Article pubs.acs.org/JPCC
Realistic Models of Bioactive Glass Radioisotope Vectors in Practical Conditions: Structural Effects of Ion Exchange Antonio Tilocca* *
Department of Chemistry, University College London, London WC1H 0AJ, United Kingdom S Supporting Information *
ABSTRACT: Yttrium-doped silicate glasses with bioactive properties (YBGs) have potential applications in radiotherapy. The effect of gradual release of sodium and calcium from a YBG composition, modeling the ion-exchange processes occurring upon contact with an aqueous biological medium, was investigated through molecular dynamics simulations. The ion-exchanged hydroxylated structures exhibit some important differences with respect to the pristine dry composition, such as increased cross-linking and yttrium clustering, whose potential effects on their performances as radioisotope vectors are discussed. These findings also highlight the importance of taking into account the structural effects of ion exchange with an aqueous environment in order to develop more realistic predictive models of the behavior of bioactive glasses and related materials in practical conditions.
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INTRODUCTION The efficacy of conventional (external) radiotherapy to treat malignant tumor tissues is limited by the need to employ low doses of radiation, in order to reduce damage to healthy tissues surrounding the affected area.1 Internal or in situ radiotherapy allows safe delivering of much larger radiation doses, by incorporating active radioisotopes such as 90Y into microspheres of an inert vector that carries them through the blood vessels and directly into the tumor.2,3 Once they have reached the destination, the microspheres attack the target in two ways: by hitting it with localized high radiation and by blocking (embolizing) the capillaries that supply the tumor. These synergetic “radioembolization” effects rely on the physical and chemical stability of the radioisotope vector, in order to prevent the unwanted release of radioactive ions in the bloodstream before and after the microspheres reach the target area. In fact, highly durable yttrium aluminosilicate glass (YAS, Y2O3− Al2O3−SiO2) compositions have proven clinically effective for this application.4 However, a very stable radioisotope vector also presents some drawbacks: YAS glasses are not biodegradable, and as such the microspheres can reside in the host for considerably long times after the end of the treatment (i.e., after the radioactivity levels have decayed). Moreover, aluminosilicate glasses such as YAS are biologically inactive and do not establish any favorable interaction with their biological host, for instance, promoting regeneration of healthy tissues to replace those destroyed by the treatment. On the other hand, biodegradability, strong interaction with existing tissues, and ability to promote growth of new tissues represent known features of bioactive glasses (BGs).5,6 The use of BGs as radioisotope vectors could then be highly beneficial for the applications described above. The gradual biodegradation of the vector can minimize the unknown long-term effects of the particles residing in the host, and this property could also be © XXXX American Chemical Society
exploited in temporary implants containing the radiation source (such as those employed in brachytherapy of brain and other tumors),7 which would spontaneously dissolve after treatment without requiring secondary surgery for their removal. Further potential benefits of BGs as radioisotope vectors could derive directly from their bioactivity: their ability to form strong bonds with existing tissues8 might improve the embolization of some types of tumor such as bone, by effectively anchoring the Ycontaining particles to the local tissues; moreover, the ability of BGs to trigger the natural regeneration of new tissues9 can improve the post-treatment recovery. Fueled by these premises, some studies have started to explore the development of bioactive and biodegradable glasses as radioisotope vectors.7,10−12 Some interesting findings, like the lack of negative effects of doping with radioisotopes on the bioactivity of the glass, have emerged. The key issue is the need to achieve a balance between two opposite goals: short-term stability in a biological environment (necessary to effectively treat the tumor while the vector containing the radioactive source is stable) with medium- to long-term biodegradability of the glasses (needed to achieve the benefits of their dissolution, in terms of particle resorption and tissue regeneration). For this reason, the features and requirements of bioactive glasses suitable for radiotherapy applications are very different from those of conventional bioglasses such as 45S5: the latter would not represent a suitable vector for yttrium radioisotopes due to its high bioreactivity, and other base compositions of lower bioreactivity are required to ensure the short-term stability of the vectors. In particular, a higher-silica yttrium-doped bioactive glass (YBG) composition, labeled LYS, emerged in previous Received: August 11, 2015 Revised: November 16, 2015
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DOI: 10.1021/acs.jpcc.5b07804 J. Phys. Chem. C XXXX, XXX, XXX−XXX
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Y2O3),11 YBG compositions resulting from gradual release of modifier Na and Ca cations upon exchange with protons were formulated as Y0.09Na0.32−xCa0.16−yP0.02Si0.62O1.745−z(OH)z, with z = (x + 2y) describing the replacement of NaO− and CaO by OH−, that is, Na+ and Ca2+ by H+. In particular, three compositions with z = 0.1, 0.32, and 0.5 were modeled with (x, y) = (0.1, 0), (0.32, 0), and (0.32, 0.09), respectively, based on the reasonable assumption that calcium ions are released after sodium (Table 1).
