Nuclear Glass Durability: New Insight into Alteration Layer Properties

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Nuclear Glass Durability: New Insight into Alteration Layer Properties Stephane Gin,*,† Claire Guittonneau,† Nicole Godon,† Delphine Neff,‡ Diane Rebiscoul,† Martiane Cabie,§ and Smail Mostefaoui^ †

CEA Marcoule DTCD SECM LCLT, F-30207 Bagnols Sur Ceze, France CEA Saclay, LAPA, SIS2M UMR 3299, F-91191 Gif Sur Yvette, France § Universite d’Aix Marseille III, CP2M, F-13397 Marseille, France ^ Museum National d’Histoire Naturelle, CNRS, LMCM, UMR 7202, F-75005 Paris, France ‡

ABSTRACT: We have performed TEM, Raman microspectroscopy, and NanoSIMS characterization of borosilicate glass samples altered for nearly 26 years at 90 C in a confined granitic medium in order to better understand the rate-limiting mechanisms under conditions representative of a deep geological repository for vitrified radioactive waste. For the first time, we show a thick interphase that behaves like a diffusion barrier between the pristine glass and the other alteration products (porous gel and crystalline phases). Our findings indicate that the glass undergoes two distinct irreversible reactions: (i) hydration of the pristine glass controlled by water diffusion with a diffusion coefficient of 2 10 21 m2/s and (ii) transformation of the hydrated glass into a macroporous gel with major structural changes. Both materials are nonstoichiometric and metastable. A final reversible reaction leads to the formation of crystalline phases that consume elements forming the gel layer and the hydrated glass. All these reactions must be combined in a model to predict long-term rates of nuclear glass in natural environments.

’ INTRODUCTION For decades, borosilicate glass has been used in several countries (France, Belgium, England, United States) to confine high-level radioactive waste remaining after reprocessing spent nuclear fuel. Management of such materials is a key issue as safety must be guaranteed over a geological time scale, typically up to a million years. The most consensual way to achieve such a goal is to place the glass canisters in a deep geological repository having suitable characteristics: a sufficiently wide and thick homogeneous and low-water-permeable host rock, a low probability of seismic events, a location below the water table, etc. The main long-term risk is contamination of the biosphere by long-lived radionuclides (i) released from the glass due to its alteration by water and (ii) dispersed due to the water mobility in the geosphere. Borosilicate glass alteration has been extensively studied since the beginning of the 1980s, leading to a generally accepted understanding of the basic processes,1 although some of them remain open to debate, namely, the rate-limiting step depending on the reaction progress.1,2 Moreover, most of the mechanisms involved appear to be similar for all silicate glasses, including basaltic glass, obsidian, nuclear borosilicate glass, and soda-lime glass.3 10 In conditions relevant to a geological disposal environment, four types of reactions appear to occur simultaneously: (i) ion exchange, also called interdiffusion, corresponding to a preferential exchange between alkalis from the glass and hydrogen species from solution;11 14 (ii) hydrolysis by OH of ionic-covalent bonds in the silicate network (such a reaction can be catalyzed or inhibited by other aqueous species);15 18 (iii) condensation of sparingly soluble species r 2011 American Chemical Society

(mainly Si, Al, Ca), leading to steady-state concentrations of these elements in solution and to the formation of a low-porosity, amorphous hydrated material called a gel or hydrated glass depending on the authors;6,19 22 and (iv) precipitation of crystalline phases, mostly smectite-type clay minerals, rare earth phosphates, and sometimes zeolites.23,24 Throughout the community working on silicate water interactions, the role of the alteration layer remains at the center of the debate: the literature reports evidence of protective25 29 or nonprotective30 32 properties of this interphase depending on the nature of the silicate and the alteration conditions, but no general theory exists to predict a priori such behaviors. Only the consequence of this effect via an apparent diffusion coefficient can be modeled. Yet, it is of primary importance to better understand this phenomenology because the existence of protective coatings can drastically modify the way to model the alteration kinetics and mass balance. To understand how the alteration layer may affect the glass dissolution rate, we have selected samples from an experiment lasting nearly 26 years carried out on the French reference SON68 borosilicate glass altered at 90 C in a slowly renewed granitic solution in the presence of sand and fragments of granite and corrosion products. These samples were altered at a mean rate of 2 3 orders of magnitude below the forward rate. When selecting them, we hoped that these very rare samples (age and Received: June 11, 2011 Revised: July 28, 2011 Published: August 15, 2011 18696

dx.doi.org/10.1021/jp205477q | J. Phys. Chem. C 2011, 115, 18696–18706

The Journal of Physical Chemistry C

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Table 1. Chemical Composition of SON68 Glass (in Oxide wt %) SiO2

