, J.-L. Tailhan, F. Le Maou, L. Gaillet, E. Martin, Basic creep behavior of concretes investigation of the physical mechanisms by using acoustic emission, Cement and Concrete Research 42 (2012) 61-73. https://doi.org/10.1016/j.cemconres.2011.07.011.

725

726

727

728

729

[52] L. Li, K. Hannawi, A. Darquennes, Delayed deformations of Na- and K- sulphates activated blast-furnace slag mortars, in: A. Jędrzejewska, F. Kanavaris, M. Azenha, F. Benboudjema, D. Schlicke (Eds.), International RILEM Conference on Synergising Expertise towards Sustainability and Robustness of Cement-Based Materials and Concrete Structures, Springer Nature Switzerland, Cham, 2023: pp. 403-414. https://doi.org/10.1007/978-3-031-33187-9_38.

730

731

[53] M. Elzeadani, D.V. Bompa, A.Y. Elghazouli, Creep response of rubberized alkali-activated concrete, J. Mater. Civ. Eng. 35 (2023) 04023434. https://doi.org/10.1061/JMCEE7.MTENG-16353.

732

733

734

[54] A.M. Humad, J.L. Provis, K. Habermehl-Cwirzen, M. Rajczakowska, A. Cwirzen, Creep and long-term properties of alkali-activated swedish-slag concrete, J. Mater. Civ. Eng. 33 (2021) 04020475. https://doi.org/10.1061/(ASCE)MT.1943-5533.0003381.

735

736

737

738

[55] H. Sounean, H. Kinda, D. Aveline, The Cracking Sensitivity of a Na-Geopolymer, in: F. Kanavaris, F. Benboudjema, M. Azenha (Eds.), International RILEM Conference on Early-Age and Long-Term Cracking in RC Structures, Springer International Publishing, Cham, 2021: pp. 165-174. https://doi.org/10.1007/978-3-030-72921-9_14.

739

740

741

[56] D.H. Nguyen, V.T. Nguyen, P. Lura, V.T.N. Dao, Temperature-stress testing machine – A state-of-the-art design and its unique applications in concrete research, Cement and Concrete Composites 102 (2019) 28-38. https://doi.org/10.1016/j.cemconcomp.2019.04.019.

742

743

744

[57] D.-B. Nguyen, C.-J. Wu, W.-C. Liao, Experimental study and development of prediction model of self-compacting concrete incorporating fly ash on compressive creep behaviour, ACT 21 (2023) 596-610. https://doi.org/10.3151/jact.21.596.

745

746

[58] M. Hojati, F. Rajabipour, A. Radlińska, Creep of alkali-activated cement mixtures, Case Studies in Construction Materials 16 (2022) e00954. https://doi.org/10.1016/j.cscm.2022.e00954.

747

748

[59] L. Li, A. Darquennes, K. Hannawi, C. Che, Effect of the Alkali-Sulphate Activators on the Hydration Process of Blast-Furnace Slag Mortars and Pastes, Materials 18 (2025) 514. https://doi.org/10.3390/ma18030514.

749

750

751

[60] J. Carette, F. Soleilhet, F. Benboudjema, X. Ma, G. Nahas, K. Abahri, A. Darquennes, R. Bennacer, Identifying the mechanisms of concrete drying: An experimental-numerical approach, Construction and Building Materials 230 (2020) 117001. https://doi.org/10.1016/j.conbuildmat.2019.117001.
752

753

754

[61] O.A. Mohamed, O. Najm, H.A. Zuaiter, Setting time, sulfuric acid resistance, and strength development of alkali-activated mortar with slag & fly ash binders, Results in Engineering 21 (2024) 101711. https://doi.org/10.1016/j.rineng.2023.101711.

755

756

757

[62] C. Shi, Y. Liu, L. Qin, M.T. Ghafoor, Z. Zhang, Time-dependent evolution of the microstructural characteristics of C-A-S-H and N-A-S-H gels with various Ca/Si and Al/Si ratios, Construction and Building Materials 502 (2025) 144408. https://doi.org/10.1016/j.conbuildmat.2025.144408.

758

759

760

[63] J. Fu, M.W. Bligh, I. Shikhov, A.M. Jones, C. Holt, L.M. Keyte, F. Moghaddam, C.H. Arns, S.J. Foster, T.D. Waite, A microstructural investigation of a Na₂SO₄ activated cement-slag blend, Cement and Concrete Research 150 (2021) 106609. https://doi.org/10.1016/j.cemconres.2021.106609.

