The great challenge of cemented tailings backfill (CTB) is difficult simultaneously maintaining its excellent mechanical properties and low cost. Fly ash (FA) can potentially address this problem and further replace cement in favor of low carbon development. However, its mechanism on CTB with low cement dosage and low Ca system remains unclear. Consequently, this study conducted uniaxial compression, X-ray diffraction (XRD), and scanning electron microscopy (SEM)–energy dispersive spectrometer (EDS) tests to investigate the effect of FA dosage on the mechanical property and microstructure of CTB. A molecular model of FA-CSH was constructed to reproduce the molecular structure evolution of CTB with FA based on the test results. The influences of FA dosage and calcium/silica molar ratio (Ca/Si ratio) on the matrix strength and failure model were analyzed to reveal the mechanism of FA on calcium silicate hydrated (C–S–H). The results show that the strength of CTB increases initially and then decreases with FA dosage, and the FA supplement leads to a decrease in Ca(OH)2 diffraction intensity and Ca/Si ratio around the FA particles. XRD and SEM–EDS findings show that the Ca/Si ratio of C–S–H decreases with the progression of hydration. The FA-CSH model indicates that FA can reinforce the silica chain of C–S–H to increase the matrix strength. However, this enhancement is weakened by supplementing excessive FA dosage. In addition, the hydrogen bonds among water molecules deteriorate, reducing the matrix strength. A low Ca/Si ratio results in an increase in water molecules and a decrease in the ionic bonds combined with Ca2+. The hydrogen bonds among water molecules cannot withstand high stresses, resulting in a reduction in strength. The water absorption of the FA-CSH model is negatively correlated with the FA dosage and Ca/Si ratio. The use of optimal FA dosage and Ca/Si ratio leads to suitable water absorption, which further affects the failure mode of FA-CSH.
With the gradual depletion of shallow resources, the development of deep earth resources has become a necessity for human development. Mining deep earth resources exacerbates environmental damage and solid waste emissions. Recycling mining wastes, such as gangue and tailings, to produce filling materials is the most effective sustainable approach for addressing waste hazards while ensuring the safety of underground structures [1–4].
The cemented tailings backfill (CTB) has consistently encountered the challenge of maintaining a balance between low cost and high strength [5–6]. Researchers have demonstrated that the addition of nanomaterials, such as graphene oxide [7] and carbon nanotubes [8], can considerably enhance the strength of cement-based materials; however, the reality of low-cost engineering constrains the use of expensive additives. Other researchers have attempted to improve the performance of cemented filling materials by adding chemical admixtures [9–10], optimizing the particle size distribution [11], and enhancing curing conditions. However, these complex procedures can hamper engineering progress and potentially harm the environment. The partial substitution of cement with fly ash (FA) as a supplementary cementitious material [12] can improve the settlement of slurry aggregates, reduce the water/cement ratio of the slurry, and enhance the late strength of the filling material [13]. Moreover, this substitution can reduce the carbon emissions of cement production, promote the recycling of FA, and lower the engineering cost of cemented filling materials [14].
Rodger and Groves [15] demonstrated that in comparison with the corresponding properties of pure cement pastes, the microstructure of cement–FA blends is similar, but the calcium/silica molar ratio (Ca/Si ratio) of the C–S–H gel material is lower. The impact of FA, which serves as a potential activator, on CTB performance in low Ca/Si ratio environments remains unclear; even at low FA dosages, the early compressive strength of cement–FA composites is usually low due to the slow reaction of FA [13,16–18]. FA may compete with Ca2+ and degrade C–S–H, thereby affecting the durability of CTB. Despite the multitude of multiscale investigations aiming to elucidate this matter from the perspective of macroscopic mechanics to the microstructural level, experimental results are sometimes challenged by material preparation and experimental equipment accuracy related to the relevant scales [19]. The effect of FA on C–S–H lacks molecular-scale evidence. In recent years, molecular dynamics (MD) simulations have become a powerful tool to provide complementary explanations for experimental results [20–21].
Uniaxial compression, X-ray diffraction (XRD), and scanning electron microscopy–energy dispersive spectrometer (SEM–EDS) tests were conducted in this study to investigate the effect of FA dosage on the mechanical properties and microstructure of CTB. Additionally, a molecular model of FA-CSH was also constructed to reproduce the molecular structure evolution of CTB with FA based on the test results. The influences of FA dosage and Ca/Si ratio on the matrix strength, Mises strain field, and failure model were analyzed to reveal the mechanism of FA influence on C–S–H.
2.
Experimental
2.1
Materials
Composite Portland cement (CPC32.5) and FA were used as cementitious materials. Figs. S1(a) and S2(a) show the XRD spectrum and particle size distribution (PSD) of CPC32.5, respectively; Figs. S1(b) and S2(b) display the XRD spectrum and PSD of FA, respectively. The XRD spectrum of FA provides an important basis for establishing a noncrystalline model based on partial melting. Fig. S2(c) illustrates the PSD of the aggregate whose tailings were used as the cementitious filler material. Table 1 shows the chemical composition of the particulate materials used in this study.
A uniaxial compression strength (UCS) test was conducted on CTB using an MTS816 mechanical test system at a loading rate of 0.3 mm/min. In brief, 1 cm3 fragments were soaked in anhydrous ethanol before being dried at 313.15 K. One fragment from each group was then encapsulated in epoxy resin, polished, and gold-sputtered before being observed using TESCAN MAIA3 FSEM to study the microstructure morphology of CTB. The elemental distribution in the energy spectrum analysis region was also analyzed. The remaining fragments were ground to 45 μm and studied using a D8 Advance diffractometer to investigate the XRD spectra of the materials.
