Molecular dynamics study on calcium aluminosilicate hydrate at elevated temperatures: Structure, dynamics and mechanical properties

Materials Chemistry and Physics 233 (2019) 276–287

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Molecular dynamics study on calcium aluminosilicate hydrate at elevated temperatures: Structure, dynamics and mechanical properties Jianhua Zhang a, Jun Yang b, Dongshuai Hou c, *, Qingjun Ding b a

College of Aerospace and Civil Engineering, Harbin Engineering University, Harbin, China School of Materials Science and Engineering, Wuhan University of Technology, No. 122, Luoshi Road, Hongshan District, Wuhan, China, Postcode: 430070 c Department of Civil Engineering, Qingdao Technological University, No. 11, Fushun Road, Shibei District, Qingdao, China, Postcode: 266033 b

H I G H L I G H T S

G R A P H I C A L A B S T R A C T

� The interlayer region of CASH gel is expanded at high temperature. � High temperature destroys the H-bond network of interlayer water. � The chemical strength of Al–O bonds degrades at high temperature. � High temperature accelerates the depolymerization and hydrolytic reaction for CASH. � The strength and stiffness of the CASH gel is weakened at high temperature.

A R T I C L E I N F O

A B S T R A C T

Keywords: H-bond network Interlayer water dynamics Aluminosilicate skeleton Inhomogeneous calcium dynamics Hydrolytic reaction pathway

In order to gain a molecular-level insight on the structure and performance of cement-based materials in high temperature environment, this paper investigates the structure, dynamics and mechanical properties of their main hydration product, calcium aluminosilicate hydrate (C-A-S-H), at elevated temperatures by using reactive molecular dynamics simulation. The results show that rising temperature can destroy the H-bond network of interlayer water molecules in C-A-S-H, leading to pronounced expansion of the interlayer regions. Meanwhile, interlayer water molecules lose its glassy water dynamics and exhibit dramatically high diffusivity with rising temperature. In addition, the calcium atoms show inhomogeneous dynamics at high temperature. At 1500K, the interlayer calcium atoms can escape from their coordination “cages” and diffuse, while the motion of those located in the principal layer are restricted by the Si–O and Al–O bonds throughout the simulation. Furthermore, reactive force field couples the mechanical response and chemical reaction during the uniaxial tensile test for CA-S-H gel. The de-polymerization and hydrolytic reaction happens frequently in the deformed C-A-S-H gel at high temperature, resulting in degradation of stiffness and strength. On the other hand, the atomic rearrangement at elevated temperature contributes to the re-connection of broken chemical bonds, enhancing ductility of C-A-S-H.

1. Introduction Concrete, ubiquitously utilized all over the world, is the essential

construction material for infrastructures. The cement manufacture re­ sults in approximately 6–8% of the yearly man-made global CO2 emis­ sions [1]. With the increasing demand of concrete materials, the amount

* Corresponding author. E-mail addresses: [email protected] (J. Zhang), [email protected] (J. Yang), [email protected] (D. Hou), [email protected] (Q. Ding). https://doi.org/10.1016/j.matchemphys.2019.05.020 Received 25 March 2019; Received in revised form 10 May 2019; Accepted 14 May 2019 Available online 16 May 2019 0254-0584/© 2019 Elsevier B.V. All rights reserved.

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of carbon emission is still increasing [2]. In order to lower the envi­ ronmental impact of the construction industry, industrial wastes or by-products are utilized to substitute the cement and alternative binders are developed [3]. The wide usage of some supplementary cementitious materials (SCMs), such as fly ash and blast furnace slag [4], has not only reduced CO2 emission [5,6], but also improve the workability and durability of cement-based materials, such as resistance to sulfate attack [7] or alkali aggregate reaction (ASR) [8,9]. It should be noted that SCM, including the active silicate and aluminate minerals, often bring about aluminate species in the cement-based material [4]. The incorporation of Al changes the stoichiometry of calcium silicate hydrate (C–S–H) gel, the main hydration product of cement-based materials. The Ca/Si ratio of the C–S–H gel is reduced and calcium aluminosilicate hydrate (C-A-S-H) is produced [4]. The nanostructure and composition change induced by the Al incorporation can influence the properties of cement-based materials. As a consequence, many studies have been devoted to decode the molecular structure of C-A-S-H gel. So far, the molecular representation of the C-AS-H gel has been elucidated by the analogue based on the aluminumcontaining tobermorite minerals [10,11]. Now it is widely believed that aluminum can substitute the position of the bridging silicate tet­ rahedron in C-A-S-H and it exists in the form of four-/five- and six-coordination [12–14]. It is also found that aluminum atoms can heal the defective silicate chains to form long aluminosilicate chains [12,13]. Since the covalently bonded aluminosilicate chains are the backbone of the C-A-S-H structure, the highly polymerized aluminosilicate chains is believed to enhance the mechanical properties of the C-A-S-H gel [15, 16]. Furthermore, incorporation of aluminum in the interlayer region of C-A-S-H gel can facilitate the connection of aluminosilicate chains in adjacent calcium silicate sheets and transform the layered molecular structure to the cross-linked C-A-S-H [17]. Both experimental and simulation work have exhibited the enhanced stiffness and cohesive strength of cross-linked C-A-S-H as compared to its non-cross-linked counterpart [18,19]. Additionally, the structure and dynamics of water and ions in the gel pore of C-A-S-H has also been studied using molecular dynamics method [20]. It is found that aluminum incorpo­ ration can enhance the polarity of C–S–H, and C-A-S-H shows better hydrophilicity and ion immobilization properties [21]. Even though the nanoscale nature of the C-A-S-H gel has been studied by many years, there are still outstanding questions concerning to the properties of C-A-S-H gel at extremely severely condition, such as high temperatures. It is necessary to comprehensively study the performance of sustainable cement-based materials at elevated temperatures, especially, when they suffer to fire attacking [22,23]. Arioz [24] has studied the effect of elevated temperatures on con­ crete and found that the concrete specimens exhibit dramatic reductions in both weight loss and mechanical strength after exposure to high temperature as high as 1200 � C. Khoury and Ulm et al. [25–27] inves­ tigated the mechanical properties of the concrete sample at different temperatures and found that when temperature ranges from 100 to 200 � C, the strength of concrete sample increases, but its strength de­ creases at still higher temperatures. They also tested the influence of temperature when the concrete samples are sealed. Unexpectedly, sealed concrete sample exhibits lower resistance to high temperature, as compared with the unsealed sample. This is attributed to pronounced transient thermal creep in sealed concrete [25]. Ladaoui et al. [28] have reported that the creep of concrete sample is doubled at 50 � C compared to that at ambient temperature. These studies imply that the destruction of concrete at high temperature arises from the change of the nano­ structure of C–S–H gels. On the other hand, when the temperature in­ creases to 560 � C, geopolymer would transform from a stiff material to a viscoelastic one abruptly (namely, glass transition) [29]. It suggests the aluminosilicate chains may show a different nature from the silicate chains at high temperature. Hence, to interpret the influence of tem­ perature on the SCM blended cement, it is necessary to investigate the properties of C-A-S-H at high temperature.