experimental studies as a potential candidate for radiotherapy applications, on the basis of its measured low yttrium release rate combined with sufficient bioactivity.11 Atomistic simulations based on molecular dynamics (MD) have proven particularly effective for determining the structural properties of bioactive glasses13,14 and also allowed us to rationalize the suitability of LYS as a Y vector.15,16 These computational investigations focused on specific structural features of the dry, pristine glass that are known to control the behavior of the glass in a biological host. The underlying assumption is that these same features will also have a strong impact on the initial response of the glass once set in contact with the host. While this assumption is certainly valid, it is well-known that the structure of a bioactive glass rapidly changes as a result of interaction with the host: alkali ions are immediately leached from the glass through exchange with protons from the contact fluid.17−19 This results in an alkalidepleted hydrated gel exposed to the biological environment: as this modified layer directly controls the interaction between glass and host, its properties (also in relation to those of the dry substrate) need to be understood in order to predict the behavior of an implanted glass, such as a radioisotope vector. A recent MD study of Na/H-exchanged conventional 45S5 bioglass has revealed how structural changes induced by ion exchange can steer the formation of a silica-rich gel and affect the bioactivity of that composition.20 Here we apply a similar approach to study the effect of the initial exchange between modifier ions in the glass and protons from solution on the structure of Y-doped bioactive glasses with potential applications in radiotherapy, using the LYS glass as reference. It is important to note that the first stages of the dissolution of various different (bioactive and nonbioactive, synthetic and natural) silicate glasses in an aqueous medium involve the same mechanism/processes, ion exchange and cation depletion,17,21−25 thus justifying the use of a common computational framework for modeling the effects of initial aqueous corrosion of different glass substrates. What is not known, and needs to be understood, is the effect that such initial processes have on structural properties of the glass that control further corrosion, biodegradation, and bioactivity in each case. These effects are not the same for any silicate glass composition: as a matter of fact, the different resistance to corrosion or degradability of different glasses has important consequences for industrial applications, geological processes, and nuclear waste storage.22,25 Separate investigations are thus particularly critical for glasses belonging to different categories, such as conventional BGs (like 45S5) and the present ones. Moreover, here we are mainly interested in the effect of initial aqueous corrosion processes on the ability of the glass to host its radioisotope load and then eventually degrade after radioactive decay: no previous atomic-scale investigations have focused on these issues. The present results extend our understanding of the atomistic properties of YBG materials from dry to ionexchanged hydrated glass and highlight the importance of taking into account the structural effects of ion exchange and hydration in order to rationalize the behavior of these materials under realistic conditions.