2.3
Sample-making procedure for CTB
Four CTB specimens featuring various FA dosages were designed with a maximum FA-to-cementitious material mass ratio of 120% to illustrate the effects of different dosages of FA on the strength performance of CTB. In addition, four curing ages (3, 7, 14, and 28 d) were designed under the same FA dosage to observe the effects of curing age on the strength performance of CTB. CTB specimens with varying amounts of FA were prepared according to the mix proportions specified in Table 2. For accurate test data, three CTB specimens were prepared for each condition, and the test results were averaged.
Table
2.
Preparation program of CTB samples
Group
FA dosage / g
Mass of cement / g
Mass of tailings / g
W/B
Curing time / d
A
15
50
450
0.8
3/7/14/28
B
30
50
450
0.8
3/7/14/28
C
45
50
450
0.8
3/7/14/28
D
60
50
450
0.8
3/7/14/28
Note: W/B represents the mass ratio of water to binder.
The specific steps for preparing the CTB specimens were as follows: (1) weighing the aggregates, cement, FA, and water according to the mass ratios specified in Table 2; (2) pouring the FA and cement into the water and mixing them using a mixer at 120 r/min for 5 min to form a homogeneous slurry; (3) adding the tailings to the slurry and mixing them at the same time and speed as in step (2); (4) pouring the mixture into a cylindrical mold of dimensions ϕ50 mm × 100 mm and fixing the mold on a vibrating table for compaction; (5) demolding the CTB specimens after 24 h and placing them in a curing chamber maintained at a standard temperature of (293.15 ± 2) K and relative humidity of 95% for the designated curing period.
2.4
Molecular dynamics simulation
The mechanical properties of C–S–H structures were simulated using the Clay Force Field (CLAYFF), an ionic force field with parameters obtained via density functional theory and quantum chemical calculations. Under the CLAYFF force field, all chemical bond models are deleted after the modeling is completed, all atoms are represented by point charges and allowed to translate completely and freely, and the interactions between point charges enable the system to rapidly attain equilibrium. The simple point charge model [22] was used for H2O molecules, and the CLAYFF force field parameters were adopted from literature [23].
The amorphous C–S–H model was constructed using the method proposed by Pellenq [24], in which a Tobermorite monoclinic cell with space group P21 and dimensions of α = β = 90°, γ =123.49°, a = 6.69, b = 7.39, and c = 123.49 (α, β, γ, a, b, and c are the relevant cell parameters) was selected. Firstly, the cell parameters were modified and orthorhombicized (α = β = γ =90°, a = 5.58, b = 7.39, and c = 123.49), and all water molecules and hydrogen atoms in the hydroxyl (–OH) groups in the orthorhombic unit cell were removed. A 2 × 3 × 1 supercell was then created along the a, b, and c axes, resulting in an orthorhombic supercell that was used as the initial model. Si–O–Si linkages were randomly broken in the orthorhombic supercell, satisfying the Qn distribution (the number of bridging oxygens contained in the silicone tetrahedra) to obtain a model with a Ca/Si ratio of 1.62. With this model as a basis, a certain number of Ca atoms were randomly removed to obtain four different orthorhombic C–S–H cells with Ca/Si ratios of 1.62, 1.42, 1.22, and 1.02, as shown in Fig. S3.
FA is a type of volcanic ash material with a predominantly amorphous aluminum silicate structure that serves as an auxiliary cementitious material. As shown in Fig. S1(b), 50%–90% of the FA structure consists of an amorphous phase (glass), and the remaining phases are crystalline phases of quartz and mullite [25–27]. MD simulation is a powerful tool for studying the microscopic structure and properties of materials and has been widely implemented in the analysis of the structure and physicochemical properties of phase-change energy storage molten salts and silicate melts. In most studies, the suggested procedure for constructing glass atomic structures is the melt-quench method [21,28]. Similarly, in the present work, the initial cell parameters of Anorthite mineral (Ca8Si16Al6O64) was modified (α = 93.113°, β = 115.913°, γ =91.261°, a = 8.173, b = 12.869, and c = 14.165) to match the chemical composition of FA. The modified molecular formula was Ca4Si24Al8O64, which is the initial structure of FA, i.e., 4CaO·24SiO2·4Al2O3. The initial structure was melted and quenched using the Born–Mayer–Huggins (BMH) potential [29] to obtain an amorphous structure. The BMH potential is commonly used to describe the conformation and characteristics of several silicate systems; its correctness has been widely validated and was hence applied in this study [30], and the expression for the BMH potential function is given in Eq. (1).
Uij(r)=qiqjrij+Aijexp(−Bij⋅rij)−Cijr6l
(1)
where Uij(r) represents the interatomic potential between atomic pairs; qi and qj are the effective charges of atomic pairs i and j, respectively; Aij and Bij are the repulsive constants between the i-th and j-th pairs; rij is a parameter that describes the distance between the i-th and j-th pairs; Cij is the van der Waals potential parameter; rlis the side length of the simulated box, nm. Table 3 lists the potential energy parameters between different atoms [31].
Table
3.
BMH potential parameters for atomic pairs in CaO–Al2O3–SiO2 melt
The modified crystal structure of the Anorthite mineral was first heated from 300 to 5000 K at a rate of 50 K/ps under the NPT (constant-pressure, constant-temperature) ensemble and then cooled to 1873 K at the same rate, followed by 60000 steps of relaxation at 1873 K to obtain the FA model used in this simulation. Figs. S4 and S5 show the modeling process.
