Electronic Origin of Doping-Induced Enhancements of Reactivity

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Electronic Origin of Doping-Induced Enhancements of Reactivity: Case Study of Tricalcium Silicate Jian Huang,† Bu Wang,‡ Yingtian Yu,‡ Loredana Valenzano,§ Mathieu Bauchy,*,‡ and Gaurav Sant*,†,⊥ †

Laboratory for the Chemistry of Construction Materials (LC2), Department of Civil and Environmental Engineering, and Laboratory for the Physics of AmoRphous and Inorganic Solids (PARISlab), Department of Civil and Environmental Engineering, University of California, Los Angeles, Los Angeles, California 90095, United States ⊥ California Nanosystems Institute (CNSI), University of California, Los Angeles, Los Angeles, California 90005, United States § Department of Chemistry, Michigan Technological University, Houghton, Michigan 49931, United States ‡

S Supporting Information *

ABSTRACT: Systematic manipulation of the reactivity of silicate materials in aqueous environment remains a challenging topic. Herein, by combining first-principles and reactive molecular dynamics simulations, we present a complete picture of the influence of impurity species on hydration reactivity, using the reactive triclinic tricalcium silicate phase as an example. We show that although initial hydration is influenced by the surface’s chemistry and structure, longer-term hydration is primarily controlled by proton transport through the bulk solid. Both shorter- and longer-term hydration processes are noted as being intrinsically correlated with electronic features. These outcomes provide the first direct evidence of the linkages between electronic structure and the longer-term (i.e., on the order of several nanoseconds) hydration behavior and sensitivity of hydrophilic crystalline materials and also offer a pathway to efficient compositional design for similar materials.

1. INTRODUCTION AND BACKGROUND The behavior of silicate materials in aqueous environment has tremendous impact on a wide range of natural and engineering processes ranging from mineral dissolution1 to metabolism of living organisms. 2 A ubiquitous understanding of the interaction mechanisms between these materials and water is critical for applications such as manipulating chemical reactivity3−5 and novel material synthesis.6,7 In particular, the hydration mechanisms of highly reactive silicate phases such as calcium silicates remain less understood, despite their wide, hydration-based applications as technical cements,4,8 scaffolds for bone repair,7,9 and synthetic dental tissues,10−12 as well as continued interests in their hydration products.13,14 This renders efficient modifications of these materials through compositional design especially difficult. Indeed, although doping has been previously examined and some thermodynamic descriptors (i.e., the enthalpy of formation), charge localization, and charge transfer behavior have been used to postulate indications of reactivity in aqueous environments,3 these hypotheses have remained unconfirmed. More generally, the contribution of the bulk properties (if any) on hydration, a reaction generally initiated from the surfaces, remains elusive. By combining density functional theory (DFT) and molecular dynamics (MD) using a reactive potential, we unambiguously elucidate the influences of impurities on hydration reactivity, using as a case study tricalcium silicate © 2015 American Chemical Society

(Ca3SiO5, or C3S), the most reactive phase in anhydrous ordinary portland cement. First, stoichiometric and impure C3S variants of low-index orientations were constructed, and the faceted (100) surfaces were then contacted with molecular water using DFT. A detailed analysis of the energetics of water sorption with C3S was performed to elucidate reaction mechanisms and the affinity of a given surface for water. Second, the outcomes of DFT were used as inputs to directly observe the hydration of pure and impure C3S surfaces in contact with bulk water via reactive MD simulations. This approach allows one to combine DFT and MD to study both the initial and the longer-term hydration behavior and establishes the first direct correlations between dopant-induced electronic property modifications and the chemical reactivity of silicate phases. This has translational impacts on the compositional manipulation of calcium silicates and similar materials for enhanced reactivity and thus engineering performance.