Al2O3

B2O3

Na2O

CaO

ZnO

Li2O

Fe2O3

P2O5

NiO

Cr2O3

ZrO2

Cs2O

SrO

Y2O3

45.48

4.91

14.02

9.86

4.04

2.50

1.98

2.91

0.28

0.74

0.51

2.65

1.42

0.33

0.20

MoO3

MnO2

CoO

Ag2O

CdO

SnO2

Sb2O3

TeO2

BaO

La2O3

Ce2O3

Pr2O3

Nd2O3

UO2

ThO2

1.70

0.72

0.12

0.03

0.03

0.02

0.01

0.23

0.60

0.90

0.93

0.44

1.59

0.52

0.33

Figure 1. Experimental setup.

leaching conditions) would favor identification of the key processes controlling the glass dissolution kinetics in a confined medium. The morphology, chemistry, and microstructure of the alteration layer have been investigated by transmission electron microscopy (TEM), Raman microspectroscopy, and nanoscale secondary ion mass spectroscopy (NanoSIMS).

’ MATERIALS AND METHODS Glass. This study was conducted on SON68 glass, the inactive surrogate of the French R7T7 glass produced at La Hague, where spent nuclear fuel in France is reprocessed and the waste vitrified.33 Its composition is shown in Table 1. Leaching Experiment. A cylindrical SON68 glass block 70 mm in diameter and 80 mm high was leached for 25.75 years in slowly renewed “granitic groundwater” (at 90 C, 100 bar) in contact with sand and some pieces of granite and steel as repository environmental materials (Figure 2). A detailed description of the experiment is given in a previous paper.34 The granitic groundwater was, in fact, a commercial mineral water, called Volvic; its composition has changed a little during the test. The main constituents of this solution are SiO2 (23.0 31.7 mg 3 L 1), Ca (9.8 11.5 mg 3 L 1), Na (9.2 11.6 mg 3 L 1), Mg (5.4 8.0 mg 3 L 1), K (5.5 6.2 mg 3 L 1), Cl (7.0 13.5 mg 3 L 1), SO42 (7.2 8.1 mg 3 L 1), HCO3 (65.9 71 mg 3 L 1), and NO3 (1.0 6.3 mg 3 L 1), and the pH ranged between 7.0 and 7.2. During the test, 163 solution samples were used to calculate from the boron release the mean altered glass thickness (28 ( 9 μm) and the glass dissolution rate.34 For the last 24 years, the mean alteration rate has remained very constant at 6 10 3 g 3 m 2 3 d 1, about 20 times higher than the residual rate (i.e., the minimum and virtually constant rate reached once the solution is saturated with respect to silicon) measured in a batch reactor at the same temperature and about 200 times lower than the initial rate. At the end of the experiment, SEM analyses were performed on many specimens sampled from the entire glass block. Surprisingly, the glass alteration layer has neither a uniform

thickness nor a homogeneous morphology. The location of the sampling valve (at midheight of the glass block) appears to “divide” the glass block into two parts. In the upper half (above the sampling valve), the general morphology of the alteration layer consists of a relatively uniform gel and some crystalline phases, which are rare-earth phosphates; the mean measured thickness of this alteration layer is 7 μm. In the lower half of the glass block, however, the alteration layer is more than 10 times thicker. It has been shown that these differences were likely due to chemical gradients in the solution resulting from low hydrodynamic transport of aqueous species.34 Glass Samples. Of the many specimens obtained from the large glass sample, two are taken as representative of the upper and lower parts (Figure 1). Sample 2 is considered interesting because it was altered at a rate 20 times higher than Sample 1. The idea motivating the analysis of this sample was to attempt the generalization of the concepts. Focused Ion Beam Milling. Focused ion beam (FIB) milling was performed with an FEI 200 TEM FIB system at the University of Aix-Marseille III. We used the same procedure as described in Benzerara et al.35 The same area as observed by SEM was located by means of FIB imaging capabilities. The FIB cross section was prepared with a 30 kV Ga+ beam operating at ∼20 nA. The prepared cross section measuring approximately 12 μm 5 μm 0.15 μm was transferred at atmospheric pressure with a micromanipulator to the membrane of a carboncoated 200 mesh copper grid. Three cross sections were prepared by this technique (Figure 2). The first included the entire alteration film belonging to sample 1. The second included the pristine glass/gel interface belonging to sample 2. The third was taken near the second (on sample 2) at the gel/crystalline phase interface. TEM and EDX Analyses. Two transmission electron microscopes were used in the analyses of the two specimens. The morphological observations were made on a JEOL JEM 2010F microscope operating at 200 kV, equipped with a field emission gun, a high-resolution UHR pole piece, and a Gatan GIF200 energy filter. Energy-dispersive X-ray (EDX) analyses were carried out on a Technai G2 (FEI), equipped with a LaB6 source operating at 200 kV. The detectors were a Gatan CCD camera, a STEM BF-DF detector, and an EDAX Genesis for the EDX analyses. The spatial resolution was 0.27 nm. Both instruments were used to detect crystals by an electron diffraction technique. Raman Spectroscopy. Raman microspectroscopy analyses were performed at room temperature on a Renishaw Invia spectrometer equipped with a double Nd:YAG laser (532 nm) and a 2400 grating. Analyses were performed under a Leica 100/0.90 microscope objective working with a beam size of 1.5 μm in diameter. The spectral resolution was better than 2 cm 1. Each Raman spectrum was obtained with a total acquisition time of 500 s (10 s per spectrum and 50 successive acquisitions to reduce spurious noise). A linear baseline was subtracted from the Raman spectra with LabSpec software. 18697