The final FA-CSH combination model was matched with different Ca/Si ratios and FA dosages according to Table 4. The grand canonical Monte Carlo method [32–33] was employed to simulate the adsorption of water molecules at saturation under standard atmospheric pressure. The simulation was conducted at a temperature of 300 K, with a chemical potential (μ) value of 0.5 eV. After water adsorption, the system was equilibrated under the isothermal-isobaric NPT ensemble at 300 K. The combination model was then extended in the x, y, and z directions by a 3 × 1 × 2 periodicity to obtain stable results, following which it was relaxed again at 300 K until an equilibrium state was reached. Fig. 1 shows the initial state of all molecular models in the stretching simulations. Stretching simulations were performed using the Lammps calculation program, with the CLAYFF force field and the 12-6L-J potential and Coulombic potential as the interatomic interaction potentials. For the elimination of the boundary effects of the simulation system, the simulation was conducted under periodic boundary conditions with a cutoff radius of 8.0 for the 12-6L-J potential and 10.0 for the Coulombic potential. The long-range Coulombic interactions were calculated using the Ewald summation method. The time step is set as 0.1 fs, followed by equilibration under the NPT ensemble. After energy minimization, the models were stretched under the NPT ensemble at a strain rate of 0.08 ps−1. The simulation lasted for 1000 ps, with dynamic data collected every 0.1 ps and the dynamic trajectory of all atomic configurations obtained during the simulation process. Each model was simulated three times to reduce errors and ensure the accuracy of the simulation results.
Table
4.
FA-CSH molecular model construction scheme
3.1.1
Effect of FA dosage on the compressive strength of CTB
In Fig. 2, the curing age exhibited a considerable impact on the compressive strength of CTB. As the curing time progressed, a marked increase was observed in compressive strength. Within the same curing period, the compressive strength of CTB initially exhibited an upward trend, followed by a subsequent decline as the FA dosage increased. This observation suggests that the incorporation of an optimal amount of FA can substantially enhance the compressive strength of CTB. Moreover, an FA dosage of 60 g considerably diminished the compressive strength of CTB at 3-d and 28-d intervals. This finding implies that excessive FA dosage adversely affects the early and late development of compressive strength, rendering it less favorable for achieving optimal performance.
Fig.
2.
Effects of FA dosage at different dasages (a) and curing time with different FA dosages (b) on the compressive strength of CTB. mfa refers to the dosage of fly ash.
In Fig. 3(a), XRD was conducted on group B specimens at 3, 7, 14, and 28 d. The hydration products exhibited similar mineral phases for the different ages of the samples, with only the intensity of diffraction peaks showing differences in different ages. The mineral phases consisted of ettringite (Aft), calcium hydroxide (Ca(OH)2), calcite (CaCO3), and incompletely reacted dicalcium silicate (2CaO·SiO2, C2S), and tricalcium silicate (3CaO·SiO2, C3S). A “flat convex” was observed in the 2θ range of 18°–20° and 33°–35°, indicating the presence of a large amount of low-crystallinity or amorphous substances in the hydration products generally believed to be amorphous C–S–H gel. Fig. 3(b) shows the XRD results of CTB specimens from groups A, B, C, and D. Fig. 3(c) displays that the diffraction peak of Ca(OH)2 in the hydration products gradually became blunt with the increase in curing age, and its crystallinity decreased gradually. This finding indicates the continuous hydration of C2S and C3S, generating C–S–H gel and Ca(OH)2. Ca(OH)2, which serves as the hydration product, also reacted with the FA in CTB as the reactant of the secondary hydration reaction to form the C–S–H gel. The weak diffraction peak of the amorphous C–S–H gel was enlarged in Fig. 3(e). The appearance of CaCO3 was attributed to the carbonation during sample curing.
Fig.
3.
XRD of CTBs: (a) XRD of CTBs at different ages; (b) XRD of CTBs with different FA dosages; (c) Ca(OH)2 characteristic peak intensity in (a); (d) Ca(OH)2 characteristic peak intensity in (b); (e) C–S–H characteristic peak intensity in (a); (f) C–S–H characteristic peak intensity in (b).
Fig. 3(d) and (f) demonstrates that a proper increase in the dosage of FA led to a gradual decrease in the peak value of Ca(OH)2 and a sharp C–S–H peak. Excessive FA dosage resulted in the smoothening of the C–S–H peak. These results indicate that the overconsumption of Ca(OH)2 can be attributed to reactions between the active components of SiO2 and Al2O3 in FA and cement and Ca(OH)2 in aqueous solution, which leads to the formation of C–S–H. However, excessive FA dosage is not conducive to the formation of these products, which is consistent with the macroscopic experimental results and explains the decline in compressive strength at 28 d when the dosage reached 60 g.
3.1.3
SEM microstructure analysis
As shown in Fig. 4, the Ca/Si ratio of the hydration products surrounding the FA gradually decreased with the increase in FA dosage. This phenomenon can be attributed to the competition between FA and Ca(OH)2 for surrounding Ca2+, leading to the generation of C–S–H with a low Ca/Si ratio. The microstructure of the CTB specimen with low FA dosage in Fig. 4(a) exhibits abundant hydration products. Defects such as micropores and microcracks are closely filled by hydration products, indicating high strength performance. As shown in Fig. 4(b), the Ca/Si ratio continued to decrease with the increase in FA dosage. In Fig. 4(b) and (c), numerous microcracks and microporosity distributions can be observed, and the formed defects are difficult to fill by the hydration products, thus presenting a loose and porous microstructure feature. Furthermore, Fig. 4(c) reveals the presence of a small amount of spherical FA particles in the high FA dosage area, indicating that the excess FA particles have not yet participated in the hydration reaction. This observation suggests that the excessive dosage of FA can adversely affect the production of hydration products and ultimately degrade the strength performance of the CTB specimen.
Fig.
4.
SEM–EDS of CTBs with different FA dosages: (a) low FA dosage area; (b) medium FA dosage area; (c) high FA dosage area.