In this paper, the structure, dynamics and mechanical properties of C-A-S-H gel is studied at different temperatures ranging from 300K up to 1500K, using reactive molecular dynamics (MD) simulation. It should be noted that the simulations do not change the composition of C-A-S-H with temperature for following two reasons: (1) there is no reliable data which matches the dehydration degree of C-A-S-H with the temperature. (2) To avoid the influence of dehydration on the dynamics and reactivity for interlayer water. Furthermore, uniaxial tensile test is utilized to study the mechanical properties of C-A-S-H gel with increasing tem­ perature. The reactive force field can combine the chemical response and mechanical response of the C-A-S-H gel at elevated temperatures. 2. Simulation method Reactive Force field. In past decades the quality of empirical po­ tentials has greatly improved by employing bond order concepts that depend on the local chemical environment in reactive simulations. These bond-order potentials allow for the dissociation and creation of chemical bonds during molecular dynamics simulations and are thus also referred to as reactive potentials. The traditional reactive forces fields include Stillinger-Weber, Tersoff-Brenner, EDIP [30,31] and ABOP [32]. Although both the charge-free and fixed-charge bond-order for­ malisms have demonstrated the flexibility to model many systems, the obvious limitation of these traditional reactive potentials is the lack of a dynamic charge scheme to readjust the atomic charge according to its environment. As compared with the traditional potential, The ReaxFF (reactive force field) method, as originally developed by van Duin and colleagues [33], and the charge-optimized many-body (COMB) poten­ tials developed by Sinnott, Phillpot and coworkers [34] are two well-documented [35] variable-charge reactive potentials that are able to simultaneously treat a variety of elements and multifunctional sys­ tems. As illustrated in Reference [34], the COMB and ReaxFF potentials are comparable in predicting the heat-of-formation energy of a wide range of reaction systems. ReaxFF potentials are more specifically fit to individual transition states and are consequently generally more reliable than COMB at modeling this aspect of chemical reactions. In order to accurate capture the chemical reactions of the C-A-S-H gel at high temperatures, the ReaxFF was selected for the MD simulation. The ReaxFF is initially developed to simulate the hydrocarbons. So far, it has been widely utilized to model various systems, such as Oxide/H2O in­ terfaces [36,37] and catalytic processes [38,39]. The set of parameters for calcium-aluminate-silicate-hydrate systems [40–42] has also been developed and used to study the hydration products and minerals in cement system [16,43–45]. Manzano et al. [44] have calculated the structure and mechanical properties of C–S–H by using ReaxFF and DFT separately. The properties predicted from ReaxFF are comparable with those obtained from the quantum mechanics calculation. It confirms the reliability of the ReaxFF in describing the interactions of atoms in C–S–H gel. Yang et al. [19] have established five cross-linked C-A-S-H molec­ ular models by using ReaxFF force field. The structural parameters are in good agreement with experimental results, including basal spacing, silicate tetrahedral polymerization state and coordination distribution of aluminate species. The study from Manzano et al. [46] has also shown the competence of ReaxFF to charaterize the hydrogen bonds between water molecules. Model Construction. The C-A-S-H model was constructed by following the method proposed in previous research [19,47]. The orthorhombic 11 Å tobermorite [48,49] was taken as the initial struc­ ture of the model. Firstly, the tobermorite crystallographic unit cell was replicated 2, 3 and 1 times along x, y and z dimensions, respectively. Secondly, bridging SiO2 was randomly removed from the tobermorite supercell to generate a Qn (mAl) species distribution of Q1 ¼ 36.07%, Q2 ¼ 63.93%, which is close to the distribution of Q1 ¼ 37.62% and Q2 ¼ 62.38% in C–S–H with Ca/Si ratio ¼ 1.0 [50]. Here, the Qn (mAl) is used to quantify the connectivity of silicate tetrahedra. For a silicate tetrahedron, n denotes the number of neighboring silicate or aluminate 277

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position of atom i; fiI is the component in I direction of the force exerted on the I th atom and mi is the mass of I th atom. In order to investigate the influence of strain rate on the mechanical properties of C-A-S-H gel, two extra tension processes at 300K and 1500K with strain rate of 0.008/ps were conducted. The discussion on the influence of tensile strain rate is provided in supporting information file. Data analysis. To quantitatively evaluate the dynamical properties of C-A-S-H, the trajectories of the atoms are analyzed according to the following functions: Mean squared displacement (MSD), described by equation (2), is used to represent the mobility of atoms. � �2 MSDðtÞ ¼ �ri ðtÞ ri ð0Þ� (2) Where the ri(t) is the position of I th atom at time t. A higher MSD value of the atoms indicates higher mobility of these atoms. Furthermore, the diffusion coefficient (D) can be obtained from the MSD as the function of time in equation (3) [56]: � �2 2nDt ¼ �ri ðtÞ ri ð0Þ� (3)

Fig. 1. Molecular structure of the initial C-A-S-H model; Chemical composition ¼ (CaO)0.98⋅SiO2⋅[(AlO2)-]0.10(Hþ)0.10⋅1.26H2O; Cell size ¼ 21.63 � 22.67 � 25.28 Å3; The same atomic representations are used in the following figures.

polyhedron and m denotes the number of neighboring aluminate poly­ hedron. After the removal of silicate tetrahedra, aluminum atoms were randomly added into the defective sites until the Al/Si ratio gets 0.1. When each Al atom was incorporated, two oxygen atoms and one hydrogen atom were simultaneously added to compose a tetrahedron and to maintain charge neutrality. Thirdly, the interlayer regions in the supercell were slightly expanded and water molecules were inserted into until the H/Si ratio reaches 2.6. The supercell was then relaxed at ab­ solute zero and equilibrated at NPT ensemble (T ¼ 300K and P ¼ 1 atm) for 1 ns Finally, the generated C-A-S-H model has a chemical formula of (CaO)0.98⋅SiO2⋅[(AlO2)-]0.10(Hþ)0.10⋅1.26H2O and interlayer spacing of 12.64 Å, which falls in the range of experimental results [13,50]. Fig. 1 is the atomic configuration of this C-A-S-H model. The size of the model is 21.63 � 22.67 � 25.28 Å3. The molecular model represents the basic building block of the cement hydrate. According to literature [51,52], the intrinsic unit of C-A-S-H molecular structure is around 2–5 nm. To study the influence of temperature on the nanoscale properties of C-A-S-H gel, the newly constructed C-A-S-H model was simulated at 300, 400, 500, 600, 800, 1000, 1300 and 1500K, respectively. The anisotropic pressure was set to 1 atm, 1 atm and 1 GPa along x, y and z directions, respectively. Note that the z directional pressure is set at high value to prevent interlayer water from evaporating at high temperature. This value is close to the cohesive pressure between C-A-S-H nano-particles [53], and thus it can be considered as the constraint from surrounding C-A-S-H particles. The abovementioned MD simulations were performed with REAX [54] package in LAMMPS [55] software, using a time step of 0.25 fs The first 250 ps of MD run were used to equilibrate the system. Subsequently, further 1000ps of production run were performed and the atomic trajectories were recorded for data analysis. Mechanical properties. To obtain the stress-strain relation, the uniaxial tension tests along y direction at different temperatures were performed on the C-A-S-H. Firstly, the C-A-S-H models at different tem­ peratures were taken as the starting configuration. These models were periodically extended three times along x, y and z directions to construct supercells. It should be noted that the supercell of this size is to moderate the system temperature and pressure fluctuation. Secondly, the super­ cells were relaxed for 100ps and the model reached equilibrium at each temperature. Finally, the supercells were tensioned with a strain rate of 0.08/ps. During the elongation, the pressure in x, y and z directions were set as 1 atm, 1 atm and 1 GPa, respectively. The stress tensor component was calculated and modified according to following equation: PN PN riI fiJ i mi viI viJ PIJ ¼ þ i (1) V V

where n represents dimension. Time-correlated function (TCF) is used to estimate the stability of chemical bonds as following equation: CðtÞ ¼