Table 1. Atomic Compositions of YBG Glasses Modeled in This Work z
composition
0 0.1 0.32 0.5
Y90Na320Ca160P20Si620O1745 Y90Na220Ca160P20Si620O1645(OH)100 Y90Ca160P20Si620O1425(OH)320 Y90Ca70P20Si620O1245(OH)500
Models of approximately 3000 atoms were obtained by classical MD simulations with the DLPOLY code.26 The force field employed was a shell-model (SM) potential developed27 and successfully applied to model dry bioactive and yttriumdoped silicate glasses,15,16,28−32 extended with additional terms describing hydroxyl groups and their interactions with the other ionic species in the glass.20,33 Each oxygen ion (i.e., nonhydroxyl O and hydroxyl OH) is represented as a core−shell harmonic oscillator. Protons are bonded to the hydroxyl oxygen shell via a Morse potential.34 Buckingham terms describe van der Waals interactions between all oxygen shells (O and OH) and between oxygen shells and cations (M = Si4+, P5+, Na+, Ca2+, or Y3+). Truncated three-body terms control O−Si−O and O−P−O angles. This combination of shell-model and hydroxyl potentials has proven accurate in a large number of computational studies of hydrated silicates, phosphates, and other oxides35−39 as well as in our recent simulations20,33 of hydrated glasses. MD simulations in the NVT ensemble were carried out with a Berendsen thermostat, using damped adiabatic40 dynamics with a 0.2 fs time step to control the high-frequency core−shell motion. As in our previous study,20 the approach to model glasses incorporating different amounts of OH involved heating and thermal annealing at 600−900 K of an initial random configuration of the composition with the required OH content, inserted in a cubic box of ∼33 Å. The hightemperature stage is mainly aimed at accelerating structural relaxation by rapidly sampling the configurational space and efficiently removing unstable interactions. After the thermal relaxation, a final NVT run of 400 ps was used first to equilibrate the system at room temperature and then to generate the final MD trajectory from which the structural data discussed in the following were extracted.
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RESULTS AND DISCUSSION Figure 1 shows the oxygen coordination number of Si, P, Na, Ca, and Y, highlighting the individual contribution of hydroxyl (OH) and non-hydroxyl (O) oxygens in each case. The left panel shows that the increasing OH content, introduced by replacing Na+ and Ca2+ with protons, is associated with both network formers (Si and P) and modifiers (M = Na, Ca, Y). In the first case, the total coordination number remains exactly 4, because the formation of Si−OH and P−OH bonds is balanced by a corresponding loss of Si−O and P−O bonds, that is, OH
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COMPUTATIONAL METHODS Starting from the base dry LYS glass (Y0.09Na0.32Ca0.16P0.02Si0.62O1.745, or 62.3 mol % SiO2, 1.0 mol % P2O5, 16.1 mol % Na2O, 16.1 mol % CaO, and 4.5 mol % B
DOI: 10.1021/acs.jpcc.5b07804 J. Phys. Chem. C XXXX, XXX, XXX−XXX
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Figure 1. Number of oxygen atoms coordinated to (left) Si and P and (right) Na, Ca, and Y. The total coordination number (dashed lines) is split into contributions from hydroxyl (OH) and non-hydroxyl (O) oxygen species.
replaces O within the coordination shell of network formers. On the other hand, the total M−O coordination number increases with OH content, because the increased number of OH that enter the coordination shell of M is not fully balanced by a lower number of non-hydroxyl O atoms in the same shell. The association of hydroxyls with the network formers thus involves oxygen substitution, whereas their association with nonnetwork formers may also involve some degree of expansion in the coordination shell of the M ion. Figure 2 shows the
Figure 3. Qn distribution of Si (bottom) and P (top) for different OH contents. (Insets) Corresponding network connectivity (NC), total and split into self (Si−Si, P−P) and cross (Si−P, P−Si) contributions.41