The reason for the early-stage deterioration caused by excessive FA dosage in 3-day-old specimens cannot be explained by microscopic experiments alone. Therefore, MD simulations were employed to analyze the consumption of Ca2+ by FA in the surrounding environment. Four C–S–H matrix models with Ca/Si ratios of 1.62, 1.42, 1.22, and 1.02 were established to simulate different Ca/Si ratio environments. Each Ca/Si ratio was doped with FA at mass fractions of 0, 22%, 44%, and 66%, and stretching simulations were conducted in the z direction to analyze the stretching results under different Ca/Si ratios and FA dosage conditions during the early reaction stage. The group classes in Fig. S6 were selected to discuss the two causes separately. However, the model corresponding to a Ca/Si ratio of 1.02 cannot truly simulate the reaction during the late stage competing for a large amount of Ca2+ because the tensile strength in the later stage is determined by the amount of C–S–H generated. Meanwhile, the amount of hydrated products constructed for all models in this study was approximately the same.
3.2.1
Analysis of stress–strain curve
The fracture behavior and tensile strength of the FA-CSH model were considerably impacted by the Ca/Si ratio and FA dosage, as indicated by the stress–strain curves obtained from tensile simulations (as depicted in Fig. 5). The FA-CSH model with low FA dosage and Ca/Si ratios of 1.02 and 1.22 showed a fracture mode of plastic flow failure, which transitioned to brittle tensile failure with the increase in FA dosage. Conversely, when the Ca/Si ratio was 1.42 and 1.62, the fracture mode exhibited brittle tensile failure regardless of the FA dosage. The tensile strength of the FA-CSH model did not exhibit a considerable improvement when the FA dosage was increased from 44wt% to 66wt%, except for when the Ca/Si ratio was 1.02 and 1.22. Therefore, the Ca/Si ratio has a considerable influence on the fracture mode of the FA-CSH model when the FA dosage is low. This may be attributed to the transition from the C–S–H(I) to C–S–H(II) structure. Moreover, a Ca/Si ratio of 1.5 is the critical value for these two structures [34]. At a constant FA dosage, the pores within the C–S–H structure at different Ca/Si ratios possessed varying amounts of adsorbed water molecules under identical conditions. During water absorption, the removal of Ca atoms in the dry state created vacancies filled by water molecules. Therefore, the lower the Ca/Si ratio, the greater the number of adsorbed water molecules. The water molecules in the z-direction weakening the interaction between silicon and oxygen atoms may account for the aforementioned changes.
Fig.
5.
Stress–strain of FA-CSH model in axial tension under different Ca/Si ratios: (a) 1.02; (b) 1.22; (c) 1.42; (d) 1.62.
Fig. 6 shows the impact of FA dosage and Ca/Si ratio on water absorption and illustrates that FA doping exerted a considerable impact on the number of water molecules adsorbed by the FA-CSH molecular model in dry state. A reduction in the amount of water absorbed by the material may lead to a tendency toward brittle damage and substantial premature tensile deformation. Consequently, the peak strength of the model was primarily governed by the Ca/Si ratio and FA dosage. Nevertheless, these results do not imply that the interaction forces between water molecules have no impact on determining the tensile strength of the material.
Fig.
6.
Number of H2O molecules (NH2O) in FA-CSH under different Ca/Si ratios and FA dosages.
As shown in Fig. 7(a), the tensile strength of the FA-CSH in molecular model was positively correlated with the Ca/Si ratio for a certain amount of FA dosage, indicating that the Ca in the skeleton can bear most of the load and the load-bearing capacity in the tensile process increases with the Ca/Si ratio. Fig. 7(b) reveals a nonlinear relationship between FA dosage and tensile strength. Under Ca/Si ratios of 1.62 and 1.42, the tensile strength of the FA-CSH in molecular model increased with the FA dosage from 0 to 44wt% but decreased with a further increase in the FA dosage up to 66wt%. Similarly, under a Ca/Si ratio of 1.02, the tensile strength decreased when the FA dosage increased from 22wt% to 44wt%.
Fig.
7.
Tensile strength in FA-CSH molecular model: (a) influence of Ca/Si ratio; (b) influence of FA dosage.
The early stage with water absorption can be characterized by low Ca2+ consumption and few hydration products, which are consistent with macroscopic test results and explains the early strength deterioration caused by high FA dosage. The water molecules in the system also contribute to tensile strength, as demonstrated by the resistance to deformation. Macroscopic experiments indicate that the mechanical properties of samples in group D deteriorated due to excessive FA dosage. High FA dosage regions were observed at the microscale, elucidating the observed deterioration of early strength. Therefore, water molecules contribute to tensile strength by resisting deformation to a certain extent.
3.2.3
Visual analysis of the failure mode
Fig. 8 shows the relationship between the stretching time and model height and reveals the excellent ductile properties of FA under low FA and low Ca/Si ratio conditions. Furthermore, Fig. 9 demonstrates that models with different FA dosages exhibited different failure modes at the same stretching time for a Ca/Si ratio of 1.42. With the addition of Ca and Al to the original water molecular chain, the tensile skeleton developed Al–O–Ca with stronger bonding energy than hydrogen bonding, resulting in increased stiffness and great deformation, leading to early cracking. After setting the hydrogen bond formation distance to 2.5 Å [35], local magnification and cloud diagrams show the hydrogen bond skeleton between water molecules, which limits the development of cracks during stretching.
Fig.
8.
Failure models of FA-CSH with different FA dosages and Ca/Si ratios.
Fig.
9.
Strain field of FA-CSH with a Ca/Si ratio of 1.42 during the failure process corresponding to the failure mode: (a) FA dosage of 0; (b) FA dosage of 22wt%.
In the local strain cloud map, the strain value represents the strain of each individual atom. For the observation of the overall deformation or distortion of the FA-CSH molecular model, the von Mises strain can be calculated by determining the linear strains in various directions of the atoms [36]. Compared to the individual atom strain that is specific to that particular atom and describes its unique strain state, the von Mises strain takes into account the overall strain situation of the material by considering the strain state of all atoms in the material. The strain tensor was derived using Ovito software, and εij was defined as the strain component in six directions for each atom. The von Mises strain εMises can then be calculated using Eq. (2).
