Drivers of Water Transport in Glass: Chemical or ... - ACS Publications

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Drivers of Water Transport in Glass: Chemical or Topological Effect of the Glass Network? Clémentine Mansas,† Jean-Marc Delaye,‡ Thibault Charpentier,§ Florence Bruguier,‡ Olivier Bouty,‡ Bruno Penelon,‡ Héléne Arena,† and Diane Rébiscoul*,† †

CEA, ICSM − UMR 5257 CEA-CNRS-UM-ENSCM, 30207 Bagnols-sur-Cèze Cedex, France NIMBE, CEA, CNRS, Université Paris-Saclay, CEA Saclay, 967, 91191 Gif-sur-Yvette Cedex, France ‡ CEA, DEN, DE2D/SEVT, Marcoule, F-30207 Bagnols sur Cèze, France §

S Supporting Information *

ABSTRACT: The relation existing between water transport in glass, topology, chemical elements (network former and charge compensator), and their structural role in glass were investigated through glass topology modeling, glass and water structure analyses, and water diffusion characterization. Two series of aluminosilicate glasses with and without boron and having various ratios of charge compensators CaO/Na2 O were used. The glass structure was characterized using Raman spectroscopy and nuclear magnetic resonance. Their topologies, i.e., the density of bottlenecks and the interstitial sites, were obtained by molecular dynamics. For glasses without boron, the substitution of Na by Ca leads to the strengthening of the glass network. The results are similar for glasses with boron except if the amount of Na is not sufficient. In this case, a part of Ca is required as a charge compensator for AlO4 and BO4 units. This last result is important since it highlights that the presence of boron and the CaO/Na2O ratio drives the roles of Ca inside the glass structure. Moreover, glass topology is driven by boron presence and to a lesser extent by CaO/Na2O ratio. Water transport characterized by the duration of the predominance of the hydration/interdiffusion processes, the apparent water diffusion coefficient, and the water structure in the hydrated glass were studied and determined using X-ray reflectivity and attenuated total reflectance infrared spectroscopy. The results show that the duration of the predominance of the hydration/interdiffusion processes is driven by the ability of Si-O-X bonds to be hydrolyzed as much as the fraction of free water clusters in hydrated glass. Moreover, for aluminosilicate glasses, we show that water transport is mainly driven by glass topology through the role of Ca and its amount in glass. Indeed, Ca strengthens the glass network by decreasing the density of the bottleneck allowing the diffusion of water molecules. When boron is added to the glass, water transport may be mainly driven by the chemical interactions between the water molecules, Ca, and the network former of glass matrix.

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INTRODUCTION

interdiffusion process was proposed. It supposes that glasses dissolve congruently within a thin interfacial film of water and alteration products then form by precipitation from species released into this film.6,7 However, the main part of the international community has agreed about the existence of the hydration/interdiffusion process. Several studies about the hydration/interdiffusion of float,5 archeological,8 natural,9−11 simplified,12−15 and nuclear16,17 glasses reported the measurement of the apparent water diffusion coefficients generally between 10−24 and 10−20 m2 s−1 at 298 K. While the effects of glass composition5,8,18−20 and solution12,21,22 on water diffusion were already studied and the interaction of water molecules with simplified glass matrix (sodium borosilicate and calcium aluminosilicate) was already modeled by atomistic modeling,23,24 the effect of glass topology

Water interaction with a material surface is of great interest to assess the material durability in various fields such as materials sciences, civil engineering, earth sciences, archeology, nuclear wastes, etc. This is particularly the case for the silicate glasses, materials existing in nature since several million years, used by humans for more than five thousand years and used in a wide range of industrial applications (coatings, window glass, nuclear wastes, etc.). Exposed to atmosphere or to solution, the glass surface is modified. The first stage of this modification is driven by two concomitant processes with one predominant over the other depending on glass composition and alteration conditions: (i) the water hydration/interdiffusion which consists of water diffusion through the glass network1,2 and ion-exchange between alkali ions contained in glass and positively charged water species coming from the solution3−5 and (ii) the hydrolysis of covalent bonds of the matrix. Recently another theory which does not take into account the hydration/ © XXXX American Chemical Society

Received: February 27, 2017 Revised: May 27, 2017 Published: May 31, 2017 A

DOI: 10.1021/acs.jpcc.7b01914 J. Phys. Chem. C XXXX, XXX, XXX−XXX

Article

The Journal of Physical Chemistry C

Table 1. References, Molar Compositions, Charge (Na + Ca), Glass Transition Temperature (Tg), Mass Density (ρm), and Electron Density (ρe) of G and GB Series F

SiO2

B2O3

Al2O3

Na2O

CaO

CaO/Na2O

charge (Na+Ca)

Tg (°C)

ρm (g cm‑3)

ρe (e Å‑3)

G0 G0.5 G1 GB0 GB0.5 GB1

73.2 72.5 72.6 62.9 62.9 61.6

0 0 0 15.0 15.3 15.6

5.2 5.4 5.4 4.5 4.2 4.3

21.7 15.4 11.2 17.6 12.0 9.2

0 6.7 10.8 0 5.6 9.3

0 0.46 0.96 0 0.46 1.01

21.7 22.1 22.0 17.6 17.6 18.5

504 593 634 540 588 606

2.431 2.458 2.470 2.478 2.466 2.454

0.7374 0.7351 0.7321 0.7221 0.7250 0.7344

network former modifies the structural role of Ca in glass, and it is partially removed during the hydration/interdiffusion stage.16 First, the structures of these glasses were characterized by Raman spectroscopy and nuclear magnetic resonance. Second, glasses were modeled using molecular dynamics in order to obtain their topologies, i.e., the density of bottlenecks and interstitial sites. Third, using the method developed in Rebiscoul and al.15,16 combining X-ray reflectivity and attenuated total reflection infrared spectroscopy (ATR-FTIR) the duration of the predominance of the hydration/ interdiffusion stage regarding the hydrolysis of glass matrix, the apparent water diffusion coefficient and the water structure in hydrated glass were determined. Finally, the effects of glass properties impacting these last results were discussed.

on transport was not investigated. As investigated by Kerrache and co-workers on borosilicate glasses (SiO2, B2O3, and Na2O),25 glass interstitial spaces are lower than 1 nm, and their size distribution depends on the glass composition. Actually, 1 nm is the critical size where liquid-like water does not exist and structural and dynamics properties of this interfacial water are influenced by confinement.26 Thus, water transport in glass is impacted by the glass properties, but the underlying processes need to be clarified taking into account the glass topology. Water transport can be driven by two glass properties. First, interactions between water molecules and chemical elements through their chemical nature and/or their structure in glass can play a role. Water molecules can chemically interact with the charged species such as the charge compensators,27 the non-brindging oxygen (NBO), and the surface chemistry of the glass network, i.e., the network formers. Recently, we showed that the mobility of water confined in nanoporous materials can be modified by the surface of the confinement (Si-OH, Al-OH, and Zr-OH)28 and the presence of ions (Ca and Na)29,30 which depends otherwise on the ions ability to be solvated.31 These ions modify the interactions of water with surfaces and then its dynamics and activity. Such interactions can be highlighted by the results obtained in the following studies32,33 showing that He, an inert gas having the same size that a water molecule34,35 and having no chemical interactions with the glass network, diffuses faster than water molecules in complex borosilicate glasses with diffusion coefficients between10−18 and 10−16 m2 s−1. Second, the topology of the glass may also drive water transport. Indeed, several studies showed that the atomic arrangements of atoms in silica based materials and the percentage of glass forming oxides impact the He diffusion.36 More recently, a study highlighted that the topology of simplified borosilicate glasses (SiO2-B2O3-Na2O), such as the interstitial sites accessible to He (radius of 1.28 Å), is influenced by the glass composition and more particularly by the molar fraction of Na.25 Even if such results suggest that water transport may be driven also by glass topology, the impact of this glass property on the water transport has not been studied yet. For the first time, we investigated the effects of these two glass properties on the water transport processes: the effect of chemical elements (network formers and charge compensators) and their structural role in glass and the effect of the glass topology. To reach this goal, the hydration/interdiffusion processes were studied as a function of glass composition at 30 °C and pH 3 to intensify this process and to limit the matrix hydrolysis. Aluminosilicate and borosilicate glasses with various ratios of charge compensators CaO/Na2O were used. Ca and Na were chosen as charge compensators for their different abilities to strengthen the glass network and to form complexes with water.27 Boron was added to the second glass series. This