dx.doi.org/10.1021/jp205477q |J. Phys. Chem. C 2011, 115, 18696–18706


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Figure 2. Location of the samples on the glass block and SEM images showing the location of the three cross sections prepared by FIB (outlined in red).

NanoSIMS. Chemical maps were performed on sample 1 with a Cameca NanoSIMS 50 at the Museum National d’Histoire Naturelle in Paris, France. First, under a O primary beam, focused to a spot size of 350 nm, secondary ions were sputtered from the sample surface and 11B+, 27Al+, 39K+, 90Zr+, 95 Mo+, 28 + Si , and 40Ca+ were detected in combined mode (i.e., cycling simultaneously over two B-fields, one for the first five species and the second for the last two) in electron multipliers. Second, with the same O source, a second set of cations was detected: 1H+, 6 + 23 Li , Na+, 27Al+, and 28Si+ (also in combined mode, with 1H+ species alone, followed by the other four simultaneously). Third, under a Cs+ primary beam, focused to a spot size to 80 nm, the following secondary ions were analyzed: 1H , 23Na , 28Si , and 11 B 16O2 (1H species first, followed by the other three simultaneously). The map size was 10 μm 10 μm (2562 pixel images) or 20 μm 20 μm (5122 pixel images). Maps were obtained from a presputtered area by stepping the primary beam across the sample surface. Presputtering was done to clean the surface of contaminants before analysis, implant O or Cs+ ions in the material to be analyzed, and reach a steady state of ion emission. Analyses were made with the O primary beam at 8 pA or with the Cs+ primary beam at ∼0.5 pA. For chemical profiles, the counts for the detected elements were normalized to the 28Si counts. For each element i, calibration between the measured i/28Si ratio from NanoSIMS and the actual i/28Si value for the standard NBS610 and for the pristine SON68 were used to correct the i/28Si ratios measured on the gel/glass interface. The width of B profiles within the hydrated glass were derived directly from the B (for O source) or BO2 (for Cs source) maps, avoiding errors due to systematic shift between boron and silicon profiles (50 100 nm). The true width of the boron profile was calculated by subtracting the width deduced from the NanoSIMS profile (from 16% to 84% of the maximum intensity) and the spot size. The relative error on the true width is estimated at 15%. NanoSIMS results were processed with L'IMAGE software developed by L. Nittler, Carnegie Institution of Washington, Washington DC.

’ RESULTS TEM Characterization of Sample 1. On the basis of the morphological aspect, the alteration film of this sample can be