Fig. 10 shows the current time Mises strain cloud map corresponding to the stress–strain curve for a Ca/Si ratio of 1.22 and an FA dosage of 22wt% and demonstrates that Mises strain can be used to describe the overall deformation behavior of the material.
Fig.
10.
Stress–strain and atoms colored by Mises strain of FA-CSH with a Ca/Si ratio of 1.22 and an FA dosage of 22wt%: (a) strain of 0.2; (b) strain of 0.4; (c) strain of 0.6; (d) strain of 0.7.
Fig. 11 presents the stretched states and Mises strain cloud maps of FA-CSH molecular models of a Ca/Si ratio of 1.02 with varying FA dosages at a strain of 0.3. The images demonstrate that an appropriate amount of FA can effectively enhance the strength of silica chains in C–S–H and reduce cracking in the z-direction. At an FA dosage of 22wt%, the interior atoms were densely packed, and the external ones were loose. The moisture content outside the model was higher than that inside, leading to oblique section cracking under tension. When the dosage of FA increased to 44wt%, the external water molecules were squeezed into the Ca layers on both sides, which led to cracking. The water layer possessed low moisture content, making it prone to fracture when the FA dosage was further increased to 66wt%.
Fig.
11.
Strain field of FA-CSH with a Ca/Si ratio of 1.02 during the failure process corresponding to the failure mode: (a) FA dosage of 0; (b) FA dosage of 22wt%; (c) FA dosage of 44wt%; (d) FA dosage of 66wt%.
Fig. 12 illustrates snapshots of models with different Ca/Si ratios after 3000 simulation steps under the condition of 22wt% FA dosage. When the Ca/Si ratio was increased from 1.02 to 1.42, a decrease was observed in the cracking of the water layer. However, at a Ca/Si ratio of 1.62, a negative effect of Ca2+ reinforcement on the water layer was observed because the reduction in the number of water molecules in the water layer led to early cracking at Ca/Si ratios lower than 1.42.
Fig.
12.
Snapshots of different Ca/Si ratios for an FA dosage of 22wt%: (a) 1.02; (b) 1.22; (c) 1.42; (d) 1.62.
Fig. 13 presents the snapshots of the different models from Fig. 12 as they were stretched to the point of failure. The defects caused by a low Ca/Si ratio were observed to weaken due to the filling of water molecules, resulting in a reduction in the interaction of Si–O. This phenomenon allowed for the easy formation of Ca–O bonds in the water layer, which subsequently led to the formation of a stretching skeleton aided by the hydrogen bonds between water molecules. As illustrated in Fig. 13(a), the FA-CSH model displayed high ductility with a Ca/Si ratio of 1.02. The layered silicon chains were stretched into a belt-like 3D helical structure. However, weak interactions limited the overall tensile strength. As shown in Fig. 13(b), the density of the helical structure increased, thereby enhancing the interaction between the Si–O tetrahedrons when bearing the belt-like chain rooted in the aluminum atoms of the FA, leading to an increase in the overall tensile strength in the FA-CSH molecular model.
Fig.
13.
Failure modes of FA-CSH with an FA dosage of 22wt% under different Ca/Si ratios: (a) 1.02; (b) 1.22; (c) 1.42; (d) 1.62.
Fig. 13(c) and (d) show that at high Ca/Si ratios of 1.42 and 1.62, where the absorption of water molecules was limited, the interaction between Ca and Si was difficult to weaken, and dispersing silicon chains and stretching them into multidimensional structures was challenging. Therefore, although high ductility could not be displayed, the overall tensile strength in the FA-CSH molecular model was enhanced. The hydrogen bonds between water molecules were the weakest link in the stretching skeleton and could not bind strongly with the Si–Ca atoms, resulting in the early cracking of the weakly hydrated layer. Thus, the Ca/Si ratio of 1.22–1.42 is suitable for achieving high ductility and strength.
3.2.4
Radial distribution function
The radial distribution function (RDF) can be used to characterize the structural and dynamic properties of molecular aggregation at the atomic scale [37–39]. The strength of ion hydration can be intuitively characterized by the characteristic peak value and position of the RDF, i.e., the larger the peak value and the smaller the peak position, the stronger the hydration ability; conversely, the smaller the peak value and the larger the peak position, the weaker the hydration ability. Therefore, the spatial correlation between water molecules and Ca2+defects during stretching can be investigated.
In the CLAYFF force field, Ca ions are classified into two types, CaH and CaO. Moreover, Ca2+ corresponding to CaH is referred to as Cw. RDFs for Cw–O* (oxygen atoms in water molecules) and Cw–H* (hydrogen atoms in water molecules) were computed and plotted because of their close proximity to FA dosage and interlayer Ca (Cw). The RDF represents the probability distribution function of finding oxygen or hydrogen atoms within a specific range around the Ca center atom. The combined model of an FA dosage of 44wt% and a Ca/Si ratio of 1.42 was subjected to stretching for 3500 steps until it approached breaking. As shown in Fig. 14(a), the RDF of Cw–O* exhibited two peaks: an initial high peak at radial r = 2.115 Å, where O* atoms were closely arranged due to van der Waals force, electrostatic force, and other effects; and a subsequent low peak at r = 2.692 Å with a small peak value, indicating a reduction in the order of water molecule clusters. The RDF of Cw–H* was unstable because Cw shared the same positive charge as H* despite a distinct high peak being observed at r = 2.456 Å, indicating the formation of a water layer at the fracture point. The presence of water molecules decelerated fracture occurrence. Comparison of the RDF curves of Cw–O* and Cw–H* shows that the first peak of Cw–O* emerged earlier than that of Cw–H*, suggesting that the likelihood of oxygen atoms surrounding Ca2+ was higher than that of hydrogen atoms. Hydrogen atoms were directed away from Ca2+, and oxygen atoms pointed toward Ca2+, enabling the oxygen atoms in water to simply combine with Ca2+ to form hydrated Ca2+. With the increase in distance, the weakening of intermolecular forces led to a reduction in water molecule density and the order of water molecule arrangement, resulting in a gradual decrease in the difference in the RDF values of the two curves.