2. COMPUTATIONAL METHODOLOGY 1. Density Functional Theory (DFT). The crystal structures of pure and impure T1-Ca3SiO5 (T1-C3S) developed previously were used as the starting inputs (bulk crystals) for Received: August 25, 2015 Revised: October 24, 2015 Published: October 26, 2015 25991

DOI: 10.1021/acs.jpcc.5b08286 J. Phys. Chem. C 2015, 119, 25991−25999

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The Journal of Physical Chemistry C the calculations.3 It should be noted that C3S features seven polymorphs, the most stable pure phase at room temperature being the triclinic T1 polymorph.15 Other polymorphs can be stabilized at room temperature by using dopants and are reported to result in alterations in hydration behaviors.16 To investigate the effect of doping only, that is without the contribution of different crystallographic features, we focus here on the T1 polymorph both for the pure and impure cases. Impurities in the form of Mg2+, Al3+, and Fe3+ were considered. All impurities were inserted through atomic substitutions, with Mg2+ occupying Ca2+ sites, while Al3+ and Fe3+ impurities were inserted simultaneously at Ca2+ and Si4+ sites, respectively (i.e., one atom at each site). The insertion of impurities was carried out at high-symmetry positions, while maintaining electroneutrality of the cell. For Fe3+ species, spin polarization was imposed in the antiferromagnetic state as this condition was determined to be the most stable.3 All calculations were performed at a density functional (DFT) level of theory using the Vienna Ab initio Simulation Package (VASP).17 A planewave basis set and projected augmented wave (PAW) pseudopotentials were used.18 The Perdew−Burke−Ernzerhof (PBE) generalized gradient approximation (GGA) was adopted to treat exchange-correlation functional.19 A cutoff of 600 eV was imposed on the kinetic energy. The integration of the first Brillouin zone (BZ) was performed at the gamma point. The convergence tolerance on the total energy was set to 10−6 eV. Atomic relaxations were allowed through structural optimizations, until the residual force on each atom was below 0.02 eV/ Å. To elucidate the role of magnetism, the most preferred antiferromagnetic (AF) solution was imposed when Fe3+ impurities were introduced. Graphical visualizations of structures and charges were performed using VESTA.20 Low index surfaces (see section 3.1) were constructed by cutting slabs along the corresponding crystalline plane from the optimized structures of pure and impure T1-C3S. To avoid interaction between the upper and lower surfaces, a 14 Å thick vacuum was inserted between the cleaved surfaces, after carefully evaluating its thickness. [SiO4] tetrahedra were preserved at surface terminations since the breakage of Si−O bonds would result in unstable surfaces with spuriously high surface energies. Since the slab thickness is critical for the calculation of surface energy, a slab constructed from two unit cells was selected. For computational efficiency, such a supercell (2 × 162 atoms) was considered large enough for the scope of the study, as a larger system (3 × 162 atoms) did not feature significant variations of the surface energy (≤0.01 J/m2). During geometrical relaxation, surface reconstruction was allowed by fully relaxing the slab structure (i.e., atomic positions and lattice constants dynamic optimization). The surface energy γs was then calculated as γs =