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MATERIALS AND METHODS Materials. Samples Preparation. Two series of glasses with and without boron and containing various CaO/Na2O ratios were elaborated: aluminosilicate glasses (G series) 73SiO2− 5Al2O3−(22 − x)Na2O−xCaO where (x = 0, 7, 11) and borosilicate glasses (GB series) 63SiO2−15B2O3−4Al2O3−(18 − x)Na2O−xCaO where (x = 0, 6, 9). The references of these glasses are presented in Table 1. The glasses were prepared by mixing the following analytical grade oxides: SiO2, B(OH)5 (AnalaR NORMAPUR), Na2CO3 (Sigma Aldrich), CaO (Alfa Aesar), and Al2O3 (Sigma Aldrich). Glasses G were melted in a platinium crucible at 1500 °C for 2 h 30 min followed by 30 min at 1550 °C and glasses GB at 1400 °C for 3 h. Afterward, all glasses were poured into graphite crucibles and annealed for 1 h at Tg + 20 °C as indicated in Table 1. Molar compositions were confirmed by alkaline fusion and ICP analyses (Table 1). Monoliths measuring 25 × 12.5 × 2 mm3 were cut from the glass ingot and polished with ethanol to obtain an optical grade surface. Then, monoliths were cleaned ultrasonically in acetone and in ethanol for 5 min. Glass powders used for NMR analyses were obtained by grinding and sieving. Their sizes range between 63 and 125 μm. Glass Alteration. Glass monoliths were altered at a glasssurface-area-to-solution-volume (SA/V) of 1.2 cm−1 in a solution of ultrapure water having a pH(30 °C) 3 adjusted using HCl (Normapur) in PTFE reactors. Before alteration, solutions were placed in an oven at 30 °C for 24 h. The alteration durations were adjusted as a function of the glass and varied from a few minutes to several months. For each alteration kinetics, the same glass monolith was taken from the reactor, rinsed with ultrapure water, dried at the laboratory atmosphere, analyzed using FTIR-ATR and XRR, and then immersed again in its leachate. This method allows monitoring the surface evolution of the same glass sample alters in the same leachate. B


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Table 2. Atomic Compositions (in Number of Atoms) and Volumes of the Boxes (V) Used for the Simulations by Classical Molecular Dynamics ref

Si

O

G0 G0.5 G1 GB0 GB0.5 GB1

2357 2384 2420 1854 1886 1863

5913 6028 6113 5956 6055 6093

B

Na

Al

885 917 944

1396 1012 747 1037 721 558

334 356 359 268 252 260

Ca 220 361 169 282

total

V (nm3)

10000 10000 10000 10000 10000 10000

139 138 138 126 126 126

Spectra were recorded from 400 to 4000 cm−1, adding 32 scans with a 2 cm−1 resolution correcting from the background spectrum for each substrate. For each sample, three different locations were analyzed on the monoliths. The pressure on the piston was adjusted until the stabilization of the spectrum. In our experimental conditions, the incident IR beam has a limited penetration depth through the glass (500 to 5000 nm) allowing the characterization of the totality of the altered zone. Baseline adjustments and bands decomposition were performed using the OPUS software. The O−H stretching bands deconvolution were carried out using Gaussian functions adjusting the intensity and the width of fitting curves. The “best” fit was considered when the statistical parameter R was the lowest. X-ray reflectometry (XRR) was performed by a Bruker D8 diffractometer using Cu Kα1 (λKα1 = 0.154056 nm) radiation and standard θ-2θ scan for the data collections. Step size of 0.005 and count time of 2 s were used. Reflectivity curves were presented as the evolution of the logarithmic of intensity received by the detector as a function of the wave vector transfer q = 4π sinθ/λ with θ the incident angle and λ the X-ray wavelength. Experimental curves were fitted using the Fire4c_6 software based on the Parratt algorithm adjusting the electron densities ρe (e Å−3) of n layers on a substrate.39,40 In the case of our samples, we have used the model already proposed by Rébiscoul et al., 200716 and 201215 correctly fitting the altered glass zone. The uncertainties are ±0.005 e Å−3 for the electron density and ±2 Å for the thickness. The electron density of the substrate presented in Table 1, i.e., the pristine glass, was fixed and calculated from (eq.1)

Solid Characterizations. To determine the structure of the glasses, Raman spectroscopy and NMR were used. Raman spectroscopy was recorded at room temperature using a Horiba Jobin-Yvon LabRAM Aramis instrument (1800 groove/mm grating, edge filter, and motorized mount) equipped with a frequency doubled Nd:YAG laser at 532 nm. The laser was focused on the sample with an Olympus BX41 microscope. The beam diameter and penetration depth were around 1 μm. The spectral resolution of the CCD detector was about 0.2 cm−1. Typically, spectra were counted during 60 × 5 s and recorded from 100 to 1800 cm−1 with a 200 μm confocal hole and a 150 μm slit without filter (11.0 mW). 11 B, 23Na, and 27Al magic angle spinning (MAS) NMR spectra were collected on a Bruker 500WB Avance II spectrometer (operating at a magnetic field of 11.72 T) using a 4 mm (outer diameter of the rotor) Bruker CPMAS probe at a spinning frequency of 12.5 kHz. Data were acquired using a short single pulse excitation (1 μs, corresponding to a tip angle less than π/12 to ensure quantitativeness of the spectra) and recycle delay of 2 s for 11B and 1 s for 23Na and 27Al. 11B, 23Na, and 27Al chemical shifts were referenced to an external sample of 1 M boric acid solution (19.6 ppm), 1 M NaCl solution (0 ppm), and 1 M AlCl3 solution (0 ppm). 29Si MAS NMR spectra were collected on a Bruker 300WB Avance I spectrometer operating at a magnetic field of 7.02 T. A Bruker 4 mm CPMAS probe was used at a spinning frequency of 10 kHz. Silicon data were acquired using the CPMG sequence,37 by typically accumulating 32 echoes with an echo delay of 4 ms between consecutive 180° pulses. The echoes were then summed up and Fourier transformed to obtain the spectra. A 20 s recycling time was used; checks were carried out that no spectral deformation was observed for longer recycling delays (up to 1200 s). The spectra were referenced to an external tetrakis(trimethylsilyl)silane (TKS) sample, for which the highest intensity peak was situated at −9.9 ppm from that of TMS. Data were processed and fitted using an in-house written code (T. Charpentier) as detailed in refs 37 and 38. Wide angle X-ray scattering (WAXS) measurements were carried out on glasses on a lab-based Philips XPert Pro MPD diffractometer, equipped with a Mo anode tube ((λkα = 0.71 Å), a Zr filter located on the diffracted beam to avoid Kβ radiation, and an XCelerator semiconductor detector. Reflection configuration and θ-θ mode were used. The angular range varied between 2° and 156°, giving an absolute scattering vector Q from 0.5 to 17 Å−1. To minimize intensity loss at high angles, a variable divergence slit was tuned to maintain a constant irradiated surface on the powdered sample. Altered glass surfaces were analyzed using a FTIR spectrometer Bruker VERTEX 70 in ATR mode equipped with a DTGS/KBR detector in order to determine the evolution of O−H vibrations at the glass surface. The samples were placed on surface diamond and pressed with a piston.