Fig.
14.
Radial distribution function: (a) model with an FA dosage of 44wt% and a Ca/Si ratio of 1.42; (b) model with different Ca/Si ratios at an FA dosage of 22wt%.
The RDF can also be used to quantify the chemical bond characteristics between atomic pairs by evaluating the probability density of the separation distance function for interacting pairs [40]. Fig. 14(b) shows the RDFs between different ions or atoms in models with different Ca/Si ratios and 22wt% FA dosage. Ca refers to all Ca2+ in the model, O includes the oxygen atoms in the silicate tetrahedrons of C–S–H, and O* refers to oxygen atoms in water molecules. The first peak represents the hydrogen bonds within water molecules and between water molecules, which increases as the Ca/Si ratio decreases. The second peak represents the Si–O bonds in the silicate chains and remains nearly constant. The third peak represents the ion bonds between Ca and water.
Under the same FA dosage, lower Ca/Si ratios resulted in increased absorbed water molecules and strong interaction forces between water molecules. When the Ca/Si ratio was in the range of 1.22 to 1.42, many water molecules can be enriched around each Ca2+ during stretching, which fully utilized the water-retaining capacity and increased the ductility and tensile strength of the material. By contrast, higher Ca/Si ratios caused a decrease in the number of water molecules, resulting in a weakened interaction between Ca and oxygen ions and the inability to exhibit high ductility.
4.
Conclusions
(1) The UCS of CTB initially increased and then decreased with the increasing FA dosage. An appropriate amount of FA can substantially improve CTB, and excessive FA can result in its deterioration.
(2) The FA-CSH model showed that the strength of CTB also increased initially and then decreased with FA dosage. FA reinforced the silica chain of C–S–H to increase the matrix strength. However, this enhancement was weakened by excessive FA dosage and decreased hydrogen bonding among water molecules.
(3) Optimal CTB ductility and strength were obtained at an FA dosage of 22wt% and a Ca/Si ratio of 1.22–1.42. Low Ca/Si ratios resulted in decreased ionic bonds and strength, and high ratios led to decreased ductility.
(4) Water absorption was negatively correlated with the FA dosage and Ca/Si ratio. Optimal amounts of FA and Ca/Si ratio resulted in proper water absorption and a helical 3D structure with high ductility and strength, exhibiting plastic flow failure.
Acknowledgements
This work was financially supported by the National Natural Science Foundation of China (Nos. 52004272, 52122404, 52061135111, 52174092, and 52074259), the Natural Science Foundation of Jiangsu Province, China (Nos. BK20200660 and BK20220157), the Xuzhou Science and Technology Project, China (Nos. KC22005 and KC21033), the Open Foundation of Shandong Key Laboratory of Mining Disaster Prevention and Control, China (No. SMDPC202104).
Conflict of Interest
Jiangyu Wu is a youth editorial board member for this journal and was not involved in the editorial review or the decision to publish this article. All authors confirm that they have no competing interests or financial ties that could influence the outcomes or interpretation of this research.
A.X. Wu, Y. Wang, H.J. Wang, S.H. Yin, and X.X. Miao, Coupled effects of cement type and water quality on the properties of cemented paste backfill, Int. J. Miner. Process., 143(2015), p. 65. DOI: 10.1016/j.minpro.2015.09.004
[2]
L. Liu, J. Xin, C. Huan, Y.J. Zhao, et al., Effect of curing time on the mesoscopic parameters of cemented paste backfill simulated using the particle flow code technique, Int. J. Miner. Metall. Mater., 28(2021), No. 4, p. 590. DOI: 10.1007/s12613-020-2007-2
[3]
A.A. Wang, S. Cao, and E. Yilmaz, Effect of height to diameter ratio on dynamic characteristics of cemented tailings backfills with fiber reinforcement through impact loading, Constr. Build. Mater., 322(2022), art. No. 126448. DOI: 10.1016/j.conbuildmat.2022.126448
[4]
H.Z. Jiao, W.L. Chen, A.X. Wu, et al., Flocculated unclassified tailings settling efficiency improvement by particle collision optimization in the feedwell, Int. J. Miner. Metall. Mater., 29(2022), No. 12, p. 2126. DOI: 10.1007/s12613-021-2402-3
[5]
H.Z. Jiao, W.B. Yang, Z.E. Ruan, J.X. Yu, J.H. Liu, and Y.X. Yang, Microscale mechanism of tailing thickening in metal mines, Int. J. Miner. Metall. Mater., 30(2023), No. 8, p. 1538. DOI: 10.1007/s12613-022-2587-0
[6]
B.X. Yan, H.W. Jia, E. Yilmaz, X.P. Lai, P.F. Shan, and C. Hou, Numerical study on microscale and macroscale strength behaviors of hardening cemented paste backfill, Constr. Build. Mater., 321(2022), art. No. 126327. DOI: 10.1016/j.conbuildmat.2022.126327
[7]
A. Gladwin Alex, A. Kedir, and T. Gebrehiwet Tewele, Review on effects of graphene oxide on mechanical and microstructure of cement-based materials, Constr. Build. Mater., 360(2022), art. No. 129609. DOI: 10.1016/j.conbuildmat.2022.129609