surfaces near different surface sites. The site sensitivity for water sorption was then quantified, i.e., as being in the neighborhood of a Ca, Si, or doping atom site. The influence of the orientation of the water molecule, that is, with the oxygen (H+up) or hydrogen atoms (H+-down) pointing toward the surface (see Figure 3a), was also considered to assess the role of molecular orientation on the nature/energetics of the sorption process.21 Since the upper and lower surfaces have asymmetric configurations due to the low-symmetry triclinic structure of T1-C3S, the cleavage gives rise to an artificial dipole moment. To limit such effects, which may affect the computed sorption energies, the results obtained for the two cleaved surfaces were averaged for each configuration in the final assessment. 2. Reactive Molecular Dynamics (MD). Because of the high computational cost of DFT calculations, performing a direct ab initio simulation to study the reactivity of the cleaved surface in contact with bulk water would be challenging. As such, the cleaved structures of the pure and impure systems, as predicted by DFT, were used as inputs for semiclassical MD simulations. The usual classical interatomic potentials used for silicate−hydrates22 cannot be used herein as within the framework offered by such potentials: (i) atoms have fixed charges, that is no charge transfer is possible, (ii) bonding energies do not depend on the environment of the atoms, and (iii) no bond rupture or formation is allowed, thereby preventing, e.g., observations of water’s dissociation into hydroxyl groups and protons. Reactive potentials overcome these limitations while remaining far more computationally efficient than DFT and ab initio molecular dynamics. In particular, ReaxFF23 has been applied to a wide range of material compositions and structures. Recently, ReaxFF has also been applied to study the dissociation of water molecules in C−S−H,24 dicalcium silicate,25 and pure C3S.26 Indeed, by dynamically updating bond orders and bond-order-dependent terms, ReaxFF is capable of modeling the dissociation of water as well as the recombination of O2− and OH− units. Following this approach, the surface reactivity of C3S can be examined for a relatively large system as shown below. The surface reactivity of pure and impure C3S was studied by bringing these crystal structures in contact with bulk water. Bulk water was created by relaxing a liquid water structure with a density of 0.99 g/cm3 (comprising 405 molecules) in a triclinic box adjusted to the shape of the corresponding C3S structure. The structure was relaxed at 300 K using a thermostat. The relaxed structures from DFT were used as initial configurations for the C3S systems (pure and impure) and were brought in contact with bulk water by stacking along the z-axis. A barostat maintaining a zero stress was applied along the z-axis direction, while the other two dimensions were fixed. The hydration of the surface was studied by running 5 ns long simulations, followed by another 1 ns for collecting the statistics. All the simulations were performed with a time step of 0.25 fs. Note that although we utilize ReaxFF MD to overcome the limitations of ab initio methods and, thereby, extend the study beyond the very initial stage of hydration (i.e., adsorption of water molecules on the surface), we nonetheless restrict our discussions to the hydration process occurring over a time scale of a few nanoseconds. As such, we omit processes like dissolution and reprecipitation, which are likely to become significant at longer time scales (microseconds).

Eslab − nE bulk 2A

where Eslab is the total energy of the slab, Ebulk is the total energy of the bulk crystal, n is the number of unit cells (in this case, n = 2), and A is the surface area. To investigate the influence of impurities on surface properties, the dopants were located at suitable (i.e., vacuum exposed) positions on the (100) surface. Similarly to the pure case, for Al and Fe impurities, [AlO4] and [FeO4] tetrahedra were preserved to avoid spuriously high surface energies. Water sorption was studied by placing a single water molecule in the proximity of the cleaved top and bottom 25992

DOI: 10.1021/acs.jpcc.5b08286 J. Phys. Chem. C 2015, 119, 25991−25999

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Table 1. Surface Energy (γs, J/m2) and Coordination Number (CN) of Ca Atoms on C3S Surfaces and the Area Fraction (S) of Facets on the Wulff Construction of Pure C3S

a

surface

(100)

(010)

(001)

(110)

(101)

(011)

(111)

γ CNa S

1.17 4.30/4.50 0.262

1.23 4.30/4.49 0.217

1.22 4.29/3.86 0.252

1.28 4.21/4.27 0.091

1.50 4.08/3.50 0.0

1.15 4.65/4.38 0.110

1.20 4.13/4.39 0.069

The two values correspond to the Ca atoms exposed on the lower and upper surfaces, respectively.