decreased amount of linear chainlike Q2 units, with a corresponding increase in the amount of branched Q4 species. The Q2 → Q4 conversion results in a more interconnected network, as quantified by the increase in total network connectivity (NC = ∑nQn%, representing the average number of BO per Si) shown in the inset, from 2.6 for the dry (z = 0) LYS glass to ∼2.9 for the z = 0.3 and 0.5 compositions. The increased connectivity appears to be determined mostly by new “self” Si−O−Si linkages,41 whereas Si−O−P cross-links are less frequent. Phosphorus speciation and connectivity show a similar effect, even though more marked, with a net increase in higher-n Q species (especially n = 2 and 3) replacing the lown ones (orthophosphates and terminal Q0 and Q1 phosphates), which results in a NC increase from 0.8 for LYS to >2 for z = 0.32. At variance from Si, in the case of phosphorus the increase in connectivity appears as a result of P−O−Si cross-links, whereas self P−O−P links are absent. Another interesting effect is that the increasing Si and P connectivity appears to reach a maximum for z = 0.32, after which (with Ca being exchanged with H after all Na has been lost) no further increase appears possible. The presence of a marked fraction of non-Q0 P species already in the dry LYS glass deserves some comment. The occurrence of Si−O−P bonds in 45S5 Bioglass represents a controversial issue,42 although recent work has started to clarify the nature of the discrepancy between NMR and computational data.43,44 However, the present LYS base composition contains a much higher silica amount than 45S5 (62.3% vs 46.1%), and both experiments43 and simulations28 show that increasing the silica fraction and connectivity leads to a marked increase in the fraction of P−O−Si bonds. Whereas the substantial nonorthophosphate fraction determined by MD for dry LYS
Figure 2. Detail of the coordination shell surrounding a group of nonbonded hydroxyl groups, extracted from the MD trajectory of z = 0.3 glass. Si, P, O, Ca, and H atoms are colored yellow, brown, red, blue, and white, respectively, whereas the O atoms of bonded and nonbonded OH are colored purple and green, respectively.
atomic structure of the LYS z = 0.3 glass as obtained from the MD simulation: both Si-bonded and nonbonded (nb) OH are visible. The former appear to enter Si coordination in a similar fashion as nonbridging oxygens do in modified glasses; on the other hand, nbOH are associated in small clusters, surrounded by Ca2+ ions. These structural features will be further explored in the following. The way in which gradual ion exchange affects the glass network is illustrated with the help of Figure 3, which shows the Qn distribution for Si and P (a Qn species is a Si or P atom bonded to n bridging oxygens, BO). In the case of Si, the loss of Na and Ca modifier ions upon hydroxylation results in a C
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simultaneous increasing need to fill the tetrahedral oxygen shell of Si. The right panel shows the OH population within the coordination shells of the cations, highlighting that the increasing OH content is partitioned among both network formers and modifier cations. The greater slope of the curve corresponding to the modifiers highlights their strong(er) association with OH, compared to Si and P. The evolution of the hydroxyl distribution with OH content is graphically illustrated in Figure 5, which shows that while at low OH content the hydroxyl groups are partitioned more or less equally between their nonbonded and bonded forms, the additional OH groups at higher contents mainly bond to Si. The mutual association between different species in the glass and the hydroxyl groups can be analyzed on the basis of the Rmn factors in Table 2. Rmn is the ratio between the number of species n found in the coordination sphere of m (averaged over the MD run) and the number that would occupy the same region if species n were homogeneously distributed in all available volume.30,50 Rmn thus provides an unbiased measure of how close the arrangement of species n around m is to a uniform, homogeneous distribution. The latter corresponds to Rmn = 1, which indicates no mutual preference or avoidance between species m and n. Values of Rmn higher (or lower) than unity denote that species n favors (or avoids) having species m in proximity. Table 2 shows that whereas OH groups always preferentially associate with modifier ions, this tendency is reduced with increasing OH content: in other words, low OH contents will initially associate with modifiers rather than with Si/P, but at higher contents the additional OH groups will be shared between network modifiers and formers, resulting in a more uniform distribution of the hydroxyls across the whole glass structure. This is confirmed by the ROH−OH factors, which change from a value of 2.7 at z = 0.1, denoting a highly inhomogeneous (clustered) distribution, to a value close to 1 at z = 0.5, denoting a homogeneous distribution of OH. At the same time, the yttrium distribution does not appear to be affected by the level of ion exchange, with a constant RY−Y value around 1.7, denoting a more clustered arrangement of yttrium atoms compared to the base dry LYS distribution (for