< δbðtÞδbð0Þ > < δbð0Þδbð0Þ >

(4)

Where δb(t) isb(t)-, b(t) is a logical value that gives one if an atomic pair (eg., O–H) is connected and zero otherwise, and is the average value of b throughout the simulation time and all atomic pairs. The self-part of van Hove functions for atoms are calculated by using the following equation [57]: Gs ðr; tÞ ¼

N � 1 X δ r N i¼1

� �ri ðtÞ

�� ri ð0Þ�

(5)

Where N is the number of these atoms and ri(t) is the position of I th atom at time t. This function can define the probability for an atom that has moved by a distance r at time t with respect to its original position. 3. Results and discussion 3.1. Structural evolution at elevated temperatures The layered structural evolution. It can be observed in Fig. 2 that the C-A-S-H exhibits pronounced expansion along z direction at high temperature, and expands slightly along x and y directions. The change

Here, N is the number of atoms in the simulation box and V is the volume of the box; I and J represent generic coordinate directions (x, y and z); viI, and riI are the component in I direction of the velocity and

Fig. 2. The cell size evolution of the C-A-S-H along x, y and z directions at different temperatures. 278

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Fig. 3. Side views of the C-A-S-H molecular structure at (a) 300K, (b) 800K, (c) 1000K and (d) 1500K.

of the structure is mainly attributed to the expansion of interlayer re­ gions concentrated with water molecules. The z-direction size of C-A-SH gel increases by 4.41 Å as the temperature is elevated from 300K to 1500K. It is much higher than the interlayer expansion value (~1.3 Å) obtained from the previous simulation on C–S–H [58]. The discrepancy of the interlayer expansion can be explained by different Ca/Si ratios in current model (Ca/Si ¼ 1.0) and that in previous study (Ca/Si ¼ 1.3) [58]. According to the fluid pressure theory [59], water molecules produce the disjoining pressure in the pore with a diameter less than 5 Å, and the Ca2þ ions play a cohesive role in connecting the neighboring calcium silicate sheets. Due to small number of interlayer calcium atoms in current model, the cohesion between calcium silicate sheets is weaker and hence the interlayer expansion tends to be more pronounced at high temperature. The molecular structures along with their atomic density profiles of the C-A-S-H are given in Figs. 3 and 4, respectively. As shown in Fig. 3a, the C-A-S-H shows ordered ‘sandwich’ structure at 300K. In the calcium silicate layers, calcium atoms (Cs) are distributed parallel to xy plane and aluminosilicate chains are aligned along y direction. Between the calcium silicate sheets, interlayer water molecules and calcium atoms (Cw) are randomly distributed. As shown in the atomic density profile, the alternative intensity maxima of Ca, Si(Al) and O also indicate the ordered layered structure of C-A-S-H (Fig. 4a). As the temperature rises to 800K and 1000K, the structural ordering of C-A-S-H maintains un­ changed (Fig. 3b–c). It can be observed in Fig. 4b–c that the intensities

and widths of the density peaks do not exhibit substantial differences as compared with those at 300K. Nevertheless, when the temperature reaches 1500K, the C-A-S-H structure transforms from layered structure to amorphous state. As shown in Fig. 3d, the pairing silicate tetrahedron move from their original orientations and the Cs atoms also displace from their original planes. The disorder structure is confirmed in Fig. 4d that the density peaks of calcium silicate sheets turn broadening. Furthermore, in Fig. 4d, more fluctuations in the density profiles of Ow (Oxygen atom in water molecule) and Oh (Oxygen atom in hydroxyl) the interlayer water and calcium atoms are more probable to displace frequently as compared with those in the calcium silicate sheets. The simulation results indicates that calcium silicate sheet can maintain its ordered layered structure even at temperature up to 1000K. It is much higher than the experimental results [22,60]. The dehydration of C-A-S-H happening at high temperature is not considered in current simulation. The presence of water molecules in the interlayer region can maintain the layered order structure to resist the attacking at high temperature. Aluminosilicate chains morphology. In order to characterize the aluminosilicate chains morphology, the local structures of Si and Al atoms are firstly analyzed. The radial distribution functions (RDF) of Si–O and Al–O correlations are shown in Fig. 5a. At 300K, the first peak indicates Si–O bond with a distance of 1.632 Å, which is close to the value of 1.64 Å from the experimental work [61]. With increasing temperature, the RDF peak of Si–O exhibits a broadening bond length 279

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Fig. 4. Atomic density profiles of the C-A-S-H at temperatures of (a) 300K, (b) 800K, (c) 1000K and (d) 1500K, where Ow and Oh denote oxygen atoms in water molecules and hydroxyl groups, respectively.

Fig. 5. (a) The radial distribution functions (RDF) of Si–O and Al–O correlations in C-A-S-H at different temperature; (b) Coordination number of aluminum atoms in C-A-S-H at different temperature (Ob: bridging oxygen atoms; Onb: non-bridging oxygen atoms; Total: both the Ob and Onb).

distribution. It indicates that the Si–O bonds are more probably to be stretched at high temperature. The average Si–O bond length slightly increases from 1.632 Å to 1.655 Å. The elevated temperature can also result in the increase of the width of Al–O RDF peak. As shown in Fig. 5a, the average Al–O bond length ranges from 1.873 Å to 1.894 Å. Previous X-ray diffraction (XRD) study indicates that the distances of Al–O bonds range from 1.75 to 1.90 Å and the Al–O bond length is closely related with coordination number (CN) of Al atoms [62]. The coordination distribution is an important parameter to evaluate the local structure of aluminum atoms in C-A-S-H structure. At 300K, there are tetra-/penta-/octa-coordinated aluminate species (AlIV, AlV and AlVI) in the C-A-S-H gel. This agrees reasonably well with previous simulation and experiments [13,15,50]. The CN number of Al atoms is plotted as a function of temperature in Fig. 5b. It can be observed that the Al atoms have an average CN ranging from 5.00 to 5.42 and elevated temperature does not show pronounced influence on the CN. On the other hand, the CN of Ob suddenly increases as the temperature rises to

1500K, implying the formation of more Si–O–Al bonds. It means the enhancing polymerization degree of aluminate polyhedral at extremely high temperature. The Q species distribution and hydroxyls number evolution at different temperatures are exhibited in Fig. 6a–b, respectively. Fig. 6c–e give snapshots of the local structure for the newly formed Q species. As shown in Fig. 6a, the evolution of aluminate structure can be categorized into three stages. Firstly, as the temperature is below 800K, the per­ centages of Q species maintain almost unchanged. Subsequently, as the temperature rises to 1300K, the proportions of Q2 and Q2(1Al) are reduced and that of Q1 increases, indicating the depolymerization of aluminosilicate chains as shown in Fig. 6c. At 1500K, the amount of Q3 and Q4species increases and that of Q1 decreases. It can be observed in Fig. 6d–e that the aluminosilicate branched chains and network are formed at 1500K. It should be noted that the number of hydroxyl groups is strongly correlated with the Q species distribution, as shown in Fig. 6b. When the Q species distribution remains constant, no significant 280

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Fig. 6. The evolution of (a) Q species distribution and (b) hydroxyl groups in C-A-S-H with elevated temperature; Snapshots of the (c) Q1 species, (d) aluminosilicate branched chains and (e) silicate branched chains formed at temperature.