761

762

763

[64] U. De Filippis, E. Prud'homme, S. Meille, Sodium hydroxide substitution in slag activating mixes: A potential pathway to more sustainable slag-based binders, Construction and Building Materials 300 (2021) 124183. https://doi.org/10.1016/j.conbuildmat.2021.124183.

764

765

766

[65] F. Kong, R. Xu, A. Wang, T. Hong, X. Shi, Effects of different activators on workability, mechanical properties and microstructure of alkali-activated slag-copper slag, Construction and Building Materials 463 (2025) 140151. https://doi.org/10.1016/j.conbuildmat.2025.140151.

767

768

769

[66] Y. Zheng, H. Zhong, K. Sun, B. Shen, K. Cui, Y. Zhao, G. Xiong, Q. Qin, Effect of composite activator on hydration kinetics and micromechanical properties of alkali-activated slag, Cement and Concrete Composites 160 (2025) 106047. https://doi.org/10.1016/j.cemconcomp.2025.106047.

770

771

772

[67] Y. Wei, X. Gao, S. Liang, Nanoindentation-based study of the micro-mechanical properties, structure, and hydration degree of slag-blended cementitious materials, J Mater Sci 51 (2016) 3349-3361. https://doi.org/10.1007/s10853-015-9650-4.

773

774

775

[68] C. Ma, B. Zhao, S. Guo, G. Long, Y. Xie, Properties and characterization of green one-part geopolymer activated by composite activators, Journal of Cleaner Production 220 (2019) 188-199. https://doi.org/10.1016/j.jclepro.2019.02.159.

776

777

778

[69] T. A, S. Om, Effect of varying molarity and curing conditions on the mechanical and microstructural characteristics of alkali activated GGBS binder, Mater. Res. Express 10 (2023) 095305. https://doi.org/10.1088/2053-1591/acf6f4.

779

780

781

[70] H. Lee, V. Vimonsatit, P. Chindaprasirt, T. Ngo, P. Mendis, Creep properties of cement and alkali activated fly ash materials using nanoindentation technique, Construction and Building Materials 168 (2018) 547-555. https://doi.org/10.1016/j.conbuildmat.2018.02.166.

782

783

784

[71] Y. Zhang, M. Liang, Y. Gan, O. Çopuroğlu, Micro-Mechanical Properties of Slag Rim Formed in Cement-Slag System Evaluated by Nanoindentation Combined with SEM, Materials 15 (2022) 6347. https://doi.org/10.3390/ma15186347.

785

786

787

[72] D. Hou, H. Li, L. Zhang, J. Zhang, Nano-scale mechanical properties investigation of C-S-H from hydrated tri-calcium silicate by nano-indentation and molecular dynamics simulation, Construction and Building Materials 189 (2018) 265-275. https://doi.org/10.1016/j.conbuildmat.2018.08.215.

788

789

790

791

[73] G. Kim, S. Im, H. Jee, H. Suh, S. Cho, M. Kanematsu, S. Morooka, T. Koyama, Y. Nishio, A. Machida, J. Kim, S. Bae, Effect of magnesium silicate hydrate (M-S-H) formation on the local atomic arrangements and mechanical properties of calcium silicate hydrate (C-S-H): In situ X-ray scattering study, Cement and Concrete Research 159 (2022) 106869. https://doi.org/10.1016/j.cemconres.2022.106869.

792

793

794

[74] S. Ding, Z. Hu, Z. Geng, J. Liu, W. Xu, Monitoring the effect of MgO expansive additive on micromechanical properties of cementitious materials using a combined method of EDS-nanoindentation, Construction and Building Materials 411 (2024) 134589. https://doi.org/10.1016/j.conbuildmat.2023.134589.
795

796

797

798

799

800

801

802

803

804

805

806

807

808

809

810

811

812

813

814

815

816

817

818

819

820

821

822

823

824

825

826

827

828

829

830

831

832

833

834

835

836

837

[75] L. Li, A. Darquennes, K. Hannawi, C. Che, Effect of the Alkali-Sulphate Activators on the Hydration Process of Blast-Furnace Slag Mortars and Pastes, Materials 18 (2025) 514. https://doi.org/10.3390/ma18030514.

[76] S. Barzgar, B. Lothenbach, M. Tarik, A. Di Giacomo, C. Ludwig, The effect of sodium hydroxide on Al uptake by calcium silicate hydrates (C-S-H), Journal of Colloid and Interface Science 572 (2020) 246–256. https://doi.org/10.1016/j.jcis.2020.03.057.