ρm =

ρe ∑i CiAi Na ∑i CiZi

(1)

with ρm (g Å−3) the mass density (Table 1), cj the element amount j in the material, Zj the atomic number, and Aj the atomic mass. Regarding the numerous fitting results obtained in this study, we will only present the total thickness of the altered zone. Computational Details. Classical Molecular Dynamics Calculations. Empirical potentials developed recently for SiO2B2O3-Na2O glasses41 were used for classical molecular dynamics calculations. These potentials are suitable to represent the structures and mechanical properties of a large set of SiO2-B2O3-Na2O compositions. Good results were also obtained for the density simulations and distributions of the Qn entities.42 The atomic compositions of the simulated glasses corresponding to the experimental compositions (presented in Table 1) are listed in Table 2. The potential parameters for the SiO2-B2O3-Na2O basis of a xSiO2-yB2O3-zNa2O-tAl2O3-(1-x-y-z-t)CaO glass were chosen according to the following rules: C


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Table 3. EPSR Parameters, Lennard-Jones Well Depth, ε, Range, σ, and Effective Ionic Charge, Q, for the Different Atomic Species Present in the Glassesa

(i) x′SiO2-y′B2O3-(1-x′-y′)Na2O basis was calculated by saving the [SiO2]/[B2O3] and [Na2O]/[B2O3] ratios. x, y, z, t are the stoichiometric coefficients used to simulate glasses with Na and Ca and x′, y′ are defined as: x′/y′ = x/y and (1-x′-y′)/y′ = z/y. (ii) The empirical parameters for the x′SiO2-y′B2O3-(1-x′y′)Na2O basis were determined according to the work described in ref 41. In order to complete the empirical potentials representing the interatomic interactions involving Al and Ca atoms, the parameters fitted by Guillot and Sator were used.43 This model was validated on a large set of silicate compositions with SiO2 concentration ranging from 29.29 to 74.51 wt %, Al2O3 concentration ranging from 0.01 to 17.41 wt %, CaO concentration ranging from 0.06 to 11.87 wt %, and Na2O concentration ranging from 0.23 to 4.15 wt %. The Al and Ca charges were taken to be equal to 1.4175e and 0.945e respectively. Al-B and Ca-B interactions were taken to be purely Coulombic. The glasses were prepared using the following method. A liquid was prepared using a density lower by roughly 5% than the experimental density. The liquid was equilibrated at 4000 K during 100 000 time steps. The time step was equal to 1 fs for all stages of glass preparation. Then the liquid was quenched at 5 × 1012 K s−1 until 300 K, step by step in the NVT ensemble (constant number of atoms (N), constant volume (V), and constant temperature (T)). To perform the quench, a series of steps of duration equal to 20 000 time steps was applied by decreasing the temperature 100 K by 100 K. A 20 000 time step equilibration in the NPT ensemble (constant number of atoms (N), constant pressure (P), and constant temperature (T)) was applied at 300 K in order to determine the new equilibrium volume. A final relaxation of 5000 time steps was carried out in the NVE ensemble (constant number of atoms (N), constant volume (V), and constant energy (E)) to definitively relax the atoms around their equilibrium positions. For each composition listed in Table 2, four independent configurations containing 10 000 atoms were prepared. Except for structure factors, the results presented below correspond to the averages on these four configurations. Empirical Potential Structure Refinement. In order to confirm the results obtained by classical molecular dynamics and reproduce the experimental IVB proportion, simulations were performed with the help of the so-called empirical potential refinement structure (EPSR) software.44,45 Three glass compositions were chosen (GB0, GB1, and G0.5) as their respective X-ray structure factor was made available. EPSR is based on a reverse Monte Carlo algorithm (RMC), but with a cost function driven by a simplified molecular dynamics potential (Lennard-Jones plus Coulomb potential) combined with a perturbation term which evolves during the simulation to match experimental and simulated glass structure factors. A thorough presentation of EPSR atomistic simulations for borosilicate glasses can be found in previous paper.46 For the different glasses, both well depth ϵ and range σ parameters defining Lennard-Jones potential were the same as in former studies.46 On the contrary, the boron effective ionic charge Q was chosen to reproduce as closely as possible the experimental IVB proportion measured by 11B MAS NMR. The parameters used for EPSR calculations are listed in Table 3.

glasses G0.5 GB0 GB1 G0.5 GB0 GB1 a

EPSR parameters

Si

B

Al

Na

Ca

O

ϵ σ

0.25 1.4

0.25 0.85

0.25 1.55

0.25 2.25

0.25 2.66

0.25 3.2

Q

+2

0 +1.65 +1.35

+1.5

+0.5

+1

−1 −1.022 −0.978

All ε values were fixed to 0.25.

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RESULTS Glasses Structure and Topology. The effect of glass composition on structure and topology of aluminosilicate and borosilicate glasses were investigated by combining Raman spectroscopy, NMR, and atomistic modeling. Characterization of the Glass Structure. Depolymerization. Depolymerization was estimated with the percentage of non-bridging oxygen (% NBO) in the different glasses, knowing the glass composition with the following formula (eq 2). In case of GB glasses, the BIII and BIV were determined from the 11B MAS NMR data (Figure 4a). The results are presented in Table 4. %NBO =

2(Na 2O + CaO − Al 2O3 − %BIV B2O3) × 100 ∑O (2)

The G series presents the same % NBO = 18 due to the similar total amount of charge compensators (Na+ and Ca2+). This value is higher than those of the GB series comprised between 2.3 and 7.5% NBO. These results are consistent with the Raman spectra (Figure 1a,b) showing a higher depolymerization for G than GB series. Indeed, the broad band between 830 and 1275 cm−1 linked to the elongation modes of the Si-O bond in Qn species presents a main peak around 1085 cm−1, assigned to Q3 (Si2O52‑ sheet), having a higher intensity for the G series than for the GB series. Moreover, the presence of a peak at 944 cm−1 related to the Q2 of SiO4 tetrahedron in Raman spectra of the G series and not present on the GB series spectra can also attest to a higher depolymerization for the G series. For the G series, this high depolymerization can be explained by the fact that the modifying cations first compensate for AlO4 and the remaining cations lead to the formation of NBO by breaking Si-O bridging bonds.47 Indeed, the 27Al MAS NMR data (Figure 4e) present a peak shift toward negative frequency values with the increase of the CaO/Na2O ratio indicating that AlO4 are compensated by Na.48 This is also coherent with the shift toward negative frequency of the peak obtained by 23Na MAS NMR data (Figure 4c). For these glasses, the 29Si NMR data mainly show the presence of Si(Q3) at 90 ppm and Si(Q4) at 100 ppm (Figure 4b). The 29Si chemical shift is primarily sensitive to the degree of polymerization49 and to Al/Si mixing50 (and references therein). In this case as the Al/Si ratio is constant, the decrease of chemical shift observed is mainly due to the substitution of Na by Ca, as a second-neighbor of a silicon atom which influences the position toward lower chemical shifts51 and the shape of the curve.38,48 D