Fig. 7. (a) The radial distribution function (RDF) of O–H correlation and (b) Average Hbond number per water molecule in C-A-S-H at different temperatures, where Ow-d-Ob: water molecule donate proton to form H-bond with bridging oxygen atom; Ow-d-Ow: water molecule donate proton to form H-bond with water molecule; Ow-a-Ow: water molecule accept proton to form H-bond with water molecule; Ow-d-Oh: water molecule donate proton to form H-bond with hydroxyl group; Ow-a-Oh: water molecule accept proton to form H-bond with hydroxyl group.

change happens for the number of water molecules and hydroxyls. While the depolymerization of aluminosilicate chains, water molecules start to dissociate to Si–OH groups. On the other hand, the formation of branched aluminosilicate chains and network is accompanied with reduction of Al–OH groups. It confirms that the branched chains or networks are contributed by the reaction of the polymerization Si–O–Al and water production. Local structures of water molecules. As mentioned above, the interlayer spacing swells at elevated temperature, which can induce the structural change of interlayer water molecules. It can be observed in Fig. 7a that the first RDF peak of O–H located at 0.97 Å slightly shifts toward lower distance with increasing temperature. It means that the O–H bond length reduces at higher temperatures. The shortening of O–H bond has also been reported in previous study [63] that O–H bond length decreases from 0.96 Å at ambient temperature to 0.94 Å at 383K. The experiment does not provide data at further higher temperatures as the interlayer water molecules are evaporated. In addition, the second peak in O–H correlation, denoting the hydrogen bond, gradually weakens with the increase of temperature as shown in Fig. 7a, implying the destroy of the H-bond network in the interlayer region. At 1500K, the second O–H peak overlaps with the third peak, suggesting the mixing

of second coordination shell and third coordination shell for water molecules. The average H-bond number per water molecule is calculated and exhibited in Fig. 7b. In current work, the H-bond is formed according to the following criteria [64]: (1) the distance between H and Y should be less than 2.45 Å. (2) The angle of the H-X-Y should be less than 30� , where the H-bond is written as X–H…Y, X–H is the H-bond donator and Y is the H-bond acceptor. As shown in the figure, the average number of H-bond is 2.60 at 300K, which is close to the results from previous simulations [19,43,65]. This value is lower than the number for bulk water, where one water molecule has about 3.6 H-bonds with adjacent water molecules, including 1.8 proton-acceptor H-bonds and 1.8 proton-donator H-bonds. This is attributed to geometrical confinement that provides the solid oxygen atoms to form H-bond with confined water molecules and hence disturbs the connectivity of the H-bonds. On the other hand, the average number of H-bond per water decreases monotonously as the temperature rises and gets 1.28 at 1500K. Ac­ cording to the types of H-bond donator and acceptor, the H-bonds can be further categorized into five types, and are also listed in Fig. 7b. As compared with the other four H-bond components, the number of Ow-d-Ob reduces significantly. The value reduces from 0.75 to 0.17 as 281

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the temperature rising from 300K to 1500K. It indicates that the H-bond between water molecule and Ob atom is most likely to be broken among the five H-bond types. In this respect, the reduction of H-bond number suggests that elevated temperature has a destructive effect on the H-bond network within C-A-S-H. Due to the weakening of H-bond connection, the calcium silicate aluminate sheets are dispersed in great extent, leading to the interlayer expansion.

degradation of Al–O bonds stability is initiated at a temperature between 800K and 1000K. Since the glass transition of geopolymer appears at the temperature of 833.5 K [29], falling in the temperature range for the initiation of Al–O bonds instability. It implies that the glass transition is attributed to weakening of Al–O bonds. On the other hand, a previous study [66] reported that the TCF value of Al–O bonds decreases to 0.93 when the geopolymer undergoes high temperature of 1500K for 50ps. This value is much higher than that in our C-A-S-H model (TCF ¼ 0.69), indicating the higher stability of Al–O bonds in geopolymer than in C-A-S-H. It is noteworthy that most of the oxygen atoms in the geopolymer are bridging oxygen atoms, which are restricted by the aluminosilicate skeleton. At high temperature, although the Al–O connections are weakened, the oxygen atoms still cannot diffuse freely due to the constraints of neighboring Si and Al atoms. In the C-A-S-H gel, the aluminate polyhedral is connected with interlayer water molecules. Once the Al–O bonds are broken, the hy­ droxyl exchange between Al–OH and interlayer water molecules can happen frequently. Hence, the Al–O bonds in the geopolymer are more stable than their counterparts in the C-A-S-H. Additionally, the dis­ placements of bridging and non-bridging oxygen atoms (Ob and Onb) in Al–O bonds have been analyzed separately. The self-part of van Hove functions for Ob and Onb atoms at 1500K are plotted in Fig. 8c. At t ¼ 5ps, the curve for Ob presents in r range from 0 to 7 Å, indicating that most of the Ob atoms are in the coordination cage. The function for Onb extends to area with r value larger than 15 Å, indicating high mobility of Onb atoms. Furthermore, at t ¼ 50ps, the function for Ob atoms display trinomial distribution, which indicates a small part of Ob atoms are jumped out of their first and second coordination cages. Meanwhile, the displacement of Onb atoms disperses in r range of 0–45 Å and the displacement distribution is broad. It means the coordination cage ex­ hibits weaker confinement on the Onb atoms. The positions of these Ob and Onb atoms in the molecular structure are depicted in Fig. 8d that some oxygen atoms originally connected with Al atoms diffuse into the interlayer regions after 50ps of simulation. Comparing with the

3.2. Dynamical properties Dynamics properties of Aluminosilicate tetrahedron: The section mentioned above has investigated the structural evolution of alumino­ silicate chains at elevated temperature. Here the dynamical properties of Si and Al atoms are further analyzed. As shown in Fig. 8a, the MSD values for both Si and Al atoms are equivalent with each other at each temperature. Although the MSD values for Si and Al atoms at different temperatures increase with time elapsing. After 50ps, MSD values at 1500K are still lower than 3 Å2. Considering that the squared bond length of Si–O and Al–O is 6.00 Å2, majority of Si and Al atoms can only rattle at their original positions and cannot escape from their coordi­ nation cages. It agrees well with the molecular structure analysis that aluminosilicate skeleton can maintain structural integrity at the tem­ perature up to 1500K, indicating the low mobility of Si and Al atoms. Furthermore, TCF of Si–O and Al–O bonds are shown in Fig. 8b to evaluate the bond stability. The TCF curves of Si–O and Al–O bonds remain as constant of 1 during the whole simulation time as the tem­ perature is not higher than 800K. It indicates that both Si–O and Al–O bonds are very stable at this temperature range. However, when tem­ perature increases to higher than 800K, the bond stability starts to degrade. After 50ps of MD run at 1000K, the TCF value of Si–O bonds slightly reduces to 0.998 and that of Al–O bonds decreases to 0.972. At 1500K, the TCF value of Si–O bonds still maintain at larger than 0.99, while that of Al–O bonds decreases to 0.69. The Al–O bonds are less stable than Si–O bonds at high temperature. Furthermore, the