3. RESULTS 1. Surface Properties of Pure Ca3SiO5. Atomic positions and cell parameters of bulk T1-C3S were first optimized at DFT level, providing structural results (a = 11.75 Å, b = 14.31 Å, c = 13.76 Å, α = 104.8°, β = 94.4°, γ = 90.1°) in excellent agreement with experimental data (a = 11.67 Å, b = 14.24 Å, c = 13.72 Å, α = 105.5°, β = 94.3°, γ = 90.0°).27 Surfaces were then modeled by cutting slabs from the optimized bulk C3S structures. On exposed surfaces, for all the low-index crystalline planes, [SiO4] tetrahedra were preserved since Si−O bonds are much stronger than Ca−O bonds and undercoordinated Si atoms can result in spuriously high surface energy. Table 1 lists the surface energies of seven low-index planes of pure T1-C3S. Note that although an infinite number of surfaces can be constructed, we focused on the low-index surfaces, as these are often prominent among typical crystals. We observe that the surface energy of the (101) surface (1.496 J/m2) is much higher than those of other planes. Indeed, as shown in Figure 1, the (101) plane has a high cation density, thereby

As expected, and as shown in Table 1, the coordination numbers of the Ca atoms on the surface are lower than for those within the bulk crystal (see Supporting Information for more detailed data regarding Ca coordination), which feature an octahedral environment with six O neighbors. In general, we note that higher coordination numbers are representative of higher surface stability, which manifests in the form of lower surface energies. We note that Ca atoms show a different coordination on the two exposed surfaces, thereby revealing their structural differences. Based on the knowledge of the surface energies (see Table 1), the Wulff shape of pure C3S was constructed (see Supporting Information). Interestingly, the calculated Wulff shape of T1-C3S is noticeably different from that of the monoclinic polymorph,31 which suggests that they should feature different behaviors during surface-related processes, e.g., water adsorption. As the (100) surface occupies the largest relative area on the Wulff shape (≈26%), in the rest of this study, we focus on this facet to evaluate the surface reactivity of pure and impure C3S. This would offer representative insights into the surface hydration behavior of C3S, although real materials containing such phases, e.g., technical cements, could present other less favorable surfaces, depending on the preparation process. 2. Influence of Impurities on Surface Energetics and Stability. Starting from the optimized bulk structures of impure systems, (100) surfaces were prepared so as to contain one exposed impurity atom on the surface (see Figure 2). Note that for all doped samples, a constant number of four atomic sites are substituted (see section 1), so that the impurity dosage is maintained uniformly at 1.2 atom % in each case. The calculated surface energies, as obtained for the impure C3S

Figure 1. A visual representation of the (100) plane (pink) of T1-C3S. Ca, Si, and O atoms are shown as gray, blue, and red spheres, respectively.

indicating that a cleavage along this close-packed plane is more unfavorable than along other planes as it would involve the breakage of a higher number of bonds. The surface energies along the other crystalline planes ranged from 1.153 to 1.281 J/ m2, with the (011) plane showing the lowest surface energy. We note that the calculated surface energy of (100) surface of stoichiometric T1-C3S is 0.73 J/m2 higher than that of the (100) surface of CaO but remains around half of that of the (0001) surface of α-quartz (2.30 J/m2 28,29). Such a trend can be explained by the fact that the energy of the Si−O bonds is roughly 3 times larger than that of Ca−O bonds.30

Figure 2. Relaxed structures of the (100) surface for (a) Mg, (b) Al, and (c) Fe-doped T1-C3S. Ca, Si, and O atoms are shown as gray, blue, and red spheres, respectively. 25993

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The Journal of Physical Chemistry C variants, are summarized in Table 2. In general, it is noted that all substitutions induce a decrease in the surface energy. This is Table 2. The (100) Surface Energy (γs, J/m2) and Average Coordination Number on Surface (CNsurf) of Atoms on the Surface for Pure and Impure C3S, with a Constant Impurity Concentration of 1.2 atom % C3S