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The Journal of Physical Chemistry C Table 4. Molar Compositions and Structural Information of G and GB Glasses Series % mol

SiO2

B2O3

Al2O3

Na2O

CaO

CaO/Na2O

charge (Na+Ca)

G0 G0.5 G1 GB0 GB0.5 GB1

73.2 72.5 72.6 62.9 62.87 61.55

0.00 0.00 0.00 15.03 15.33 15.59

5.2 5.4 5.4 4.45 4.22 4.32

21.7 15.4 11.2 17.62 12.01 9.24

0.00 6.7 10.8 0 5.57 9.3

0.00 0.43 0.96 0 0.46 1.01

21.7 22.1 22.1 17.62 17.58 18.53

BIII

30 44 57

BIV

% NBO

% NBO/Si

70 56 43

18.0 18.2 18.2 2.3 4.7 7.5

45.0 46.0 46.0 8.4 15.2 24.4

Figure 1. Raman spectra of the G and GB series.

to BO4 units and more particularly to those belonging to the danburite rings. The intensity of this peak decreases with the increase of the CaO/Na2O ratio, highlighting a lowering of the danburite rings in the glass network. These results go with an increase of glass depolymerization with the CaO/Na2O ratio. Second, the BO4 peak shift is higher when the CaO/Na2O ratio increases (Figure 4a). This shift, sensitive to second neighbors of boron52,53 may be explained by the substitution of Na by Ca whose field strength is higher.54,55 However, IVB is usually preferentially charge compensated by Na rather than Ca.38,48 In order to determine the role of Na and Ca regarding IV B, the modified model of Dell et al,56 describing the structure when alkali are added to the glass was used. Two cases were considered

Regarding the borosilicate glasses GB, the intensity of the main band around 1100 cm−1 on the Raman spectra linked to Si(Q3) (Si2O52‑ sheet) tends to shift to low wavenumber and decreases with increasing CaO/Na2O ratio. This indicates the formation of Si(Q2) species and attests to a depolymerization of the silicate network confirming also the increase of %NBO with the CaO/Na2O ratio (Table 4). To conclude, aluminosilicate glasses (G series) are more depolymerized than borosilicate glasses (GB series). The substitution of Na by Ca has no influence on the depolymerization of aluminosilicate glasses whereas it leads to an increase of the depolymerization for borosilicate glasses. Si-O-Si Angle. From Raman spectra, the modification of the Si-O-Si angle with the CaO/Na2O ratio in the various glasses was determined using the Okuno relation based on the shift of the peak at 500 cm−1 (elongation and bending vibration modes of Si-O-Si bonds): a decrease of 5.5 cm−1 of the peak maximum corresponds to a one degree increase of the Si-O-Si angle. With the increase of the CaO/Na2O ratio, the peak maximum is shifted from 509 to 475 cm−1 for the G series and from 499 to 481 cm−1 for the GB series, i.e., an increase of the Si-O-Si angle of 6.1° and 3.2°, respectively. Boron Coordination and Borosilicate Network. 11B MAS NMR performed on the GB series (Figure 4a) gives information about the boron speciation and the interactions between elements within the borosilicate network. First, the proportion of the two boron coordination numbers IIIB and IVB were quantified from the area under the peak at 0 ppm attributed to BO4 units and the large peak between 3 and 20 ppm attributed to the BO3 units. As presented in Table 4, % BO3 increases with the increase of CaO/Na2O ratio to the detriment of % BO4. This is supported by the increase of the band intensity at 800 cm−1 and the large band between 1300 and 1600 cm−1 corresponding respectively to the vibration of the boroxol rings and the stretching vibrations of B-O-B and BO− in the BO3 units with the substitution of Na by Ca. Raman spectra of the GB series also show a peak at 630 cm−1 assigned

(i) AlO4 and BO4 units are compensated by Na: R = (Na2OAl2O3)/B2O3 (ii) AlO4 and BO4 units are compensated by Na and Ca: R* = (Na2O + CaO-Al2O3)/B2O3 The percentage of IVB obtained by NMR vs these two ratios is presented in Figure 2. A correlation is obtained only when Na is considered as a charge compensator for AlO4 and BO4 units. Thus, this shift of the BO4 peak is more likely the result of a change in the Si/B mixing as it has already been reported in

Figure 2. Correlation between the percentage of IVB obtained by NMR and the glass chemical composition R and R*. E


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The Journal of Physical Chemistry C the literature.38,52 The 23Na and 27Al MAS NMR data (Figure 4d,f) confirm that it is mainly Na that compensates for AlO4 units: negative shifts for 23Na and moderate line width of 27Al MAS NMR spectra. In order to check if the sodium amount in a GB1 glass is sufficient to compensate for all BO4 units, the correlation between the percentage of IVB calculated using R from the model of Dell and Bray and the percentage of IVB determined from NMR is plotted on Figure 3. The broken line presents the

The structure factors in the simulated configurations were calculated using 11 configurations stored at regular intervals during the final relaxation at ambient temperature. The simulated structure factors can be compared directly to the experimental ones as presented in Figure 5 (X-ray structure factor was not measured for the GB0.5 glass). For each composition, only one configuration is used to compare simulations and experiments. The details of the experimental structure factors are correctly reproduced by the simulations even if the first peaks are slightly shifted toward higher wave vectors in the simulated structures and on the peak at around 5 cm−1 too large for some structures. A slightly too large Na-O distance (2.50 Å) may be at the origin of these discrepancies. Values ranging from 2.29 to 2.62 Å can be found in the literature.41 Interstitial Sites and Bottlenecks. As Malavasi et al. and Kerrache et al. show,25,57 the density of interstitial sites and bottlenecks were calculated for the two series of glasses and simulated by classical molecular dynamics in order to assess the free space in glass. For glasses containing boron atoms (GB series), we took into account the experimental boron coordination to estimate the number of IIIB in the simulation boxes (N3B). For example, for the GB0 glass, the experimental boron coordination is equal to 3.7, and there are 885 boron atoms in the simulation box. Thus, there should be 619.5 IVB and 265.5 IIIB. In order to confirm the results obtained by classical molecular dynamics, independent atomistic configurations of the GB0, GB1, and G0.5 glasses were calculated using the EPSR code. In the GB0 and GB1 configurations prepared by RMC, the B coordinations are equal to the experimental ones. The density of interstitial sites Ns and bottlenecks Nb were calculated using the remaining atoms with a code provided by Pedone.58 This code is able to measure the free volume inside each Delaunay tetrahedron 59 by subtracting from the tetrahedron volume, the regions volumes of the atoms intersecting it. The free surface was also calculated by subtracting the triangle surface delimited by three atoms, three intersecting regions between the atoms and the triangle. Figure 6a,b illustrates the Delaunay tetrahedron formed by four atoms and its associated surface formed by three atoms. From the free volume Vfree and the free surface Sfree of the Delaunay tetrahedron, the equivalent radius of interstitial sites rIS (eq 5) and bottlenecks rB (eq 6) was deduced.