Fig. 8. (a) MSD of Si and Al atoms in C-A-SH at different temperatures. (MSD of Si is displayed as scatter and MSD of Al as dashed line.) (b) TCF of Si–O and Al–O bonds in C-AS-H at different temperatures. (c) Van Hove functions of bridging and non-bridging oxy­ gen atoms (Ob and Onb) in the Al–O bond at 1500K. (d) The molecular structure of C-A-SH at 1500K. The red balls are O atoms that originally coordinate to Al atoms; the cyan and blue balls are Ob and Onb atoms, respectively, which are bonded to the Al atoms during the simulation. (For interpre­ tation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

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Fig. 9. (a) Mean squared displacement (MSD) of Ow atoms in C-A-S-H at different temperatures; (b) The components of MSD of Ow atoms along three x, y and z directions at 1500K; (c) Diffusion coefficient evolution of water molecules with increasing temperature; (d) Time-correlated functions (TCF) of H-bonds (dashed line) and Hw-Ow bonds (scatters) in C-A-S-H at different temperatures.

quantities of new Al-Ob and Al-Onb bonds, most of the hydroxyl ex­ changes happen on the Al–OH groups rather than on the Al-Ob bond. It confirms the better stability of the Al-Ob bond than the Al-Onb bond visually. Dynamic properties of interlayer water. In order to investigate the dynamics properties of interlayer water, the MSD of Ow atoms as a function of temperature is plotted in Fig. 9a. The inset in Fig. 9a clearly depicts the glassy dynamics of interlayer water molecules at 300K: (1) within short elapsing time (t < 0.2ps), the ballistic motion with MSD proportional to t2 dominates. It indicates the inertial motion of water molecules (ballistic stage), not influenced by other surrounding mole­ cules. (2) When t ranges from 0.2ps to 2ps, the dynamics gets into “cage stage”, which is defined as a plateau and slight decrement in the MSD curve. The presence of cage stage is due to the collisions between water molecules in the confined environment. At this stage, the motion of water molecules is constrained and they can only vibrate in the cage constructed by adjacent water molecules. (3) Subsequently, water molecules escape from the cage and begin to diffuse. And the MSD curve exhibits a linear increase and steps to diffusive stage. As the temperature rises, the plateau in the MSD curve becomes less pronounced (Fig. 9a). At 1500K, the “cage stage” completely disappears from the interlayer water dynamics. Hence, the MSD curve only contains a parabolic jump and then a linear increase, similar with the dynamics of bulk water. It illustrates the gradual transformation for interlayer water dynamics from glassy state to free diffusing state with increasing temperature. At ambient temperature, the motion of the interlayer water molecules is susceptible to the confinement of H-bond network and is restricted strongly by adjacent atoms. Consequently, the interlayer water dy­ namics exhibits glassy nature. As mentioned in Section 2.1, the high temperature destroys the H-bond connectivity, weakening the chemical bonds that restrict motion of water molecules. Hence, the interlayer water molecules can move freely at high temperature. It is worth noting that the motion of interlayer water molecules is anisotropic different from the bulk water. The x, y and z directional components of MSD of Ow atoms at 1500K are shown in Fig. 9b. It can be observed that the

MSDx and MSDy show a dynamics similar to that of bulk water, including a parabolic jump and linear increase. The MSD along z direction exhibits only a parabolic increase. At t ¼ 50 ps, the MSD value in z direction is still below 10 Å2, contributing to the total MSD (¼ 307.58 Å2). It exhibits that the water molecules remain confined in the nanometer channel between calcium silicate sheets and have a quasi-two dimensional mo­ tion at temperature of as higher as 1500K. The diffusion coefficient of interlayer water molecules are calculated according to the diffusive stage in MSD curves, as shown in Fig. 9c. The diffusion coefficient of interlayer water molecules is 9.07 � 10 12 m2/s at 300K, which is two to three magnitudes lower than that of bulk water (D ¼ 3.54 � 10 9 m2/s) [67]. This implies the blocking effect of nano­ meter confinement on the motion of water molecules. With the rising of temperature, the diffusion coefficient of water molecules increases dramatically. When the temperature reaches 1500K, the diffusion co­ efficient of water molecules increases to 10.74 � 10 9 m2/s, which is in the same magnitude of the bulk water. Apart from the destruction of H-bond network, high temperature also weakens the bond strength of Ow-Hw within water molecules. It can be observed in Fig. 9d that the strength of O–H bond in the water molecules and H-bond degrade with temperature rising. It suggests that water molecules are more frequently dissociated at high temperatures, which also contributes to the increasing mobility of interlayer water molecules. Dynamic properties of calcium atoms. The MSD curves of prin­ cipal layer and interlayer calcium atoms (Cs and Cw) are shown in Fig. 10a. At 300K, the MSD values for Cs and Cw atoms are equal when time is less than 0.05ps. This is because the Cs and Cw atoms are in inertial motion during this time, thereby showing same mobility. After 0.05ps, the motion of Cs and Cw atoms are restricted by their coordi­ nation cages. Due to the different coordinate environments, the MSD curves for Cs and Cw atoms start to separate. At t ¼ 50 ps, the MSD for Cs and Cw atoms are 0.049 Å2 and 0.084 Å2, respectively. It can be concluded that both Cs and Cw atoms are tightly rooted in the C-A-S-H structure and it is difficult to escape from their fixed positions at this temperature. As the temperature rises, the MSD values of Cs and Cw 283

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Fig. 10. (a) MSD of primary layer and interlayer calcium atoms (Cs and Cw) at different temperatures. MSD of Cs is displayed as line and MSD of Cw as scatters. (b) Van Hove functions of the Cs and Cw atoms as a function of simulation time at T ¼ 1500 K. The distance of Ca–O bond approximate to the radius of the coordination cage. (c) Atomic trajectories of calcium atoms in xy plane at 1500K during the whole simulation.

atoms increases. At 1500K, while the MSD of Cw atoms is higher than 20 Å2, the MSD increment for Cs atoms is less than 2 Å2 throughout the simulation. It suggests that the Cs atoms are still trapped in the CaO octahedral “cage”. The van Hove function is employed to describe the inhomogeneous dynamics of calcium atoms at high temperature. The van Hove functions at T ¼ 1500K and t ¼ 0.05, 1, 10 and 50 ps are shown in Fig. 10b as a function of r. The displacement distributions of Cs and Cw atoms are completely overlapped at 0.05 ps. This matches well with the MSD analysis in the inertia stage (t < 0.05 ps). At t ¼ 1 ps, part of the Cw atoms jump out of the coordination cage, as the probability function has extended to the r range higher than 3.05 Å. On the other hand, all the Cs atoms have displacements less than 3.05 Å. With the elapsing time, the function of Cw atoms continuously extends to higher r values, whereas the function Cs atoms still remains in r range of 0 Å to 3 Å. The atomic trajectories of Cs and Cw atoms along xy plane at 1500K are also pro­ vided. As shown in Fig. 10c, the atomic density of Cs atoms is distributed as isolated cycles. It means the Cs atoms are vibrated at their fixed po­ sitions. The atomic density of Cw atoms is distributed as interconnected clouds, where the high intensity positions are the energetically preferred sites.