pure

Mg2+-doped

Al3+-doped

Fe3+-doped

γs,(100) CNsurf

1.17 4.40

1.14 4.41

1.16 4.44

1.08 4.22

expected as by destabilizing the C3S lattice, impurities tend to reduce the energy that is needed to break bonds.32 These results also indicate a need to study surfaces featuring atomic substitutions as they are thermodynamically preferred and, thereby, more likely to form under equilibrium conditions. It should be noted that impurity segregation on surfaces is not unexpected and has also been observed in other oxides and metallic materials.33−35 Broadly, it is noted that Al and Mg substitutions have a very slight effect on the surface energy, resulting in a decrease of the surface energy by 0.9% and 2.5%, respectively, as compared to pure C3S. In contrast, Fe doping induces a substantial reduction of 7.9% of the surface energy, which is suggestive of significant surface reconstruction. It should be noted that such a decrease cannot be explained by the substitution of Ca−O bonds by Fe− O bonds as e.g. the obtained surface energy remains much lower than that of the (100) surface of Fe2O3 (γs = 1.53 J/ m2 36). From the analysis of the surface topology, we find that the reconstruction results in a significant decrease (around 4.2%) of the coordination number of Ca atoms present on the surface. On the contrary, doping with Mg and Al substitutions induces limited, if any (20% the binding energy (1.74 eV vs 1.41 eV), signifying a substantially stronger tendency to cause water dissociation. In all cases, no substantial modifications of the bulk structure are observed. It should be noted that the occurrence of chemisorption at surface sites suggests the formation of a “hydrate structure” on the C3S surface. In the context of longer term hydration, such complexes could act to passivate the surface, reducing its reactivity, or serve as a zone of heightened reactivity, e.g., due to offering a preferred pathway for the penetration of protons (further discussion below). To better understand the relative effects of the different impurity species, a descriptor of the electronic surface resonance (SR)39 was computed by calculating the difference between the density of states (DOSS) of the surface and the bulk. Note that the SR has been shown to be a good indicator of the surface reactivity for transition metal oxides.38 Indeed, the adsorption mechanism can be interpreted as a concerted/ coupling model between the adsorbate and the SR.39 As shown in Figure 6, we find a clear trend showing that an increase of

Figure 4. Relaxed structures of a water molecule interacting with (a) the upper and (b) the lower (100) surface. The values shown for each structure are the corresponding binding energies (see text). Ca, Si, O, and H atoms are shown in gray, blue, red, and white, respectively.

that do not belong to any [SiO4] tetrahedra. Indeed, due to the strong/directional nature of the Si−O covalent bond, an O atom that is part of such a tetrahedron has fewer degrees of freedom to move, once connected to the H atom of a water molecule. In each case, chemisorption occurs via barrierless reactions, as noted previously.38 4. Influence of Impurities on the Sorption of Water. We also investigated the role of impurities on the binding energy between the surface and a water molecule. To this end, the water molecule is placed in the vicinity of the impurity site. As shown in Figure 5, only the Mg-doped system experiences physisorption, although in contrast to the pure system, the water molecule is rotated, and binding energy is slightly lower (0.75 eV vs 0.79 eV for the pure case). In contrast, Al and Fe dopants induce chemisorption. We observe that Fe induces a slight decrease of the binding energy (1.36 eV vs 1.38 eV for

Figure 6. Water adsorption energy (on tetrahedral sites) with respect to the descriptor of electronic surface resonances (SRs). The line shows a linear fit of the values.

the adsorption energy corresponds to a decrease of the SR descriptor, in agreement with previous results.38 This suggests that Al dopant induces higher reactivity, whereas Mg and Fe substitutions result in slightly lower reactivity at the exposed surfaces, as compared to the pure system. 5. Influence of Impurities on C3S Hydration. To simulate hydration occurring after the initial adsorption (i.e., first layer), we applied classical molecular dynamics using a reactive potential ReaxFF23 (see section 2). Figure 7a shows the structure of the pure C3S−water system following contact. The