[77] V. Šmilauer, Z.P. Bažant, Identification of viscoelastic C-S-H behavior in mature cement paste by FFT-based homogenization method, Cement and Concrete Research 40 (2010) 197–207. https://doi.org/10.1016/j.cemconres.2009.10.003.

[78] W. Wilson, L. Sorelli, A. Tagnit-Hamou, Unveiling micro-chemo-mechanical properties of C-(A)-S-H and other phases in blended-cement pastes, Cement and Concrete Research 107 (2018) 317–336. https://doi.org/10.1016/j.cemconres.2018.02.010.

[79] E. Bernard, B. Lothenbach, D. Rentsch, I. Pochard, A. Dauzères, Formation of magnesium silicate hydrates (M-S-H), Physics and Chemistry of the Earth, Parts A/B/C 99 (2017) 142–157. https://doi.org/10.1016/j.pce.2017.02.005.

[80] R.L. Frost, H.J. Spratt, S.J. Palmer, Infrared and near-infrared spectroscopic study of synthetic hydrotalcites with variable divalent/trivalent cationic ratios, Spectrochimica Acta Part A: Molecular and Biomolecular Spectroscopy 72 (2009) 984–988. https://doi.org/10.1016/j.saa.2008.12.018.

[81] D.Y. Yang, R. Guo, Experimental study on modulus and hardness of ettringite, Experimental Techniques 38 (2014) 6–12. https://doi.org/10.1111/j.1747-1567.2011.00744.x.

[82] X. Deng, M. Li, Y. Wang, J. Wang, J. Zhang, Z. Yang, X. He, J. Yang, H. Tan, Impact of ettringite seeding on hydration, strength and shrinkage of Na2SO4 activated slag, Composites Part B: Engineering 276 (2024) 111374. https://doi.org/10.1016/j.compositesb.2024.111374.

[83] K. Velez, S. Maximilien, D. Damidot, G. Fantozzi, F. Sorrentino, Determination by nanoindentation of elastic modulus and hardness of pure constituents of Portland cement clinker, Cement and Concrete Research 31 (2001) 555–561. https://doi.org/10.1016/S0008-8846(00)00505-6.

[84] D. Lootens, D.P. Bentz, On the relation of setting and early-age strength development to porosity and hydration in cement-based materials, Cement and Concrete Composites 68 (2016) 9–14. https://doi.org/10.1016/j.cemconcomp.2016.02.010.

[85] N. Alharbi, B. Varela, R. Hailstone, Alkali-activated slag characterization by scanning electron microscopy, X-ray microanalysis and nuclear magnetic resonance spectroscopy, Materials Characterization 168 (2020) 110504. https://doi.org/10.1016/j.matchar.2020.110504.

[86] M. Nedeljković, K. Arbi, Y. Zuo, G. Ye, Microstructural and mineralogical analysis of alkali activated fly ash-slag pastes, Microdurability 2016: the International Rilem Conference on Microstructure Related Durability of Cementitious Composites. 2016. https://resolver.tudelft.nl/uuid:11d19e9c-87cb-41ba-878a-4bb49d2c57c3.

[87] K. Yang, C.E. White, Multiscale pore structure determination of cement paste via simulation and experiment: The case of alkali-activated metakaolin, Cement and Concrete Research 137 (2020) 106212. https://doi.org/10.1016/j.cemconres.2020.106212.

[88] Y. Ma, G. Wang, G. Ye, J. Hu, A comparative study on the pore structure of alkali-activated fly ash evaluated by mercury intrusion porosimetry, N2 adsorption and image analysis, J Mater Sci 53 (2018) 5958–5972. https://doi.org/10.1007/s10853-017-1965-x.

[89] Y. Zuo, G. Ye, Pore Structure Characterization of Sodium Hydroxide Activated Slag Using Mercury Intrusion Porosimetry, Nitrogen Adsorption, and Image Analysis, Materials 11 (2018) 1035. https://doi.org/10.3390/ma11061035.
838

839

840

[90] J. Sun, S. Hou, Y. Guo, W. He, J. Bao, Y. Cui, P. Zhang, Sustainable utilization of alkali-activated steel slag material: Effects of silicate modulus and GBFS on fresh, mechanical and pore structure properties, Developments in the Built Environment 18 (2024) 100410. https://doi.org/10.1016/j.dibe.2024.100410.

841

842

843

[91] X. Hu, C. Shi, Z. Shi, L. Zhang, Compressive strength, pore structure and chloride transport properties of alkali-activated slag/fly ash mortars, Cement and Concrete Composites 104 (2019) 103392. https://doi.org/10.1016/j.cemconcomp.2019.103392.