divided into three distinct layers. From the pristine glass to the bulk solution, we observe a microporous layer about 1.5 μm thick (Figure 3A), followed by a 10 μm thick macroporous layer with pores ranging from 10 to 170 nm in diameter, and an external layer of crystalline phases of about 0.3 μm thick. For practical reasons, these three layers are referred to here as the hydrated glass layer (HGL), the gel layer (GL), and the crystalline phase layer (CPL), respectively. The three regions are clearly separated by rough, but welldefined, interfaces (Figure 3A C). The HGL is characterized by depletion of Na and Zn and enrichment of K provided by the synthetic groundwater (spot analyses not shown here). TEM does not provide any information about the behavior of the other mobile elements of the glass, such as B and Li (generally used as glass dissolution tracers), but NanoSIMS data presented in the corresponding section indicate that they are also partly removed from the HGL. Moreover, Raman spectroscopy confirms the presence of water in the HGL (see the related section). The pore size cannot be estimated in the HGL, probably because the FIB cross section is too thick. A line of pores is visible at the interface between the HGL and the pristine glass (Figure 3B). The GL exhibits large pores except in the first hundred nanometers (Figure 3C), confirming substantial microstructural modifications compared with the HGL. Such a porous structure is consistent with in situ reactions of hydrolysis and condensation of silicon.22,36,37 The GL is partly depleted in Na, Mo, and in Ni in comparison with the pristine glass and contains some Cl and K provided by the leaching solution. The GL near the CPL contains some rounded crystallized grains 30 50 nm in diameter (Figure 4A) that are rich in Ag and Te and contain S (Figure 4B). The presence of O, Si, Al, and Na in the nanoparticles is unexpected. The analysis probably included elements from the surrounding gel. Inter-reticular distances have been deduced from the diffraction pattern: d1 = 0.382 nm, d2 = 0.323 nm, d3 = 0.235 nm. On the basis of, in part, other diffraction patterns not shown here, these nanoparticles may consist of tellurium sulfide (Te4.08S27.92), tellurium oxide sulfide ((TeO2)SO4), and silver sulfide (Ag2S), and probably other phases. phyllosilicates The CPL is mainly composed of 12 15 Å bearing some crystallized nanoparticles (Figure 4C). Figure 4C particularly shows the clayey minerals mainly composed of Si, 18698



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Figure 3. TEM bright-field images of sample 1. (A) Whole alteration film. (B) Image showing the pristine glass/hydrated layer interface. (C) Image showing the hydrated layer/gel interface.

Na, Zn, Fe, Al, Mg, K, Ni, and Ca. Because the state of hydration of these minerals cannot be complete due to the sample preparation, their identification is not possible, but the inter-reticular typically corresponds to 2:1 clay minerals. This distance of 15 Å observation is consistent with the literature.23,38 40 The CPL also contains some dense crystalline phases rich in Ca, P, and rare earth elements (REE) (Figure 4D F), corresponding to previously described apatitic minerals.41,42 TEM Characterization of Sample 2. The second and third FIB cross sections on sample 2 are located at the interface between the pristine glass and the HGL and in the CPL, respectively (Figure 2). They globally show the same type of layers and the same microstructures as those described on the previous sample, but with some interesting differences. First, although the whole alteration layer is around 20 times thicker than that in sample 1, the HGL is only half as thick (700 nm versus 1.4 μm). Moreover, K appears to be absent in both this layer and the GL. Second, HRTEM images show the presence of small clay crystals homogeneously distributed within the amorphous GL (Figure 5A). This means that clay particles can precipitate throughout the GL, whereas shorter laboratory experiments generally indicated that clay minerals only precipitated outside the gel. In that sense, this microstructure strongly resembles naturally weathered old basaltic glasses.23 Third, the CPL is made of a diffuse inner clay layer and a denser outer layer (Figure 5B). These clays also bear some crystallized nanoparticles (Figure 5E), which have not been analyzed. Figure 5D shows the sheets of the diffuse clays; the inter-reticular distance is close to 10 Å. This last parameter remains unknown for the dense clays as their structure is damaged under the electron beam. We noted little chemical difference between the two clay layers; the denser one contains less Ca and Fe and more Zn, Na, and Mg than the diffuse clays. Raman Microspectroscopy. Raman microspectra have been obtained for samples 1 and 2. Here, we show only data from

sample 1, as both lead to the same results. Spots of about 1.5 μm are plotted from the gel layer (spectra 2 6) to the pristine glass and highlight structural change from the pristine glass to the gel layer. Spectrum 6 is likely related to the hydrated glass (see justification below). The region between 800 and 1200 cm 1 mainly contains data on Si O stretching bonds (Figure 6B), including at least four different contributions attributed to tetrahedral species Qn (n refers to the number of bridging oxygen atoms bonded to a given silicon atom).43 45 The spectra obtained appear to comprise 12 contributions at wave numbers 825, 852, 888, 920, 950, 976, 1030 (between 1020 and 1040), 1047, 1076, 1120, 1136, and 1148 cm 1. These data are not the result of peak deconvolution; they correspond to the peak positions observed on the spectra. According to Li et al.,46 the bands observed in borosilicate glass at 852, 920, and 976 cm 1 are Si O stretching vibrations of Q0, Q1, and Q2 species, respectively. The contribution at 976 cm 1 probably also contains the Si OH stretching bands observed in hydrated glass.47,48 The contribution at 1030 cm 1 may be due to Si O stretching vibrations of Q3 species. Li et al.46 situate this contribution at 1050 1100 cm 1, but McKeown et al.49 assign two bands to Q3 units at 1007 and 1101 cm 1.49 We can, therefore, assume that the band between 1020 and 1040 cm 1 is due to Q3 species. The bands at 1120 and 1148 cm 1 may be Si O vibrations of Q4 species. Moreover, the 950 and 1136 cm 1 may be due to stretching of B O Si bonds.50,51 The shoulder observed at 1047 cm 1 in the spectrum for point 6, and the band that appears in the gel at 1076 cm 1, could be due to Si O stretching, as observed in hydrated glass. Both of these contributions have been observed in the hydrated layer of an archeological glass specimen at 1023 and 1079 cm 1,52 but the frequency of these bands is known to depend on the glass composition.53 Two bands have not yet been assigned: the shoulder at 825 cm 1, which increases in the gel, and the band near 18699