Figure 5. Relaxed structures of a water molecule interacting with a doped (100) surface containing: (a) Mg impurity, (b) Al impurity, (c) Fe impurity, and (d) at a Ca site for Mg- and Al-doped surfaces. Ca, Si, O, and H atoms are shown in gray, blue, red, and white, respectively. It should be noted that the water molecule moves across the periodic boundary in (b) and (c). 25995

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substantially by the end of our simulations (≈10 Å for pure C3S). This is in agreement with recent observations in dicalcium silicate25 and pure C3S26 and in fact may even explain the origin of the “induction period” due to the formation of an intermediate (metastable) passivation layer on the C3S surface, following realistic (time scale) contact with water. A closer look reveals that this process primarily proceeds through proton hopping between the oxygen atoms proximate to the Ca atoms. Protonation of the surficial silicate tetrahedra occurs at the very initial stage of sorption, but it does not appear to contribute substantially to bulk hydration at medium time scales. The effects of dopants are also evident during this stage. In particular, in the case of Fe doping (see Figure 7b), although OH− groups form rapidly around the dopants species during initial sorption, the mobility of the protons in these OH− groups is limited, which results in an inhibition of any further hydration into the bulk structure; i.e., a true passivating effect is observed, even at very small (nanoseconds) time scales. As noted in Figure 8, the type of dopant plays a significant role in controlling the kinetics of the hydration of C3S. Interestingly, although the time scales are much different, the present results are in broad agreement with available experimental data regarding the hydration behavior of impure C3S crystals;40,41 that is, (i) the addition of Mg has limited effect at short time scale but results in somewhat increased reactivity at longer time scales (t > 0.3 ns),40 (ii) the addition of Al decreases reactivity at short time scales but, on the contrary, induces an increased in extent of reaction at longer time scales, and finally, (iii) the addition of Fe results in a significant decrease in reactivity. As such, the relative effect of impurities on the hydration of C3S ranks as Fe < pure < Mg < Al. To the best of our knowledge, this is the first time that detailed computational simulations have been able to successfully reproduce the accelerating or decelerating effect of impurities of the hydration kinetics of crystals and, more specifically, the anhydrous cement phases. It should also be noted that while these results shed great insight into very early age hydration, the impacts of impurities on cement reactions, which follow a dissolution−precipitation cycle in aqueous solution, may be more complex. For example, if the impurity ions, upon exsolution into the solvent (water), interfere with the nucleation or growth of the C−S−H precipitate, reaction rates may be reduced in spite of an enhancement in the intrinsic reactivity of the anhydrous C3S. Begarin et al. suggest such a mechanism, i.e., of poisoning of the precipitate as to why aluminum in C3S solid reduces its hydration rate (over a time scale of hours), versus phase pure C3S,42 though incontrovertible proof of such a “poisoning” mechanism is not available. The results therefore suggest that trivalent impurities, e.g., Al3+ and Fe3+, appear to be particularly “avoidable”if the intent is to enhance the overall C3S reactivity and in solution.

Figure 7. Hydrated structures of (a) pure and (b) Fe-doped Ca3SiO5 surfaces. The sticks represent OH− groups. Calcium, iron, and the rest of the oxygen atoms are shown as blue, gold, and red spheres, respectively, together with silicon tetrahedra. Nondissociated water molecules are shown as gray wires for clarity.

dissociation of water molecules, resulting in the formation of hydroxyl groups (OH−) and the diffusion of hydrogen (H+) from the surface, into the bulk C3S is observed. The extent of the hydration was tracked by computing the number of hydroxyl groups formed per unit of exposed surface, which is shown in Figure 8 for the pure and impure crystals. We observe

Figure 8. Number of hydroxyl groups per unit of exposed surface with respect to the time of hydration for the pure and doped (Al, Mg, and Fe) Ca3SiO5 surfaces. The inset shows the same data plotted on a logarithmic time scale.