which RY−Y was measured as 1.43). As clustering is deemed to inhibit ion dissolution from a glass,30 this finding may suggest that, upon contact with the host, the ion-exchanged glass vector may develop an even higher short-term resistance to dissolution and leaching of its radioactive yttrium load. As ion exchange progresses, the increasing OH content and the associated decrease in modifier ion content result in an increased number of OH groups shared between Ca and Y, that is, belonging to the coordination shells of both a Ca and a Y ion. This is evident from Figure 6, where OH groups associated with Ca, Y, or both ions are displayed in different colors. The figure also shows that these shared OH groups tend to be more clustered than the nonshared ones. Therefore, as ion exchange progresses, the increased need for OH groups to share the coordination of two different modifier ions may result in enhanced mixing between Ca and Y ions, at least within these shared regions, which may have important consequences on the migration mechanism of these cations in the glass matrix.51
reflects a feature also found for other models of phosphosilicate compositions with approximately 60% silica,28,43 these MD estimates of the non-orthophosphate fractions can in some cases be 2−3 times larger than those determined by conventional 31P NMR.43 This discrepancy may be associated with the combination of partial Q1(P) overestimation by MD simulations and underestimation by standard NMR.43 In general, determining the P speciation in silicate glasses containing very low phosphorus amounts (for reference, the base LYS composition examined here contains only 1% P2O5) remains a challenge for most conventional experimental techniques. As a matter of fact, it has been shown that, for the benchmark 45S5 glass composition, advanced solid-state NMR techniques44 are able to detect a small fraction of Q1(P) species that frequently escapes detection by standard NMR, leading to much better agreement with the Q1(P) fraction predicted by MD.30 Therefore, the significant Q1(P) amount identified by MD in the dry LYS glass is in no way indicative of an intrinsic inaccuracy of the simulations45 but appears as a natural reflection of the different underlying glass network compared to the less cross-linked 45S5. Whereas these P−O− Si cross-links may limit the phosphate release and reduce the rate of apatite formation and thus the bioactivity of the LYSderived glasses, this would actually represent a desirable feature for the present applications as radioisotope vectors, for which very high bioactivity (such as that of 45S5) would result in unwanted premature leaching of the radioactive load. Another issue to consider is whether the observed trend could be biased by limited statistical accuracy, especially for the speciation and network connectivity of phosphorus, whose estimation is based on a relatively small number of P species in the sample. Figure S1 in Supporting Information shows that the error bars are generally small, especially for the network connectivity, and therefore the observed trend of increasing NC with OH content is not a mere reflection of limited statistics. Figure 4 highlights the way in which hydroxyl groups are arranged in the glass structure. The left panel shows that while
Figure 4. (Left) Percentage of OH bonded to Si, to P, and to neither Si nor P (nbOH). (Right) OH population within the coordination shells of Si, P, Ca, and Y.
most OH groups are found bonded to Si, a substantial fraction of “nonbonded” hydroxylsthat is, hydroxyls not bonded to either Si or Pis present. The occurrence and stability of these species in silicate glasses, especially at low hydration levels, have been shown in previous experimental and modeling studies.46−48 The percentage of nbOH decreases with increasing z, as a correspondingly higher fraction bonds to Si. This may be due to both the increasingly limited availability of modifier cations (needed to stabilize the nbOH)48,49 and to the
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CONCLUSIONS Previous computational studies of bioglasses for radiotherapy have concerned dry compositions, and no atomic-scale data are available on how the initial ion exchange influences properties of these materials that affect their performance as radioisotope D
DOI: 10.1021/acs.jpcc.5b07804 J. Phys. Chem. C XXXX, XXX, XXX−XXX
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Figure 5. Partitioning of OH groups within the glass structure for (a) z = 0.1, (b) z = 0.32, and (c) z = 0.5. Nonbonded, Si-bonded, and P-bonded hydroxyls are colored green, purple, and orange, respectively.