844

845

846

[92] A. Keulen, Q.L. Yu, S. Zhang, S. Grünewald, Effect of admixture on the pore structure refinement and enhanced performance of alkali-activated fly ash-slag concrete, Construction and Building Materials 162 (2018) 27–36. https://doi.org/10.1016/j.conbuildmat.2017.11.136.

847

848

[93] O.A. Mohamed, R. Al-Khattab, W. Al-Hawat, Resistance to acid degradation, sorptivity, and setting time of geopolymer mortars, Front. Struct. Civ. Eng. 16 (2022) 781–791. https://doi.org/10.1007/s11709-022-0862-9.

849

850

851

[94] C. Duran Atiş, C. Bilim, Ö. Çelik, O. Karahan, Influence of activator on the strength and drying shrinkage of alkali-activated slag mortar, Construction and Building Materials 23 (2009) 548–555. https://doi.org/10.1016/j.conbuildmat.2007.10.011.

852

853

854

[95] H. Gao, I. Shikhov, E. Hamed, A. Hajimohammadi, I. Al-Damad, C. Arns, S.J. Foster, New insights on the basic creep mechanism of one-part alkali activated slag and fly ash paste, Cement and Concrete Research 186 (2024) 107691. https://doi.org/10.1016/j.cemconres.2024.107691.

855

856

[96] B. Walkley, X. Ke, J.L. Provis, S.A. Bernal, Activator Anion Influences the Nanostructure of Alkali-Activated Slag Cements, J. Phys. Chem. C 125 (2021) 20727–20739. https://doi.org/10.1021/acs.jpcc.1c07328.

857

858

[97] A.R. Brough, A. Atkinson, Sodium silicate-based, alkali-activated slag mortars Part I. Strength, hydration and microstructure, Cement and Concrete Research (2002). https://doi.org/10.1016/S0008-8846(02)00717-2.

859

860

861

[98] H. Xu, H. Yin, P. Ge, Design and microstructural analysis of the mixture proportion of alkali-activated fly ash-slag composite cementitious material, Mater. Res. Express 11 (2024) 115303. https://doi.org/10.1088/2053-1591/ad9083.

862

863

864

[99] X. Zhou, Y. Chen, S. Dong, H. Li, Geopolymerization kinetics of steel slag activated gasification coal fly ash: A case study for amorphous-rich slags, Journal of Cleaner Production 379 (2022) 134671. https://doi.org/10.1016/j.jclepro.2022.134671.

865

866

867

[100] R.J. Thomas, B.S. Gebregziabiher, A. Giffin, S. Peethamparan, Micromechanical properties of alkali-activated slag cement binders, Cement and Concrete Composites 90 (2018) 241–256. https://doi.org/10.1016/j.cemconcomp.2018.04.003.

868

869

[101] R.J. Thomas, S. Peethamparan, Alkali-activated concrete: Engineering properties and stress–strain behavior, Construction and Building Materials 93 (2015) 49–56. https://doi.org/10.1016/j.conbuildmat.2015.04.039.

870

871

872

[102] F. Puertas, M. Palacios, H. Manzano, J.S. Dolado, A. Rico, J. Rodríguez, A model for the C-A-S-H gel formed in alkali-activated slag cements, Journal of the European Ceramic Society 31 (2011) 2043–2056. https://doi.org/10.1016/j.jeurceramsoc.2011.04.036.

873

874

[103] Fernandez-Jimenez, P. Angel, L.-H. Cecilio, Engineering properties of alkali-activated fly ash concrete, (2006). ACI Materials Journal 103(2):106-112. https://doi.org/10.14359/15261.

875

876

[104] M. Sofi, J.S.J. Van Deventer, P.A. Mendis, G.C. Lukey, Engineering properties of inorganic polymer concretes (IPCs), Cement and Concrete Research 37 (2007) 251–257. https://doi.org/10.1016/j.cemconres.2006.10.008.

877

878

879

[105] N.K. Lee, H.K. Lee, Setting and mechanical properties of alkali-activated fly ash/slag concrete manufactured at room temperature, Construction and Building Materials 47 (2013) 1201–1209. https://doi.org/10.1016/j.conbuildmat.2013.05.107.
880

881

882

[106] J. Xie, J. Wang, B. Zhang, C. Fang, L. Li, Physicochemical properties of alkali activated GGBS and fly ash geopolitical recycled concrete, Construction and Building Materials 204 (2019) 384-398. https://doi.org/10.1016/j.conbuildmat.2019.01.191.