Figure 3. Correlation between the percentage of BIV calculated using R = (Na2O-Al2O3)/B2O3 from the model of Dell and Bray and the percentage of BIV determined from NMR MAS 11B for the GB series.

limit from which Ca is required to compensate for AlO4 and BO4 units (below the line) and from which glass has enough sodium to compensate for AlO4 and BO4 units (above the line). For GB1 glass, the Na amount is not sufficient to compensate for BO4 units and a fraction of Ca is necessary to play this role. Simulations Results. Glass Density. Table 5 compares the experimental data with those obtained by molecular dynamic (MD) for B coordinations (CB) and glass densities (ρm). For densities, the differences between simulated and experimental values are less than 3% in each case. For boron coordination, larger values are systematically observed in simulated structures compared to experimental ones. A too strong attraction between Ca and B ions may be one of the main causes of these discrepancies since the B coordination would be lower if the Ca-B affinity was also lower. The Kieu’s potentials completed by the Guillot and Sator’s potentials used to represent the SiO2-B2O3-Na2O-Al2O3-CaO glasses are not able to correctly reproduce the relative interactions between the formers (Al and B) and modifiers (Na and Ca) and too many BO4 entities are formed. Structure Factor. In order to calculate the total structure factors in simulated glasses, partial structure factors were measured for each pair α-β using eq 3: Sαβ(Q ) =

fαβ N

Nα

Nβ

(3)

where rj and rk are the interatomic distances, N is the atom density, Nα and Nβ are the atom density of species α and β, and fαβ is the variable equal to 1 when α = β and to 1/2 when α ≠ β. Summations were performed on Nα and Nβ. The total X-ray diffraction structure factor was obtained from the partial structure factors using eq 4: SWAXS(Q ) =

N ∑α Nαfα2 (Q )

∑ fα (Q )fβ (Q )Sαβ(Q ) αβ

(5)

4πrB2 = Sfree

(6)

To analyze the different configurations, the atomic radii considered were equal to 0.41, 1.10, 0.27, 0.50, 0.95, and 0.99 Å for the Si, O, B, Al, Na, and Ca, respectively. The Si, B, and Ca radii correspond to the values proposed by Shannon or Pauling.60,61 The O radius corresponds to the half distance between O atoms in the O2 molecule. This value was chosen to guarantee that the sum of radii of the two neighboring atoms remains lower than the distance between them. For the same reason, the Al radius was slightly reduced to 0.50 Å (instead of 0.53Å for the Pauling radius) to guarantee that each Al-O distance remains larger than the sum of their radii. The distribution of the density of interstitial sites Ns and bottlenecks Nb as a function of equivalent radius of interstitial sites and of bottleneck for the two series of glasses are

∑ ∑ ⟨exp(iQ (rj − rk))⟩ j=1 k=1

4 3 πrIS = Vfree 3

(4)

with fα(Q) being the atomic scattering factor for the α species. F


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Figure 4. (a to f) 11B NMR, 29Si NMR, 23Na NMR, and 27Al NMR spectra of G and GB series.

Table 5. Comparison between Simulated Compositions (MD) and Experimental Values (Exp) for B Coordinations (CB) and Glass Densities (ρm) refs G0 G0.5 G1 GB0 GB0.5 GB1

CB (MD)

3.87 3.85 3.78

CB (Exp)

ρm (g cm‑3) (MD)

ρm (g cm‑3) (Exp)

3.70 3.56 3.43

2.413 2.469 2.502 2.480 2.512 2.524

2.431 2.458 2.470 2.478 2.466 2.454

presented in Figure 7a,b. These results highlight the presence of two mean Ns mainly around 0.5 and to a lesser extent around 0.7−0.8 Å for all glasses. For the GB series, the curve is slightly shifted toward lower radii than for the G series. Two mean Nb around 0.7 and 0.3 Å for all glasses are obtained. The main difference between the two series of glasses is the presence of bottlenecks having a size around 0.2 Å and a higher number of bottlenecks having sizes between 0.3 and 0.7 Å for the GB series. Moreover, the substitution of sodium by calcium slightly

Figure 5. Comparison between the experimental and simulated structure factors S(Q) for the glasses G0, G0.5, G1, GB0, and GB1. The simulated structure factors were smoothed by a running average over five successive values to reduce the fluctuations. The spectra were shifted for the sake of clarity.

decreases the Ns with radius between 0.7 and 0.8 Å. This may be related to Raman spectroscopy data highlighting a G


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Figure 6. Representation of the Delaunay tetrahedron formed from four atoms (the red spheres) (a) and one of its surface formed from three atoms (b). For (a), the gray regions correspond to the intersecting volumes between the atoms and the tetrahedron, and for (b) the shading regions correspond to the intersecting surface between the atoms and the face of the tetrahedron.

strengthening of the network with the CaO/Na2O ratio increase. These results suggest that glass composition has an impact on glass topology. Water Transport and Water Structure in Glass. In this part, the water transport in glasses was characterized using FTIR-ATR and XRR. The results obtained allowed the determination of the duration of the predominance of the hydration/interdiffusion stage, the apparent diffusion coefficients of water in the various glasses, and the water structure in the altered zone. Duration of the Predominance of the Hydration/ Interdiffusion Stage. The duration of the predominance of the hydration/interdiffusion stage τH was determined by FTIRATR. To reach this goal, we considered that hydration/ interdiffusion process was predominant when only the stretching band (2600−3800 cm−1) and the bending band (1650 cm−1) of O-H in the IR spectra were modified. The increase of these bands intensity corresponds to an increase of O-H bond density into the glass and thus to an increase of the hydrated and desalkalinized zone thickness. This observation generally goes with a decrease of the B-O bond stretching vibration at 1390 cm−1 for the GB series. Such a phenomenon was already modeled and characterized during the glass hydration/interdiffusion process,16,23,62 and it is related to the hydrolysis of the Si−O−IIIB bonds. When the hydrolysis of the glass network becomes predominant, the main peak corresponding to the Si-O bond stretching vibrations at 950 cm−1 is distorted in favor of another peak around 1000 cm −1 corresponding to the Si-OH stretching vibration. From these criteria, τH1 and τH2 were determined (Figure 8): 1. τH1 corresponds to the time required to observe the first modification of the Si-O bond stretching vibrations at

Figure 8. Durations of the predominance of the hydration/ interdiffusion stage process τH1 and τH2 for G0 and GB glasses. For G0.5 and G1 glass, τH being higher than 105 min, the values are not indicated on the graph.