3.3. Uniaxial tension test Mechanical properties. To evaluate the mechanical response of CA-S-H at elevated temperature, uniaxial tension tests along y direction were performed on these models. The stress-strain relationship, char­ acterizing the constitutive relation of materials, is exhibited in Fig. 11a. At ambient temperature (300K), the stress-strain curve increases linearly as the strain increases from 0 Å/Å to 0.13 Å/Å (Fig. 11a). The stress reaches the maximum of 10.82 GPa at strain of 0.13 Å/Å. Subsequently, the stress drops fast and then maintains at a plateau of 8.40 GPa. At the strain of 0.35 Å/Å, the stress reduces again and decreases to 0 GPa at strain of 0.67 Å/Å. Comparing with the stress evolution at ambient temperature tension, the initial linear increment in the stress becomes smaller, and in the post-failure stage, the stress reduces slower during high temperature tensioning. The Young’s modulus and tensile strength for C-A-S-H are obtained from elastic stage of the stress-strain curves. As shown in Fig. 11b, the Young’s modulus and tensile strength of C-A-S-H decreases as the temperature increases, implying the degradation of the stiffness and strength of C-A-S-H. However, the stress reduction process becomes slower with increasing temperature, indicating high tempera­ ture enhances the ductility of C-A-S-H.

Fig. 11. (a) Stress-strain relationship of C-A-S-H along y direction at different temperatures; (b) Young’s modulus and tensile strength of C-A-S-H along y direction at different temperatures. 284

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Fig. 12. Molecular structure of C-A-S-H (view along x direction) under y directional tension with strain rate of 0.08/ps at (a) 300K and (b) 1500K. From top to bottom: the tensile strains are 0.2, 0.4, 0.6 and 0.8 Å/Å, respectively. (The cyan sticks represent newly formed hydroxyl groups during structural deformation). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

reaches 0.2 Å/Å, the percentages of Q2 and Q2(1Al) start to decrease and that of Q1 increases. This denotes the breaking of aluminosilicate long chains and formation of chain terminals. At strain value of 0.6 Å/Å, the variation of Q species percentages seems to terminate, which indicates the rupture of C-A-S-H structure. Furthermore, there is also ~1% of Q3þQ4 species formation during the depolymerization of aluminosili­ cate chains. It is a result of the structural rearrangement between broken chains. In respect to the elongation test at 1500K, the change of Q1, Q2 and Q2(1Al) species at about the same magnitude to those at 300K (Fig. 13b), although their original Q species percentages are different. But the evolution of Q species at 1500K appears to be more moderate compare to that at 300K. The Q species percentages are changing throughout the elongation process, which is consistent with the obser­ vations in molecular structure evolution. It confirms the gentler destruction process of C-A-S-H at higher temperature. Besides, ~2% of Q3þQ4 species is produced during the elongation, which means the structural rearrangement is more pronounced at high temperature tensioning. The structural rearrangement can explain the ductility of CA-S-H at high temperature. The hydroxyl number evolution during the elongation of C-A-S-H at 300K and 1500K are plotted in Fig. 13c–d, respectively. It can be noted that the transformation among hydroxyls is strongly correlated with that of Q species evolution. The number of water molecules decreases and that of hydroxyls (Ca–OH, Si–OH and Al–OH) increases, which are caused by the hydrolytic reactions of aluminosilicate chains and Si–O–Ca bonds in calcium silicate sheets. At 300K, the number of

Molecular structure deformation. The molecular structure evolu­ tion of C-A-S-H during elongation at 300K and 1500K are displayed in Fig. 12. At 300K, the fracture of C-A-S-H exhibits a brittle behavior (Fig. 12a). At the strain of 0.2 Å/Å, the C-A-S-H structure does not show observable cracks. As the strain increases to 0.4 Å/Å, small cracks appear in the structure, with hydroxyl groups formed around the cracks. At strain of 0.6 Å/Å, small cracks grow together and a large crack that almost penetrates the structure emerges in the C-A-S-H. Finally, the structure is completely stretched broken at strain of 0.8 Å/Å. On the contrary, the failure process of C-A-S-H at 1500K seems to be slower. The structure also shows evident small cracks at the strain of 0.4 Å/Å, as demonstrated by the hydroxyls around them. With the strain increasing to 0.6 Å/Å, one broad crack is produced in the structure. In addition to this crack, there also a number of small cracks randomly distribute. The connection of C-A-S-H bulk structure is better than the C-A-S-H with the same tensile strain at 300K. When the strain evolves to 0.8 Å/Å, two large cracks are observed on the C-A-S-H but the structure is not totally fractured. It is worth mentioning that the structural ordering of calcium silicate sheets decreases significantly when the C-A-S-H is elongated at 1500K. This implies that the aluminosilicate chains are polymerized and depolymerized frequently at high temperature. The evolution of Q species percentages and hydroxyl number are analyzed to give a quantitative insight on the structural deformation during the tension tests. As shown in Fig. 13a, at 300K, the Q species distribution almost does not change when the strain is below 0.2 Å/Å, which corresponds to the elastic deformation of C-A-S-H. After the strain 285

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Fig. 13. Q species distribution in C-A-S-H as a function of tensile strain along y direction under (a) 300K and (b) 1500K. The number variation of hydroxyl groups in C-A-S-H as a function of tensile strain along y direction under (c) 300K and (d) 1500K.

Acknowledgements

hydroxyls changes at relatively high rate when the strain ranges from 0.2 Å/Å to 0.6 Å/Å (Fig. 13c). At other strain ranges, the variation of hydroxyl number is relatively slower. This suggests that the depoly­ merization reactions start to occur at the end of the elastic deformation for C-A-S-H and the reactions take place hurriedly. On the contrary, the change of hydroxyl number becomes moderate at 1500K (Fig. 13d).

Acknowledgements are given to the National Natural Science Foundation of China (under Grant 51778513,51678317 and U1806225), Natural Science Foundation of Shandong Province under grant ZR2017JL024, the Fok Ying-Tong Education Foundation for Young Teachers in the Higher Education Institutions of China (Grant No. 161069) and the China Ministry of Science and Technology (under Grant 2015CB655101) for their financial support.

4. Conclusions Utilizing reactive molecular dynamics simulation, the structure, dynamics and mechanical properties of C-A-S-H at elevated temperature is studied. The simulation results can be concluded as follows:

Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.matchemphys.2019.05.020.

1. With the increase of temperature, both the density and stability of Hbonds between interlayer water molecules decrease. Owing to the weakened confinement from the H-bond network, the mobility of interlayer water molecules increases exponentially with rising tem­ perature. Meanwhile, the dynamics of interlayer water molecules transforms gradually from one featured of glassy water to one of bulk water. 2. At ambient temperature, calcium atoms in the primary layers and in the interlayer regions show similar dynamical properties. Both of them are get trapped in the coordination cages and vibrate at the fixed positions. As the temperature rises, the calcium atoms gradu­ ally exhibit inhomogeneous dynamics. At temperature of up to 1500K, part of the interlayer calcium atoms can jump out the “cages”. Due to the geometrical confinement from aluminosilicate clusters, the principal layer calcium atoms show slight increase in mobility with increasing temperature. At 1500K, they are still stuck within the CaO octahedra. 3. On the one hand, high temperature weakens the cohesive strength of calcium silicate sheets, resulting in decreasing stiffness and strength of the C-A-S-H. On the other hand, high temperature facilitates the structural rearrangement of aluminosilicate skeleton, leading to increasing ductility of C-A-S-H.