a two-step hydration mechanism. (i) At small time scales (t < 0.3 ns), fast hydration of the exposed surface occurs. At this stage, some water molecules dissociate to stabilize the exposed undercoordinated atoms; that is, hydroxyl groups connect to the undercoordinated cations, while the remaining protons connect to available terminating oxygens. After this initial stage of hydration, that is, for t ≈ 0.3 ns, all the pure and impure C3S surfaces show a hydroxyl density of ≈7 OH−/nm2. (ii) At larger time scale (t > 0.3 ns), after a short plateau, the formation of hydroxyl groups continues, but with a deceleratingroughly logarithmicrate (see Figure 8 inset). Upon careful examination of the hydration process, we observe that hydration at medium time scales involves continued protonation beneath the cleaved surface. As shown in Figure 7a, after initial hydroxylation of the surface, a substantial amount of protons diffuse into the bulk and form OH− groups with the oxygen atoms inside the C3S structure. As a result, the thickness of the hydrated layers increases

4. DISCUSSION By combining DFT and reactive molecular dynamics simulations, we have shown that the hydration of C3S is a complex process involving multiple stages. On one hand, we note that the initial adsorption of water is highly influenced by the surface structure and chemistry. This process is very rapid and reaches equilibrium within 0.3 ns. On the other hand, we observe that further hydration, at medium time scale (1 ns or longer), continues by proton diffusion into the bulk structure. Because of the substantial differences in the time scales, hydration is dominated by bulk proton diffusion in the case of 25996

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The Journal of Physical Chemistry C C3S. Consequently, as shown in Figure 9a,b, it is unsurprising that the hydration sensitivity as described by OH− group

between these oxygen atoms would facilitate such hopping. More specifically, decreased localization of the VBM bands should facilitate proton hopping and therefore enhance bulk hydration. Electronic localization, as quantified by the maximum isosurface value (MIV, a larger MIV corresponds to less localization, and vice versa), has been used previously to identify C3S compositions with enhanced reactivity.4 That study, however, focused on the conduction band minimum (CBM).4 On the basis of the work presented herein, and as shown in Figure 9c, we propose that hydration sensitivity should be correlated with the MIV of the VBM, rather than that of the CBM. Indeed, as reported in our previous work,3 the MIV of the VBM for the four systems ranks as Fe (0.222 e/Å3) > pure (0.085 e/Å3) > Mg (0.082 e/Å3) > Al (0.077 e/Å3), such that a decreasing MIVVBM value corresponds with increasing reactivity as also noted above. To the best of our knowledge, this is the first time that a direct correlation between charge localization and reactivity has been clearly established, rather than assumed. More generally, the present results emphasize the predominant role of the electronic properties of materials in determining their reactivity. It is also significant that although initial hydration is indeed controlled by the properties of the surface (i.e., SR), longer-term reactivity is primarily dictated by the bulk structure, i.e., through the localization (as indicated by MIV) of the VBM. As such, and significantly, the methodology proposed herein serves as a rigorous means to uncover and understand the means and mechanisms by which dopants alter the hydration sensitivity of condensed phases with water.

Figure 9. OH− density as a function of (a) surface energy and (b) surface resonance. (c) The OH− group density as a function of the maximum isosurface value of the valence band maximum (MIVVBM), with the inset showing same data zoomed to the low MIVVBM region. The OH− group density is obtained from the last 1 ns of the MD simulations.