Table 2. Rmn Factors Calculated for Different m−n Pairs and OH Content OH content (z)
RSi−OH
RP−OH
RNa−OH
RCa−OH
RY−OH
ROH−OH
RY−Y
0.1 0.32 0.5
0.99 1.11 1.29
0 0.67 0.81
1.15
1.44 1.46 1.36
2.06 1.84 1.59
2.7 1.29 1.11
1.68 1.69 1.69
Si coordination, compared to the dry glass.20 This effect is predominant at low ion-exchange levels, whereas at higher OH contents the amount of nonbonded OH formed and the corresponding increase in network connectivity appear to reach saturation (probably as a result of the lack of modifier cations able to stabilize nbOH), and higher amounts of Si−OH bonds are observed. In terms of the properties of YBG materials relevant for their application in radiotherapy, the increased connectivity of the glass matrix upon ion exchange essentially implies the formation of a more stable/durable glass vector upon contact with the aqueous medium. The increased yttrium clustering observed for all hydrated compositions can also be associated with a slower release of these cations compared to the dry glass.30 Whereas these represent favorable features in the course of radiotherapy treatment, because they will prevent leaching of Y isotopes in the bloodstream before their radioactive decay, at the same time they may negatively affect biodegradation of the YBG particles after treatment. The present findings thus highlight the need to take into account alterations to the dry glass structure introduced by contact with the physiological environment of the host, in order to develop more refined predictions of the activity and suitability of these materials compared to models solely based on the structure of the dry, non-ion-exchanged glass.15,52
vectors. The rationalization of this effect in the present work thus represents a substantial step forward in the current effort aimed to understand the behavior of these materials at a fundamental level.14 Besides the new insight specific to radioisotope vectors, the results of this work illustrate the importance and need to move beyond models of dry compositions, especially when the target is represented by materials that rapidly transform, departing from the pristine reference, under practical conditions. The present simulations show that yttrium-doped bioactive glass, depleted of Na and Ca ions after their exchange with protons from the contact medium, will contain a more crosslinked silicate matrix compared to the pristine dry material. This is the result of the intrinsic stability of nonbonded OH groups in the glass, whose occurrence induces the formation of a higher amount of Si−O−Si cross-links in order to satisfy the
Figure 6. Hydroxyl association with modifier cations for (a) z = 0.1, (b) z = 0.3, and (c) z = 0.5. OH groups found within the coordination shell of Ca, Y, or both Ca and Y are shown as magenta, blue, and yellow spheres, respectively. E
DOI: 10.1021/acs.jpcc.5b07804 J. Phys. Chem. C XXXX, XXX, XXX−XXX
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(16) Christie, J. K.; Malik, J.; Tilocca, A. Bioactive Glasses as Potential Radioisotope Vectors for in Situ Cancer Therapy: Investigating the Structural Effects of Yttrium. Phys. Chem. Chem. Phys. 2011, 13, 17749−17755. (17) Mariappan, C. R.; Yunos, D. M.; Boccaccini, A. R.; Roling, B. Bioactivity of Electro-Thermally Poled Bioactive Silicate Glass. Acta Biomater. 