References [1] A.E. Mir, S.G. Nehme, Utilization of industrial waste perlite powder in selfcompacting concrete, J. Clean. Prod. 157 (2017) 507–517. [2] K. Scrivener, The concrete conundrum, Chem. World (2008) 62–66. [3] H. Mikul�ci�c, J.J. Kleme�s, M. Vujanovi�c, K. Urbaniec, N. Dui�c, Reducing greenhouse gasses emissions by fostering the deployment of alternative raw materials and energy sources in the cleaner cement manufacturing process, J. Clean. Prod. 136 (2016) 119–132. [4] B. Lothenbach, K. Scrivener, R.D. Hooton, Supplementary cementitious materials, Cement Concr. Res. 41 (12) (2011) 1244–1256. [5] K.P. Sethy, D.S.U.C. Pasla, Utilization of high volume of industrial slag in self compacting concrete, J. Clean. Prod. 112 (5) (2016) 581–587. [6] D.N. Huntzinger, T.D. Eatmon, D.N. Huntzinger, A life-cycle assessment of portland cement manufacturing: comparing the traditional process with alternative technologies, J. Clean. Prod. 17 (7) (2009) 668–675. [7] M.C.G. Juenger, F. Winnefeld, J.L. Provis, J.H. Ideker, Advances in alternative cementitious binders, Cement Concr. Res. 41 (12) (2011) 1232–1243. [8] M.H. Shehata, M.D.A. Thomas, Use of ternary blends containing silica fume and fly ash to suppress expansion due to alkali–silica reaction in concrete, Cement Concr. Res. 32 (3) (2002) 341–349. [9] J. Duchesne, M.A. B�erub�e, Evaluation of the validity of pore solution expression method from hardened cement pastes and mortars, Cement Concr. Res. 24 (3) (1994) 456–462. [10] I.G. Richardson, Tobermorite/jennite-and tobermorite/calcium hydroxide-based models for the structure of CSH: applicability to hardened pastes of tricalcium

286

Materials Chemistry and Physics 233 (2019) 276–287

J. Zhang et al.

[11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38]

silicate, β-dicalcium silicate, Portland cement, and blends of Portland cement with BFS, metakaolin or silica fume, Cement Concr. Res. 34 (9) (2004) 1733–1777. R.J. Myers, S.A. Bernal, N.R. San, J.L. Provis, Generalized structural description of calcium-sodium aluminosilicate hydrate gels: the cross-linked substituted tobermorite model, Langmuir 29 (17) (2013) 5294–5306. P. Faucon, A. Delagrave, J.C. Petit, C. Richet, J.M. Marchand, H. Zanni, Aluminum incorporation in calcium silicate hydrates (c s h) depending on their ca/si ratio, J. Phys. Chem. B 103 (37) (1999) 7796–7802. G.K. Sun, J.F. Young, R.J. Kirkpatrick, The role of Al in C–S–H: NMR, XRD, and compositional results for precipitated samples, Cement Concr. Res. 36 (1) (2006) 18–29. M.D. Andersen, H.J. Jakobsen, J. Skibsted, A new aluminium-hydrate species in hydrated Portland cements characterized by Al and Si MAS NMR spectroscopy, Cement Concr. Res. 36 (2) (2006) 3–17. M.J.A. Qomi, F. Ulm, J. Pellenq, Evidence on the dual nature of aluminum in the calcium-silicate-hydrates based on atomistic simulations, J. Am. Ceram. Soc. 95 (3) (2012) 1128–1137. D. Hou, Z. Li, T. Zhao, Reactive force field simulation on polymerization and hydrolytic reactions in calcium aluminate silicate hydrate (C-A-S-H) gel: structure, dynamics and mechanical properties, RSC Adv. 5 (1) (2015) 448–461. R.J. Myers, E. L’H^ opital, J.L. Provis, B. Lothenbach, Effect of temperature and aluminium on calcium (alumino)silicate hydrate chemistry under equilibrium conditions, Cement Concr. Res. 68 (2015) 83–93. G. Geng, R.J. Myers, J. Li, R. Maboudian, C. Carraro, D.A. Shapiro, Aluminuminduced dreierketten chain cross-links increase the mechanical properties of nanocrystalline calcium aluminosilicate hydrate, Sci. Rep. (2017) 7. J. Yang, D. Hou, Q. Ding, Structure, dynamics, and mechanical properties of crosslinked calcium aluminosilicate hydrate: a molecular dynamics study, ACS Sustain. Chem. Eng. 6 (2018) 9403–9417. D. Hou, T. Li, Influence of aluminates on the structure and dynamics of water and ions in the nanometer channel of calcium silicate hydrate (c-s-h) gel, Phys. Chem. Chem. Phys. 20 (4) (2018) 2373. D. Hou, T. Li, P. Wang, Molecular dynamics study on the structure and dynamics of NaCl solution transport in the nanometer channel of CASH gel, ACS Sustain. Chem. Eng. 6 (7) (2018) 9498–9509. C. Alonso, L. Fernandez, Dehydration and rehydration processes of cement paste exposed to high temperature environments, J. Mater. Sci. 39 (9) (2004) 3015–3024. P.A. Bonnaud, Q. Ji, K.J. Van Vliet, Effects of elevated temperature on the structure and properties of calcium-silicate-hydrate gels: the role of confined water, Soft Matter 9 (28) (2013) 6418–6429. O. Arioz, Effects of elevated temperatures on properties of concrete, Fire Saf. J. 42 (8) (2007) 516–522. G.A. Khoury, Compressive strength of concrete at high temperature: a reassessment, Magzine Concrete Research 44 (161) (1992) 291–309. F. Ulm, O. Coussy, Z. Bazant, The chunnel fire: chemoplastic softening in rapidly heated concrete, J. Eng. Mech. 3 (1999) 125. F.J. Ulm, P. Acker, M. Levy, The "chunnel" fire. II: analysis of concrete damage, J. Eng. Mech. 125 (3) (1999) 283–289. W. Ladaoui, T. Vidal, A. Sellier, X. Bourbon, Effect of a temperature change from 20 to 50� c on the basic creep of hpc and hpfrc, Mater. Struct. 44 (9) (2011) 1629–1639. Z. Pan, J.G. Sanjayan, Stress–strain behaviour and abrupt loss of stiffness of geopolymer at elevated temperatures, Cement Concr. Compos. 32 (9) (2010) 657–664. J.F. Justo, M.Z. Bazant, E.E. Kaxiras, V.V. Bulatov, S. Yip, Interatomic potential for silicon defects and disordered phases, Phys. Rev. B 58 (5) (1997) 2539–2550. M.Z. Bazant, E. Kaxiras, J.F. Justo, Environment-dependent interatomic potential for bulk silicon, Phys. Rev. B 56 (14) (1997) 8542–8552. M. Mrovec, M. Moseler, C. ElsSser, P. Gumbsch, Atomistic modeling of hydrocarbon systems using analytic bond-order potentials, Prog. Mater. Sci. 52 (2–3) (2007) 230–254. A.C.T.V. Duin, S. Dasgupta, F. Lorant, W.A. Goddard, Reaxff:, a reactive force field for hydrocarbons, J. Phys. Chem. A 105 (41) (2001) 9396–9409. T. Liang, B. Devine, S.R. Phillpot, S.B. Sinnott, Variable charge reactive potential for hydrocarbons to simulate organic-copper interactions, J. Phys. Chem. A 116 (2012) 7976–7991. Y.K. Shin, T.R. Shan, T. Liang, M.J. Noordhoek, S. Sinnott, Variable charge manybody interatomic potentials, MRS Bull. 37 (2012) 504–512. J.A.F.D. Keith, T. Jacob, A.C.T. Van Duin, Reactive forcefield for simulating gold surfaces and nanoparticles, Phys. Rev. B 81 (23) (2010) 235404. D. Raymand, A.C.T.V. Duin, W.A. Goddard, K. Hermansson, D. Spangberg, Hydroxylation structure and proton transfer reactivity at the zinc oxide water interface, J. Phys. Chem. C 115 (17) (2011) 8573–8579. W.A. Goddard, J.E. Mueller, A.C.T. van Duin, Development and validation of ReaxFF reactive force field for hydrocarbon chemistry catalyzed by nickel, J. Phys. Chem. C 114 (2010) 4939–4949.