5. SUMMARY AND CONCLUSIONS First-principles calculations at a density functional level of theory were combined with reactive molecular dynamics to obtain a detailed understanding of the influences of Mg, Al, Fe species, i.e., substitutional impurities, on the hydration sensitivity of T1-Ca3SiO5. The surface energy, which characterizes the stability of the surface, is expectedly correlated to the density of defects on the surface (i.e., undercoordinated atoms) and is strongly influenced by the type of dopant used. However, the surface energy does not control the short- or longer-term hydration. In contrast, reactivity is strongly influenced by the electronic density of the crystal. Furthermore, although initial hydration does indeed depend on the electronic surface resonance, longer-term hydration, that is, after the initial hydroxylation of the surface, appears controlled by the bulk and, more specifically, the rate of proton conduction (transport) that is permitted therein. We also found that the bulk proton conduction is correlated with the localization of the valence band maximum. Thus, for the first time, a clear correlation between electronic property and longer-term hydration is established, which manifests in the form of enhanced reactivity when the nucleophilic valence band maximum is strongly delocalized. On a practical note, Mg and Al species are noted to enhance intrinsic reactivity, whereas Fe strongly slows down hydration, in agreement with available experimental results.40,41 This suggests that tuning the ionic (e.g., Mg) or covalent (e.g., Al, Fe) character of atomic bonds, by doping the atomic network, enables new means to enhance the reactivity of the (tri- and di-) calcium silicates and similar materials. Such manipulations of impurity distributions and crystallochemical configurations of the anhydrous calcium silicate phases may offer innovative opportunities to synthesize

formation after 5 ns of hydration (see section 5) shows no clear correlation with the surface properties (e.g., surface energy, SR, etc.). According to the present MD predictions, and as observed experimentally,40 the reactivity of the doped C3S systems investigated in this work ranks as Fe < pure < Mg < Al. Clearly, doping can affect C3S properties, which can, in turn, alter proton transport within the bulk crystal. As per frontier orbital molecular theory,43−45 it has been postulated that the protonation of molecules is initiated from their highest occupied molecular orbital (HOMO) as they represent the most nucleophilic sites.46,47 Applying this analogy to the solid phase, the HOMO is equivalent to the electronic bands at the top of valence band, i.e., the valence band maximum (VBM). Interestingly, we find that the location of VBM in C3S is indeed closely related to proton diffusion. We have previously identified that the VBM in pure C3S mainly consists of the p orbitals from the oxygen neighbors of the Ca atoms.3 This is in agreement with the outcomes of the present reactive MD simulations, which show that protons diffuse into bulk C3S primarily by jumping between these oxygen atoms, as discussed in section 5. If we consider proton hopping as a series of reactions, in which a given OH− group dissociates and protonates a nearby oxygen atom, the correlation between reactivity and the VBM, or the frontier orbitals in crystal, is then not surprising. Moreover, it is intuitive that the spatially extended VBM 25997

DOI: 10.1021/acs.jpcc.5b08286 J. Phys. Chem. C 2015, 119, 25991−25999

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The Journal of Physical Chemistry C

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such technologically relevant materials with enhanced reactivity and engineering performance to do “more with less”.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.5b08286. Calculated Wulff shape and Ca−O bond distribution on surface (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (M.B.). *E-mail: [email protected] (G.S.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors acknowledge financial support for this research provisioned by the University of California, Los Angeles (UCLA), and National Science Foundation (CAREER Award # 1253269). Access to parallel computational resources was provisioned by the Laboratory for the Chemistry of Construction Materials (LC2), the Physics of AmoRphous and Inorganic Solids Laboratory (PARISlab), the Institute for Digital Research and Education (IDRE) at UCLA, and the Extreme Science and Engineering Discovery Environment (XSEDE) supported by the U.S. National Science Foundation (OCI-1053575 and DMR-130039). The authors acknowledge the support of these facilities in making this research possible. The contents of this paper reflect the views of the authors, who are responsible for the accuracy of the data presented herein, and do not reflect the views and/or policies of funding agencies, nor do the contents constitute a specification, standard or regulation. The last author acknowledges support for this research provisioned by the Edward K. and Linda L. Rice Endowed Chair in Materials Science. G.S. acknowledges Prof. Fredrik P. Glasser (University of Aberdeen) for scientific discussions and his patient insights into crystallochemical configurations of cement phases in the foundational stages of this research.

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REFERENCES

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DOI: 10.1021/acs.jpcc.5b08286 J. Phys. Chem. C 2015, 119, 25991−25999