[8]
J.Y. Wu, H.W. Jing, Y. Gao, Q.B. Meng, Q. Yin, and Y. Du, Effects of carbon nanotube dosage and aggregate size distribution on mechanical property and microstructure of cemented rockfill, Cem. Concr. Compos., 127(2022), art. No. 104408. DOI: 10.1016/j.cemconcomp.2022.104408
[9]
J.Z. Chen, Y. Zhang, D.S. Hou, J. Yu, T.J. Zhao, and B. Yin, Experiment and molecular dynamics study on the mechanism for hydrophobic impregnation in cement-based materials: A case of octadecane carboxylic acid, Constr. Build. Mater., 229(2019), art. No. 116871. DOI: 10.1016/j.conbuildmat.2019.116871
[10]
S. Cao, E. Yilmaz, Z.Y. Yin, G.L. Xue, W.D. Song, and L.J. Sun, CT scanning of internal crack mechanism and strength behavior of cement-fiber-tailings matrix composites, Cem. Concr. Compos., 116(2021), art. No. 103865. DOI: 10.1016/j.cemconcomp.2020.103865
[11]
J.Y. Wu, H.S. Wong, Q. Yin, and D. Ma, Effects of aggregate strength and mass fraction on mesoscopic fracture characteristics of cemented rockfill from gangue as recycled aggregate, Compos. Struct., 311(2023), art. No. 116851. DOI: 10.1016/j.compstruct.2023.116851
[12]
C.L. Hu and Z.J. Li, Property investigation of individual phases in cementitious composites containing silica fume and fly ash, Cem. Concr. Compos., 57(2015), p. 17. DOI: 10.1016/j.cemconcomp.2014.11.011
[13]
T. Hemalatha and S. Sasmal, Early-age strength development in fly ash blended cement composites: Investigation through chemical activation, Mag. Concr. Res., 71(2019), No. 5, p. 260. DOI: 10.1680/jmacr.17.00336
[14]
D.K. Nayak, P.P. Abhilash, R. Singh, R. Kumar, and V. Kumar, Fly ash for sustainable construction: A review of fly ash concrete and its beneficial use case studies, Cleaner Mater., 6(2022), art. No. 100143. DOI: 10.1016/j.clema.2022.100143
[15]
S.A. Rodger and G.W. Groves, Electron microscopy study of ordinary Portland cement and ordinary Portland cement-pulverized fuel ash blended pastes, J. Am. Ceram. Soc., 72(1989), No. 6, p. 1037. DOI: 10.1111/j.1151-2916.1989.tb06265.x
[16]
M.C.G. Juenger and R. Siddique, Recent advances in understanding the role of supplementary cementitious materials in concrete, Cem. Concr. Res., 78(2015), p. 71. DOI: 10.1016/j.cemconres.2015.03.018
[17]
J. Skibsted and R. Snellings, Reactivity of supplementary cementitious materials (SCMs) in cement blends, Cem. Concr. Res., 124(2019), art. No. 105799. DOI: 10.1016/j.cemconres.2019.105799
[18]
E. Sakai, S. Miyahara, S. Ohsawa, S.H. Lee, and M. Daimon, Hydration of fly ash cement, Cem. Concr. Res., 35(2005), No. 6, p. 1135. DOI: 10.1016/j.cemconres.2004.09.008
[19]
Q. Ji, R.J.M. Pellenq, and K.J. van Vliet, Comparison of computational water models for simulation of calcium-silicate-hydrate, Comput. Mater. Sci., 53(2012), No. 1, p. 234. DOI: 10.1016/j.commatsci.2011.08.024
[20]
C.C. Qi, H. Manzano, D. Spagnoli, Q.S. Chen, and A. Fourie, Initial hydration process of calcium silicates in Portland cement: A comprehensive comparison from molecular dynamics simulations, Cem. Concr. Res., 149(2021), art. No. 106576. DOI: 10.1016/j.cemconres.2021.106576
[21]
M. Shishehbor, D. Sakaniwa, D. Stefaniuk, K.J. Krakowiak, and M.J. Abdolhosseini Qomi, On the significance of interfacial chemistry on the strength of fly ash-cement composites, Cem. Concr. Res., 151(2022), art. No. 106619. DOI: 10.1016/j.cemconres.2021.106619
[22]
H.J.C. Berendsen, J.R. Grigera, and T.P. Straatsma, The missing term in effective pair potentials, J. Phys. Chem., 91(1987), No. 24, p. 6269. DOI: 10.1021/j100308a038
[23]
R.T. Cygan, J.J. Liang, and A.G. Kalinichev, Molecular models of hydroxide, oxyhydroxide, and clay phases and the development of a general force field, J. Phys. Chem. B, 108(2004), No. 4, p. 1255. DOI: 10.1021/jp0363287
[24]
R.J.M. Pellenq, A. Kushima, R. Shahsavari, et al., A realistic molecular model of cement hydrates, Proc. Natl. Acad. Sci. U. S. A., 106(2009), No. 38, p. 16102. DOI: 10.1073/pnas.0902180106
[25]
S.W. Tang, H.B. A, J.T. Chen, et al., The interactions between water molecules and C–S–H surfaces in loads-induced nanopores: A molecular dynamics study, Appl. Surf. Sci., 496(2019), art. No. 143744. DOI: 10.1016/j.apsusc.2019.143744
[26]
P. Ganster, M. Benoit, W. Kob, and J.M. Delaye, Structural properties of a calcium aluminosilicate glass from molecular-dynamics simulations: A finite size effects study, J. Chem. Phys., 120(2004), No. 21, p. 10172. DOI: 10.1063/1.1724815
[27]
Y.Q. Wu, J.L. You, and G.C. Jiang, Molecular dynamics study of the structure of calcium aluminate melts, J. Inorg. Mater., 18(2003), No. 3, p. 619.
[28]
A. Pedone, Properties calculations of silica-based glasses by atomistic simulations techniques: A review, J. Phys. Chem. C, 113(2009), No. 49, p. 20773. DOI: 10.1021/jp9071263
[29]
R.A. Buckingham. The classical equation of state of gaseous helium, neon and argon, Proc. R. Soc. A, 168(1938), No. 933, p. 264.