950 cm−1 and the formation of a peak around 1000 cm−1 connected to Si-OH stretching vibration. 2. τH2 corresponds to the time required to have a significant modification of these bands as illustrated in (the SI, Figure S1). The evolutions of τH1 and τH2 as a function of the CaO/ Na2O ratio are presented on Figure 8. Even after 105 min, no modification of Si-O bond stretching vibrations at 950 cm−1 was detectable on FTIR spectra for glasses G0.5 and G1. Thus, τH1 and τH2 were not indicated on Figure 8 for these glasses. In the following section, τH, the arithmetic average of τH1 and τH2, will be used in the following sections. The results presented in Figure 8 highlight the impact of glass composition on the duration of the predominance of the hydration/interdiffusion stage. For glasses without Ca, the presence of boron in glass (GB0) leads to a 10 times longer τH than glass without boron (G0). For glasses with Ca, τH is higher for glasses without boron (G0.5 and G1) than with boron (GB0.5 and GB1). The GB1 glass presents the lower τH. Apparent Diffusion Coefficient of Water. The apparent water diffusion coefficients DH2O were determined from XRR and from the FTIR-ATR analyses performed during τH as in Rebiscoul et al.,15,16 2007 and 2012, i.e., when the hydration/ interdiffusion stage was predominant. DH2O of G0 and GB1 glasses were only determined from FTIR-ATR measurements because of the high roughness at a scale of the X-ray wavelength probably generated by the high glass hydrolysis rate. This

Figure 7. (a and b) Distributions of the density of interstitial sites Ns and bottlenecks Nb as a function of the equivalent radius of interstitial sites and bottlenecks for the two series of glasses. H


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Table 6. Apparent Diffusion Coefficients of Water DH2O Determined from XRR and FTIR-ATR Analysis for G and GB Glass Series

phenomenon has not allowed the getting of fringes on X-ray reflectivity curves as previously characterized in Rebiscoul et al., 2007 and 2012. The thicknesses thXRR, obtained from the fits of reflectivity curves (SI, Figure S2), were plotted as a function of the square root of time (SI, Figure S3). Results highlight a linear evolution between the thickness and the square root of time characteristic of a diffusive process. Thus, the second Fick law63 was used to calculate the apparent water diffusion coefficients (eq 7): thXRR = 2

DH2Ot π

refs G0 G0.5 G1 GB0 GB0.5 GB1

(7)

A vO − H + 2.694 0.059

(1.4 (1.3 (1.0 (2.0

± ± ± ±

1.0) 1.0) 1.0) 1.0)

× × × ×

−22

10 10−24 10−18 10−19

DH2O obtained by FTIR-ATR (m2 s‑1) (6.2 (9.9 (1.8 (9.3 (2.8 (1.3

± ± ± ± ± ±

2.0) × 10−20 1.5) × 10−23 3) × 10−24 1.5) × 10−19 1.5) × 10−19 2.0) × 10−17

that water diffusion coefficients in borosilicate glasses (10−20 < DH2O < 10−24 m2 s−1).15,16,22,64 Moreover, except for the GB1 glass presenting the highest DH2O, DH2O decreases with the CaO/Na2O ratio (Figure 10).

with thXRR the thickness in m, DH2O the apparent diffusion coefficient in m2 s−1 and t is the time in seconds. In order to specify these results, we also determined DH2O from the thickness thFTIR (Å) obtained from AvO−H the peak area of the O-H stretching between 2600 and 3800 cm−1 using the relation established in Rebiscoul et al.,15 2012, on soda-lime borosilicate glasses containing zirconium (eq 8): thFTIR =

DH2O obtained by XRR (m2 s‑1)

(8)

First, in order to test the validity of eq 8 for the two series of glasses used in this study, the evolutions of AvO−H versus thXRR were plotted and are presented in Figure 9.

Figure 10. Apparent diffusion coefficients of water DH2O in log scale determined by XRR and FTIR for the G and GB glass series.

Water Structure. In order to better understand the hydration/interdiffusion stage, the water structure in the altered layer was characterized using FTIR-ATR focusing on the stretching vibration of the O-H bond between 2600 and 3800 cm−1. This band was resolved in four contributions in the same way as Crupi et al.65 and Davis and Tomozawa:12 a first subband around 3000 cm−1 attributed to the ion or charge bonded water (X-O−, Ca2+, and H3O+ as example), a second subband of free H2O clusters around 3280−3308 cm−1, a third one corresponding to water linked to the surface, i.e., water molecules belonging to an interfacial layer around 3470−3522 cm−1 and the last one at 3585−3611 cm−1, related to surface and internal O-H. In some cases, two or three peaks between 2800 and 3000 cm−1 were observed and attributed to the C-H vibration coming from the spectrometer diamond cleaning step with ethanol. To take into account these contributions, two or three supplementary bands were included in the fitting procedure. An example of the deconvolution of stretching vibration of O-H bonds in four components with the best fit is presented in the SI Figure S4 for the G1 glass. The evolutions of the ratios of the various types of hydroxyls as a function of time for G and GB series are presented in Figure 11a−f. As displayed by these figures, the ratio of interfacial water and the ratio of surface and internal O-H bonds are approximately constant and equal for the two glass series. However, the proportions of free water and ion bonded water differ depending on the presence of boron in glass: the ratio of ion

Figure 9. Evolutions of peak area of the O-H stretching between 2600 and 3800 cm−1 AvO−H as a function of the thickness determined from XRR analysis thXRR.

A good correlation is obtained between AvO−H and thXRR and eq 8 for borosilicate glasses (GB series and glasses studied previously by Rebiscoul et al.15) but not for aluminosilicate glasses (G series). This is relevant to a large number of borosilicate glasses but not to aluminosilicate glasses. These results may indicate a difference of underlying hydration processes between both glass series. Consequently, for the G series, the values of thFTIR obtained from eq 8 are not validated. However, as a first approximation, they were calculated with this method. Consequently, the determination of DH2O calculated from the slope of the linear evolution of the thickness as a function of the square root time was possible. As for thXRR, thFTIR shows a linear evolution with the square root of time characteristic of a diffusive process (SI, Figure S3). The DH2O obtained from eqs 7 and 8 are presented in Table 6. DH2O obtained from XRR and FTIR-ATR analyses are about the same order of magnitude. DH2O of the G series are between 10−20 and 10−24 m2 s−1 and are lower than DH2O of the GB series comprised between 10−19 and 10−17 m2 s−1. These apparent coefficients are about the same order of magnitude I


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Figure 11. (a−f) Evolutions of the ratio of the different types of O-H vibrations during the glass alteration for the G and GB series. 3000 cm−1, ion and charge bonded water; 3280−3308 cm−1, free H2O clusters; 3470−3522 cm−1, interfacial water; 3585−3611 cm−1, surface and internal O-H.

presence of boron and the CaO/Na2O ratio drive the roles of Ca inside the glass structure. Moreover, glass topology is also strongly impacted by the glass composition: the boron presence and in to a lesser extent the CaO/Na2O ratio. The glass surfaces analyses allowed the determination of the duration of the predominance of the hydration/interdiffusion stage, the apparent diffusion coefficients of water in the various glasses and the water structure in the altered zone. Taking into account these results summarized in Figure 12, we will discuss in a first part the global impact of glass composition on the hydration/interdiffusion stage. Afterward, glass topology and chemical interaction between water and the glass matrix will be discussed to determine which one is the main driver of the water transport. Impact of Glass Composition. Duration of the Predominance of the Hydration/Interdiffusion Stage. The duration of the predominance of the hydration/interdiffusion stage, τH, is mainly driven by the glass composition through the CaO/Na2O ratio and the presence of boron. However, two different behaviors are observed (Figures 8 and 10).

bonded water is higher than the one of free H2O clusters in hydrated glasses of the G series whereas these ratios are close for hydrated GB0 and GB0.5 glasses. In hydrated GB1 glass, the free H2O cluster ratio is higher than the ions bonded water one. These results attest that the water structure is related to the glass composition and more particularly that the presence of boron in glass leads to a higher ratio of free H2O clusters than in glass without boron.