[39] J.E. Mueller, A.C.T.V. Duin, W.A. Goddard, Application of the ReaxFF reactive force field to reactive dynamics of hydrocarbon chemisorption and decomposition, J. Phys. Chem. C 114 (12) (2010) 5675–5685. [40] A.C.T.V. Duin, A. Strachan, S. Stewman, Q. Zhang, X. Xu, W. Goddard, ReaxFF SiO reactive force field for silicon and silicon oxide systems, J. Phys. Chem. A 107 (19) (2003) 3803–3811. [41] H. Manzano, R.J.M. Pellenq, F.J. Ulm, M.J. Buehler, D.A.C.T. van, Hydration of calcium oxide surface predicted by reactive force field molecular dynamics, Langmuir the Acs Journal of Surfaces & Colloids 28 (9) (2012) 4187–4197. [42] L. Liu, A. Botero, G.W., H. Sun, Development of a ReaxFF reactive force field for ettringite and study of its mechanical failure modes from reactive dynamics simulations, J. Phys. Chem. A 116 (2012) 3918–3925. [43] D. Hou, T. Zhao, H. Ma, Z. Li, Reactive molecular simulation on water confined in the nanopores of the calcium silicate hydrate gel: structure, reactivity, and mechanical properties, J. Phys. Chem. C 119 (3) (2015) 1346–1358. [44] H. Manzano, S. Moeini, F. Marinelli, Confined water dissociation in microporous defective silicates: mechanism, dipole distribution, and impact on substrate properties, J. Am. Chem. Soc. 134 (4) (2012) 2208–2215. [45] H. Manzano, E. Durgun, M.J.A. Qomi, F.J. Ulm, R.J.M. Pellenq, J.C. Grossman, Impact of chemical impurities on the crystalline cement clinker phases determined by atomistic simulations, Cryst. Growth Des. 11 (7) (2011) 2964–2972. [46] H. Manzano, W. Zhang, M. Raju, J.S. Dolado, et al., Benchmark of reaxff force field for subcritical and supercritical water, J. Chem. Phys. 148 (23) (2018) 234503. [47] R.J.M. Pellenq, A. Kushima, R. Shahsavari, A realistic molecular model of cement hydrates, Proc. Natl. Acad. Sci. Unit. States Am. 106 (38) (2009) 16102–16107. [48] S.А. Hamid, The crystal structure of the 11 Å natural tobermorite Ca2. 25 [Si3O7. 5 (OH) 1.5]� 1H2O, Z. für Kristallogr. - Cryst. Mater. 154 (1–4) (1981) 189–198. [49] S. Murray, V. Subramani, R.,.H.K. Selvam, Molecular dynamics to understand the mechanical beha- vior of cement paste, Transp. Res. Rec. J. Transp. Res. Board 2142 (2010) 75–82. [50] E. L’H^ opital, B. Lothenbach, G. Le Saout, D. Kulik, K. Scrivener, Incorporation of aluminium in calcium-silicate-hydrates, Cement Concr. Res. 75 (2015) 91–103. [51] A. Allen, Time-resolved phenomena in cements, clays and porous rocks, J. Appl. Crystllogr 24 (1991) 624–634. [52] A. Allen, J. Thomas, H. Jennings, Composition and density of nanoscale calciumsilicate-hydrate in cement, Nat. Mater. 6 (2007) 311–316. [53] J.M. Pellenq, N. Lequeux, H.V. Damme, Engineering the bonding scheme in c–s–h: the iono-covalent framework, Cement Concr. Res. 38 (2) (2008) 159–174. [54] H.M. Aktulga, J.C. Fogarty, S.A. Pandit, A.Y. Grama, Parallel reactive molecular dynamics: numerical methods and algorithmic techniques, Parallel Comput. 38 (4–5) (2012) 245–259. [55] S. Plimpton, Fast parallel algorithms for short-range molecular dynamics, J. Comput. Phys. 117 (1) (1995) 1–19. [56] S. Kerisit, C. Liu, Molecular simulations of water and ion diffusion in nanosized mineral fractures, Environ. Sci. Technol. 43 (3) (2009) 777. [57] P. Chaudhuri, L. Berthier, W. Kob, Universal nature of particle displacements close to glass and jamming transitions, Phys. Rev. Lett. 99 (6) (2007) 060604. [58] D. Hou, D. Li, T. Zhao, Z. Li, Confined water dissociation in disordered silicates nanometer-channels at elevated temperatures: mechanism, dynamics and impact on the substrates, Langmuir the Acs Journal of Surfaces & Colloids 32 (17) (2016) 4153–4168. [59] P.A. Bonnaud, Q. Ji, B. Coasne, R.J.M. Pellenq, K.J.V. Vliet, Thermodynamics of water confined in porous calcium-silicate-hydrates, Langmuir the Acs Journal of Surfaces & Colloids 28 (31) (2012) 11422. [60] X. Cong, R.J. Kirkpatrick, Effects of the temperature and relative humidity on the structure of c-s-h gel, Cement Concr. Res. 25 (6) (1995) 1237–1245. [61] C. Meral, C.J. Benmore, P.J.M. Monteiro, The study of disorder and nanocrystallinity in c–s–h, supplementary cementitious materials and geopolymers using pair distribution function analysis, Cement Concr. Res. 41 (7) (2011) 696–710. [62] T. Radnai, P.M. May, G.T. Hefter, P. Sipos, Structure of aqueous sodium aluminate solutions: a solution x-ray diffraction study, J. Phys. Chem. A 102 (40) (1998) 7841–7850. [63] C.E. White, Effects of temperature on the atomic structure of synthetic calcium–silicate–deuterate gels: a neutron pair distribution function investigation, Cement Concr. Res. 79 (2016) 93–100. [64] J. Wang, A.G. Kalinichev, R.J. Kirkpatrick, Molecular modeling of water structure in nano-pores between brucite (001) surfaces, Geochem. Cosmochim. Acta 68 (16) (2004) 3351–3365. [65] M. Youssef, R.J.M. Pellenq, B. Yildiz, Glassy nature of water in an ultraconfining disordered material: the case of calcium silicate hydrate, J. Am. Chem. Soc. 133 (8) (2011) 2499–2510. [66] D. Hou, Y. Zhang, T. Yang, J. Zhang, H. Pei, J.L. Zhang, J. Jiang, T. Li, Molecular structure, dynamics, and mechanical behavior of sodium aluminosilicate hydrate (NASH) gel at elevated temperature: a molecular dynamics study, Phys. Chem. Chem. Phys. 20 (2018) 20695–20711. [67] J.P. Korb, L. Monteilhet, P.J. McDonald, J. Mitchell, Microstructure and texture of hydrated cement-based materials: a proton field cycling relaxometry approach, Cement Concr. Res. 37 (3) (2007) 295–302.

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