883

884

885

[107] D. Ballekere Kumarappa, S. Peethamparan, M. Ngami, Autogenous shrinkage of alkali activated slag mortars: Basic mechanisms and mitigation methods, Cement and Concrete Research 109 (2018) 1-9. https://doi.org/10.1016/j.cemconres.2018.04.004.

886

887

888

[108] N. Reviron, Nanthilde Reviron. Etude du fluage des bétons en traction. Application aux enceintes de confinement des centrales nucléaires à eau sous pression. Matériaux. École normale supérieure de Cachan - ENS Cachan, (2009) 40-60. https://theses.fr/2009DENSA006

889

890

[109] J. Ma, F. Dehn, Shrinkage and creep behavior of an alkali-activated slag concrete, Structural Concrete 18 (2017) 801-810. https://doi.org/10.1002/suco.201600147.

891

892

893

[110] A. Humad, Shrinkage and related properties of alkali-activated binders based on high-MgO blast furnace slag, Luleå University of Technology, Department of Civil, Environment and Natural Resources Engineering, Luleå, https://doi.org/2019.10.13140/RG.2.2.13804.13448.

894

895

896

[111] J.L. Tailhan, L. D'aloia, P. Autuori, Simulations numériques du comportement au jeune âge des structures en béton: Modélisation et retour d'expérience, Laboratoire Central Des Ponts et Chaussées, Paris, France (2011) 20-40. https://hal.science/hal-00612155.

897

898

[112] M. Vandamme, Two models based on local microscopic relaxations to explain long-term basic creep of concrete, Proc. R. Soc. A. 474 (2018) 20180477. https://doi.org/10.1098/rspa.2018.0477.

899

900

[113] Z. Bažant, A. Hauggaard, S. Baweja, F. Ulm, Microprestress-solidification theory for concrete creep. I: Aging and drying effects., (1997). https://doi.org/10.1061/(ASCE)0733-9399(1997)123:11(1188)/

901

902

903

[114] P. Duxson, A. Fernández-Jiménez, J.L. Provis, G.C. Lukey, A. Palomo, J.S.J. Van Deventer, Geopolymer technology: the current state of the art, J Mater Sci 42 (2007) 2917-2933. https://doi.org/10.1007/s10853-006-0637-z.

904

905

[115] D. Zheng, T. Ji, G. Wang, Effect of CaO on the autogenous shrinkage of alkali-activated slag mortar, Advances in Materials Science and Engineering (2021) 9918834. https://doi.org/10.1155/2021/9918834.

906

907

908

[116] P. Duxson, J.L. Provis, G.C. Lukey, S.W. Mallicoat, W.M. Kriven, J.S.J. Van Deventer, Understanding the relationship between geopolymer composition, microstructure and mechanical properties, Colloids and Surfaces A: Physicochemical and Engineering Aspects 269 (2005) 47-58. https://doi.org/10.1016/j.colsurfa.2005.06.060.

909

910

911

[117] F. Lolli, J.J. Thomas, K.E. Kurtis, F. Cucinotta, E. Masoero, Early age volume changes in metakaolin geopolymers: Insights from molecular simulations and experiments, Cement and Concrete Research 144 (2021) 106428. https://doi.org/10.1016/j.cemconres.2021.106428.

912

913

[118] Un, Chi Hou (2017). Creep behaviour of geopolymer concrete. Swinburne. Thesis 126-135. https://doi.org/10.25916/sut.26279014.v1.

914

915

916

[119] P.B. Bamforth, In situ measurement of the effect of partial Portland cement replacement using either fly ash or ground granulated blast-furnace slag on the performance of mass concrete, Proceedings of the Institution of Civil Engineers (1980) 69 (3): 777-800. https://doi.org/10.1680/iicep.1980.2377.

917

918

[120] Q. Zhao, X. Liu, J. Jiang, Effect of curing temperature on creep behavior of fly ash concrete, Construction and Building Materials 96 (2015) 326-333. https://doi.org/10.1016/j.conbuildmat.2015.08.030.

919

920

921

[121] A.E. Klausen, T. Kanstad, Ø. Bjøntegaard, E. Sellevold, Comparison of tensile and compressive creep of fly ash concretes in the hardening phase, Cement and Concrete Research 95 (2017) 188-194. https://doi.org/10.1016/j.cemconres.2017.02.018.

922

923
924

925

926

927

928

929

930

931

932

933

934

935

936

937

938

939

Declaration of interests

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

[ ] The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:

<figure><img src="image_17.png" /></figure>