■

DISCUSSION The characterization and the molecular dynamic simulations of the two series of glasses allowed the determination of several characteristics. The structure of the two glasses series is different depending on the CaO/Na2O ratio and on the boron presence. The addition of boron globally increases the polymerization of the glass network. For glasses without boron, the substitution of Na by Ca strengthens the glass network. The same behavior is observed for glass with boron except when the amount of Na is not sufficient. In that case, a part of Ca is required to compensate for AlO4 and BO4 units. This last result is important since it highlights that both the J


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interdiffusion stage. This fact may be due to the Si-O-B hydrolysis which liberates some sites able to host water and probably free water clusters. Such water is more available to hydrolyze the Si-O-X bond of the glass network than ion or charge bonded water, leading to the predominance of hydrolysis. These results attest that τH is equally driven by the ability of Si-O-X bonds to be hydrolyzed than by the fraction of free water clusters. In order to enlarge the number of glasses and to determine the effect of the network former, the data obtained in refs 15, 38, and 69 about four glasses (Ca series: 63SiO2-15B2O32ZrO2-(18-x)Na2O-xCaO where (x = 0, 3, 6, 9)) were integrated in the discussion. These glasses are similar to the GB series (same NBO/Si) but Al2O3 has been substituted by ZrO2. As for the GB series, the CaO/Na2O ratio and the boron presence drive the roles of calcium inside the glass structure either as an element strengthening the glass network or as a charge compensator for ZrO2 and BO4 units. Apparent Diffusion Coefficient and Water Density in Hydrated Glass. Whereas the density of hydroxyl groups in hydrated glass is not proportional to the CaO/Na2O ratio and by extension to the NBO (Figure S5), as showed previously, DH2O strongly depends on the glass composition. Indeed, Figure 13 illustrates the strong relation existing between the

Figure 12. Diagram presenting the processes driving the water transport in G and GB series obtained in this study.

For the G series, while DH2O decreases when Na is substituted by Ca, τH drastically increases showing that hydrolysis is negligible compared to hydration/interdiffusion processes. This highlights the blocking role of Ca in the glass network regarding these two processes. Concerning the hydrolysis, Ca2+ ions are much less mobile than the Na+ and the bond (Si−O−)2Ca2+ is more difficult to break than SiO −Na +13,66,67 which may explain a longer hydration/ interdiffusion stage with fewer hydrolysis. The boron addition to glass (GB series) leads to lower τH than for aluminosilicate glasses (G series) and tends to reduce the predominance of hydration/interdiffusion against hydrolysis. This fact can be explained by the presence of Si-O-B bonds that are more easily hydrolyzable2,68 than Si-O-Si bonds. Moreover, for borosilicate glasses, τH depends on the role of Ca. When Na alone plays the role of charge compensator, the increase of the CaO/Na2O ratio leads to the same behavior than for the G series, i.e., an increase of τH. However, if a part of Ca is required to compensate for AlO4 and BO4 units, the glass network hydrolysis becomes easier regarding the hydration/ interdiffusion processes leading to a shorter τH. Furthermore, τH can also be related to the fraction of the different types of water existing in the hydrated zone. As displayed in Figures 8 and 11, longer τH are obtained for the G series which present a higher fraction of ion or charge bonded water. In these hydrated glasses, the evolution of the ratio of ion or charge bonded water as function of alteration time is quite constant even if the CaO/Na2O ratio increases. As the G glasses are highly depolymerized (%NBO = 18 and %NBO/Si = 45), this fraction of ion or charge bonded water may be related to water molecules bonded to Si−O−. For the GB series, the fraction of water bonded to ions or charge species in hydrated glasses is only correlated to the polymerization degree of silica at the beginning of the alteration. When the fraction of free water becomes higher than water bonds to ions or charge species, the alteration durations is close to the time required for hydrolysis to become predominant to the hydration/

Figure 13. Evolutions of the apparent water diffusion coefficients DH2O as a function of the CaO/Na2O ratio. NBO/Si of each glass and the other factors related to the glass composition which may impact the water diffusion are indicated.

DH2O, the nature of the network formers and the role of Ca. For glasses without boron (G series), DH2O decreases with the increase of the CaO/Na2O ratio. The addition of boron to the glass (GB and Ca series) induces a DH2O increase which depends on the nature of the network formers: ZrO2 leading to lower DH2O than Al2O3. Moreover, the substitution of Na by Ca leads to a slight decrease of DH2O when Na can compensate for AlO4/ZrO6 and BO4 units. However, when a part of Ca compensates for the network formers and BO4, DH2O drastically increases for the GB series or remains constant for the Ca series. Effect of Glass Topology. To study specifically the impact of the topology on water transport, the density of interstitial sites NsH2O and of bottlenecks NbH2O having a size larger than the H2O molecule radius (1.28 Å) were determined and related to the glass composition. In the following discussion, as the bottlenecks are the glass diffusion limiting parameter, NbH2O will be related to the DH2O. In the same way, as the interstitial K


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Figure 14. Evolutions of the density of interstitial sites NsH2O (a) and bottlenecks NbH2O (b) able to host and to allow the diffusion of one or several water molecules in pristine glasses, glasses without Na, and glasses without Na and BO3 as a function of the CaO/Na2O ratio for the G and GB series.

glass and the density of interstitial sites NsH2O able to host one or several water molecules (Figure S7). Figure 15 presents the evolutions of DH2O for the two series of glass, as a function of NbH2O in pristine glasses, glasses

sites are the space able to host water molecules, NsH2O will be considered as the main factor which may affect the water density in glass. Relation between Topology and Glass Composition. Using classical molecular dynamics calculations, NsH2O and NbH2O were calculated for the following three cases: (i) the pristine glass, (ii) the glass without Na as it is known that Na+ diffuse in solution during the hydration/interdiffusion stage, and (iii) the glass without Na and BO3 for the GB series since during the hydration/interdiffusion stage the hydrolysis of the Si-O-B3 bonds can occur.16,23 In a first approach, Ca was not taken into account due to the fact that cationic charge compensators are retained into the bulk and do not exchange with hydronium ions as easily as the network cationic modifiers.13,24 For the calculations of the case (i) the total density of sites having a radius higher than 1.28 Å were summed. For the cases (ii) and (iii), the Na and, for GB series, the BO3 entities corresponding to the experimental IIIB coordination were eliminated from the list of the atoms, by choosing randomly the B sites and the first three first-neighbors O atoms around each B atoms. The distribution of Ns and Nb able to host or to allow the diffusion of water molecules are presented in Figure S6a,b. As showed in Figure 14a,b, presenting the evolutions of the total densities of interstitial sites NsH2O and bottlenecks NbH2O as a function of the CaO/Na2O ratio, NsH2O and NbH2O are significantly different depending on the case considered. For pristine glass, whatever the glass series, the NsH2O are very low and close to zero and the NbH2O slightly increases with the CaO/Na2O ratio. For glasses without Na, NsH2O and NbH2O are two to three times higher than those of pristine glasses. These values decrease with the CaO/Na2O ratio, and they are obviously correlated with the substitution of Na by Ca and consequently with the glass network strengthening. NsH2O and NbH2O of the glasses without Na and BO3 are almost two times higher than the ones of the glass without Na and the substitution of Na by Ca goes with an increase of NsH2O and NbH2O. This fact shows that the presence of boron in glass has a more significant impact on the density of sites able to host or to allow the diffusion of water molecules than the CaO/Na2O ratio. Impact of Topology on Apparent Water Diffusion Coefficient and on Water Density in Hydrated Glass. Whatever the case considered (pristine glasses, glasses without Na, and glasses without Na and BO3), no clear relation can be established between the density of hydroxyls in the hydrated