[30]
Z.J. Wang, Z.H. Li, M. Zhong, Z.S. Li, and C. Wang, Elucidating the effect of Al2O3/SiO2 mass ratio upon SiO2–MnO–CaF2–Al2O3-based welding fluxes: Structural analysis and thermodynamic evaluation, J. Non Cryst. Solids, 601(2023), art. No. 122071. DOI: 10.1016/j.jnoncrysol.2022.122071
[31]
J.M.D. Lane, Cooling rate and stress relaxation in silica melts and glasses via microsecond molecular dynamics, Phys. Rev. E, 92(2015), No. 1, art. No. 012320. DOI: 10.1103/PhysRevE.92.012320
[32]
D.J. Adams, Grand canonical ensemble Monte Carlo for a Lennard-Jones fluid, Mol. Phys., 29(1975), No. 1, p. 307. DOI: 10.1080/00268977500100221
[33]
J.D. Gale, GULP: A computer program for the symmetry-adapted simulation of solids, J. Chem. Soc., Faraday Trans., 93(1997), No. 4, p. 629. DOI: 10.1039/a606455h
[34]
A. Nonat, The structure and stoichiometry of C–S–H, Cem. Concr. Res., 34(2004), No. 9, p. 1521. DOI: 10.1016/j.cemconres.2004.04.035
[35]
P.A. Sigala, E.A. Ruben, C.W. Liu, et al., Determination of hydrogen bond structure in water versus aprotic environments to test the relationship between length and stability, J. Am. Chem. Soc., 137(2015), No. 17, p. 5730. DOI: 10.1021/ja512980h
[36]
F. Shimizu, S. Ogata, and J. Li, Theory of shear banding in metallic glasses and molecular dynamics calculations, Mater. Trans., 48(2007), No. 11, p. 2923. DOI: 10.2320/matertrans.MJ200769
[37]
A.L. Thorneywork, R. Roth, D.G.A.L. Aarts, and R.P.A. Dullens, Communication: Radial distribution functions in a two-dimensional binary colloidal hard sphere system, J. Chem. Phys., 140(2014), No. 16, art. No. 161106. DOI: 10.1063/1.4872365
[38]
M.P. Pomiès, N. Lequeux, and P. Boch, Speciation of cadmium in cement: Part I. Cd2+ uptake by C–S–H, Cem. Concr. Res., 31(2001), No. 4, p. 563. DOI: 10.1016/S0008-8846(00)00480-4
[39]
Q. Luo and J.L. Huang, Mechanisms and critical technologies of transport inhibitor agent (TIA) throughout C–S–H nano-channels, Materials, 15(2022), No. 2, art. No. 515. DOI: 10.3390/ma15020515
[40]
B.A. Steele, N. Goldman, I.F. Kuo, and M.P. Kroonblawd, Mechanochemical synthesis of glycine oligomers in a virtual rotational diamond anvil cell, Chem. Sci., 11(2020), No. 30, p. 7760. DOI: 10.1039/D0SC00755B
Xinjie Li, Yingfang Fan, Qiuchao Li, et al. Experimental study on early-age fracture behavior of cement mortar with the addition of fly ash. Construction and Building Materials, 2025, 465: 140255.
DOI:10.1016/j.conbuildmat.2025.140255
2.
Xuebing Zhang, Yang Quan, Guohui Cao, et al. Influence of waste glass on concrete strength and permeability during dry-wet cycles. European Journal of Environmental and Civil Engineering, 2025, 29(1): 21.
DOI:10.1080/19648189.2024.2365960
3.
Jiang Li, Shenggen Cao, Yang Liu, et al. Effects of Different Ratios of Fly Ash and River Sand on Mechanical Properties of Cemented Paste Backfill Materials. Advances in Civil Engineering, 2024, 2024(1)
DOI:10.1155/2024/7529543
4.
Shengtang Zhang, Shurong Chen, Lingran Min, et al. Study on the evolution of mechanical properties of backfill body under the combined action of sulfate erosion and load. Case Studies in Construction Materials, 2024, 21: e03977.
DOI:10.1016/j.cscm.2024.e03977
5.
Shushuai Wang, Renshu Yang, Yongliang Li, et al. Effects of cement content, polypropylene fiber length and dosage on fluidity and mechanical properties of fiber-toughened cemented aeolian sand backfill. International Journal of Minerals, Metallurgy and Materials, 2024, 31(11): 2404.
DOI:10.1007/s12613-024-2885-9
6.
Pengfei Wu, Xinyue Liu, Xiaoming Liu, et al. Effect of Industrial Byproduct Gypsum on the Mechanical Properties and Stabilization of Hazardous Elements of Cementitious Materials: A Review. Materials, 2024, 17(17): 4183.
DOI:10.3390/ma17174183
7.
Chao Wei, Zhiqing Xie, Jiarui Gu, et al. High-volume utilization of circulating fluidized bed fly ash for the production of autoclaved aerated concrete: Performance optimization and hydration characteristics. Construction and Building Materials, 2024, 448: 138305.
DOI:10.1016/j.conbuildmat.2024.138305
8.
Hongsheng Wang, Dengfeng Chen, Ruihong Guo, et al. A Preliminary Study on the Improvement of Gangue/Tailing Cemented Fill by Bentonite: Flow Properties, Mechanical Properties and Permeability. Materials, 2023, 16(20): 6802.
DOI:10.3390/ma16206802
9.
Weiqiang Chen, Shuaijie Lu, Shuyang Yu, et al. Effects of graphene oxide on shearing performance of C–S–H composites: a molecular dynamics study. Journal of Materials Science, 2023, 58(44): 16972.
DOI:10.1007/s10853-023-09101-2
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