2009, 5, 1274−1283. (18) Skipper, L. J.; Sowrey, F. E.; Pickup, D. M.; Drake, K. O.; Smith, M. E.; Saravanapavan, P.; Hench, L. L.; Newport, R. J. The Structure of a Bioactive Calcia-Silica Sol-Gel Glass. J. Mater. Chem. 2005, 15, 2369−2374. (19) Cerruti, M.; Bianchi, C. L.; Bonino, F.; Damin, A.; Perardi, A.; Morterra, C. Surface Modifications of Bioglass Immersed in TrisBuffered Solution. A Multitechnical Spectroscopic Study. J. Phys. Chem. B 2005, 109, 14496−14505. (20) Tilocca, A. Atomic-Scale Models of Early-Stage Alkali Depletion and SiO2-Rich Gel Formation in Bioactive Glasses. Phys. Chem. Chem. Phys. 2015, 17, 2696−2702. (21) Banchet, V.; Jallot, E.; Michel, J.; Wortham, L.; Laurent-Maquin, D.; Balossier, G. X-Ray Microanalysis in Stem of Short-Term Physicochemical Reactions at Bioactive Glass Particle/Biological Fluid Interface. Determination of O/Si Atomic Ratios. Surf. Interface Anal. 2004, 36, 658−665. (22) Hellmann, R.; Cotte, S.; Cadel, E.; Malladi, S.; Karlsson, L. S.; Lozano-Perez, S.; Cabie, M.; Seyeux, A. Nanometre-Scale Evidence for Interfacial Dissolution-Reprecipitation Control of Silicate Glass Corrosion. Nat. Mater. 2015, 14, 307−311. (23) Trotignon, L.; Petit, J. C.; Della Mea, G.; Dran, J. C. The Compared Aqueous Corrosion of Four Simple Borosilicate Glasses: Influence of Al, Ca and Fe on the Formation and Nature of Secondary Phases. J. Nucl. Mater. 1992, 190, 228−246. (24) Valle, N.; Verney-Carron, A.; Sterpenich, J.; Libourel, G.; Deloule, E.; Jollivet, P. Elemental and Isotopic (29Si and 18O) Tracing of Glass Alteration Mechanisms. Geochim. Cosmochim. Acta 2010, 74, 3412−3431. (25) Cailleteau, C.; Angeli, F.; Devreux, F.; Gin, S.; Jestin, J.; Jollivet, P.; Spalla, O. Insight into Silicate-Glass Corrosion Mechanisms. Nat. Mater. 2008, 7, 978−983. (26) Todorov, I. T.; Smith, W.; Trachenko, K.; Dove, T. M. DL_POLY_3: New Dimensions in Molecular Dynamics Simulations Via Massive Parallelism. J. Mater. Chem. 2006, 16, 1911−1918. (27) Tilocca, A.; de Leeuw, N. H.; Cormack, A. N. Shell-Model Molecular Dynamics Calculations of Modified Silicate Glasses. Phys. Rev. B: Condens. Matter Mater. Phys. 2006, 73, No. 104209. (28) Tilocca, A.; Cormack, A. N.; de Leeuw, N. H. The Structure of Bioactive Silicate Glasses: New Insight from Molecular Dynamics Simulations. Chem. Mater. 2007, 19, 95−103. (29) Tilocca, A. Molecular Dynamics Simulations of a Bioactive Glass Nanoparticle. J. Mater. Chem. 2011, 21, 12660−12667. (30) Tilocca, A. Cooling Rate and Size Effects on the Medium-Range Structure of Multicomponent Oxide Glasses Simulated by Molecular Dynamics. J. Chem. Phys. 2013, 139, No. 114501. (31) Goel, A.; Kapoor, S.; Tilocca, A.; Rajagopal, R. R.; Ferreira, J. M. F. Structural Role of Zinc in Biodegradation of Alkali-Free Bioactive Glasses. J. Mater. Chem. B 2013, 1, 3073−3082. (32) Malavasi, G.; Pedone, A.; Menziani, M. C. Study of the Structural Role of Gallium and Aluminum in 45S5 Bioactive Glasses by Molecular Dynamics Simulations. J. Phys. Chem. B 2013, 117, 4142− 4150. (33) Malik, J.; Tilocca, A. Hydration Effects on the Structural and Vibrational Properties of Yttrium Aluminosilicate Glasses for in Situ Radiotherapy. J. Phys. Chem. B 2013, 117, 14518−14528. (34) Baram, P. S.; Parker, S. C. Atomistic Simulation of Hydroxide Ions in Inorganic Solids. Philos. Mag. B 1996, 73, 49−58. (35) Chizallet, C.; Raybaud, P. Pseudo-Bridging Silanols as Versatile Brønsted Acid Sites of Amorphous Aluminosilicate Surfaces. Angew. Chem., Int. Ed. 2009, 48, 2891−2893.
ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.5b07804. One figure showing average Qn speciation and network connectivity vs OH content, calculated from two independent samples for each composition, to highlight the size of error bars (PDF)
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AUTHOR INFORMATION
Notes
The authors declare no competing financial interest.
ACKNOWLEDGMENTS I gratefully acknowledge the U.K. Royal Society for financial support (RS-URF). REFERENCES
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