Figure 15. Evolutions of the logarithm of the water diffusion coefficients DH2O for the G, GB, and Ca series as a function of the total density of bottlenecks NbH2O able to allow the diffusion of one or several water molecules in pristine glasses, glasses without Na, and glasses without Na and BO3 for the GB series.

without Na, and glasses without Na and BO3. It appears that when Na and BO3 (for GB series) are removed from the glass, DH2O can be related to NbH2O by a simple global relation. The increase of NbH2O leads to an increase of DH2O. These results may attest to the partial dependence of DH2O with the glass topology where the most soluble elements are removed, sodium and boron. Moreover, when NbH2O tends to infinity, i.e., a bulk media, DH2O tends to a value around 4 orders of magnitude (3.4 × 10−13 m2 s−1) lower than the diffusion coefficient of water in bulk (DH2O(20 °C) = 2.5 × 10−9 m2 s−1). Assuming such difference, it may suggest that the water molecules also have chemical interactions with the glass network during their diffusion. To verify the predictive properties of this simple global relation for the Ca series, the NbH2O without Na and BO3 units were calculated with a function depending on the amount of oxides in glasses of the G and GB series (dashed line in Figure 15). However it appears that this function is not able to correctly predict DH2O for these glasses. This fact may come from the method used for the NbH2O calculation and/or from a L


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Figure 16. Evolutions of (a) the parameter W as a function of the CaO/Na2O ratio and (b) the logarithm of the water diffusion coefficients DH2O as a function of W for the G, GB, and Ca series15 of glasses. The glasses requiring Ca to compensate for AlO4/ZrO6 and BO4 are surrounded. Regarding these results, W can be related to glass chemical elements and their structural role in the glass matrix.

CaO/Na2O ratio for GB and Ca series. Moreover, when Ca does not compensate for AlO4/ZrO6 and BO4 units, the increase of the fraction of bonded water in these glasses is correlated to a decrease of DH2O as presented in Figure 16b which is not the case with NbH2O (Figure 15). For borosilicate glasses, this may suggest that chemical interactions of water molecules with Ca and with glass network formers are probably the main factor affecting DH2O in these borosilicate glasses. For the G series, the trends are different. As presented in Figure 16b, DH2O increases with W. It means that even if the proportion of bonded water increases, which is not due to the amount of Ca as illustrated in Figure 16a, the DH2O still increases. For these glasses series, as DH2O is not related to the amount of Ca and to the bonded water, the glass topology may be the main factor affecting DH2O. Indeed in Figure 15, the linear evolution of the ln(DH2O) with NbH2O clearly shows the influence of the density of bottlenecks accessible to water. For aluminosilicate glasses, the topology, which is mainly driven by the presence of Ca strengthening the glass network, may be the main factor affecting the water transport in glass.

chemical effect due to the substitution of Al2O3 by ZrO2. In this last case, a chemical effect of the surface of the hydrated glass network may be at the origin of this phenomenon. Effect of Chemical Interaction. To highlight an effect of chemical interaction between water molecules and elements from the glass, we have taken into account the water structure through the factor W. W is the ratio of the band area of the OH bonds corresponding to the bonded water, over the band area of the O-H bond belonging to free water. To better understand, bonded water corresponds to water molecules linked to ions or charged species (-O−, Ca2+, and H3O+) around 3000 cm−1 and to interfacial water about 3500 cm−1. Free water is defined as free clusters around 3300 cm−1 (eq 7). W = (A O − H3000 + A O − H3500)/A O − H3300 = A O − H(bonded water)/A O − H(free water)

(9)

The evolutions of W versus CaO/Na2O ratio and of the logarithm of DH2O versus W for all of the series are presented in Figure 16a,b, respectively. First, glasses having Al2O3 as a network former (G and GB series) have a higher amount of bonded water than free water. The opposite behavior is displayed in the case of the Ca series with ZrO2 as a network former. This is not related to NBO since it is higher in glasses of Ca series than in glasses of GB series, and the proportion of NBO remains constant for glasses of the G series. This may suggest that the glass network, i.e., the surface of the sites holding water in the hydrated glass, plays a role on the water structure. Indeed, Bouyer el al.24 showed that, in calcium aluminosilicate glass, the presence of water leads to the well know hydrolysis reaction and the formation of AlO3H2O entities by breaking the Al-O-Si bonds. Moreover, according to Briman et al.,28 alumina-based surfaces more strongly interact with the water dipoles and texture the water layers, i.e., the interfacial water molecules, than zirconia-based surfaces. Also, water can be more strongly immobilized by AlOH than Zr-OH surfaces, which behave as Lewis acids forming coordination bonds with water molecules. Second, for the GB and Ca series, W increases with the substitution of Na by Ca, whereas it decreases for glass without boron (G series). These may come from different underlying processes as already suggested from Figure 9. It was already determined that water in hydrous glasses forms a strong complex with Ca, decreasing the protons mobility.27 This may explain the increase of bonded water in hydrated glass with the

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CONCLUSION

For the first time, the complex relation existing between water transport in glass, topology, the chemical elements (network former and charge compensator), and their structural role in glass were studied through glass and water structure analyses, water diffusion characterization, and glass topology modeling. The results show that the duration of the predominance of the hydration/interdiffusion stage is driven by the ability of SiO-X bonds to be hydrolyzed and the fraction of free water clusters. Moreover, for aluminosilicate glasses, water transport is mainly driven by glass topology through the amount of Ca and its glass network strengthening role which decreases the density of bottleneck allowing the diffusion of water molecules. When boron is added to the glass, water transport may be mainly driven by the chemical interactions between water molecules, Ca, and the network former of glass matrix. In order to confirm the predominant factors at the origin of the water transport through the glass, i.e., topology and chemical affinity of the glass matrix, some experiments are ongoing to determine He diffusion through these glasses. These new data will allow the determination of the contribution of these two factors on water transport. M


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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.7b01914. (S1) Criteria of hydration/hydrolysis process determined by FTIR-ATR. (S2) Experimental and simulated reflectivity curves of glasses during alteration process. (S3) Thicknesses as a function of the square root of time. (S4) Example of experimental and fitted ATR-FTIR spectra. (S5) Evolutions of the density of hydroxyl groups in the hydrated glasses. (S6) Distributions of the density of interstitial sites Ns and bottlenecks Nb. (S7) Evolutions of the density of hydroxyl groups in the hydrated glasses as a function of the density of interstitial sites. (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] ORCID

Diane Rébiscoul: 0000-0001-6604-3167 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by CEA. We are grateful to Bruno Corso for XRR maintenance. We thank M. Moskura for his help in acquiring the MAS NMR spectra and Sylvain Peuget for the fruitful discussion about Raman spectroscopy.

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REFERENCES

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