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Figure 4. TEM analyses of the crystalline phases and the gel near them on sample 1. (A) TEM image of crystallized grains in the gel, near the crystalline phases, and electron diffraction pattern (inset) of a grain. (B) EDX spectrum of a crystallized grain. (C, D) TEM images of crystalline phases. (E) Zoomin on the dark zone of these phases. (F) Electron diffraction pattern on zone E showing that they are crystallized phases.

888 cm 1, which is of much lower intensity in the gel than in the glass. The main gel characteristics observed in this spectrum band assignment are the decrease in B O Si bonds and Q3 species and the increase in Q0, Q2, and Q4 species compared with the pristine glass. First of all, the decrease in B O Si bonds is consistent with NanoSIMS chemical analysis (see below). The presence of Q2 species and the decrease in Q3 species has already been observed by Raman spectroscopy of archeological glass altered 18 centuries in seawater.52 According to Mysen and Richet,54 alkali metals are generally bound to nonbridging oxygen atoms of Q3 species and alkaline earth metals are generally bound to nonbridging oxygen atoms of Q2 species. It is thus logical to observe more Q2 species in the gel, as Ca2+ is

better retained in the gel than the alkali metals (see Ca and Li behavior in Figure 7). Concerning Q4 species, the Raman spectra reveal two contributions in the gel above 1100 cm 1, both of which increase compared with the pristine glass. The increase in Q4 species is consistent with repolymerization of the network during gel formation.36 Spectrum 6 was recorded between the pristine glass and the gel and could, therefore, characterize the hydrated glass zone, although caution is necessary because the hydrated glass cannot be precisely located by optical microscopy and the beam is the same size as the hydrated glass layer. Nevertheless, Figure 6A showing the Raman bands in the vibration zone of water molecules and OH groups44,45,55 clearly reveals the presence of water in the probed zone, confirming this result. The bands 18700



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Figure 6. Raman spectra carried out on sample 1. (A) Raman spectra of the five points analyzed at the gel/glass interface and the pristine glass analyzed far from the interface, in the range of 1500 1700 cm 1; the spectrum for point 2 is not plotted on the same graph as the others because of the strong fluorescence beyond 1200 cm 1, which appears to enhance the intensity of the bands after baseline correction. (B) Raman spectra of the same six points, in the range of 800 1200 cm 1. Figure 5. TEM bright-field images of the gel and the crystalline phases on sample 2. (A) The gel near the hydrated glass layer, bearing some thin clayey phases. (B) The clayey phases (diffuse clays and dense clays). (C) The gel with some crystallized grains (grain 1: Ag, Te, S; grain 2: Cd, S, K) and the diffuse clays. (D) Sheets of the diffuse clays. (E) Dense clays with a crystallized grain, and the electron diffraction pattern (inset). (F) Electron diffraction pattern of the dense clays.

observed show two main contributions: one at 1576 cm 1 due to O H stretching and another at 1629 cm 1 due to the deformation of water molecules (both free water and bound water in silanol groups).56 In the 800 1200 cm 1 spectrum band, the characteristics of point 6 clearly resemble those of the pristine glass more than the gel. Note the decrease in the intensity of the characteristic bands of Q3 species and B O Si bonds, which is consistent with the chemical analysis findings (see below). This observation suggests preferential hydrolysis of Si O Alk bonds (Alk = Na, Li, Cs), followed by subsequent repolymerization of silicon in the gel. NanoSIMS Characterization. NanoSIMS analyses were carried out on sample 1 in an area where the alteration film was about 10 μm thick. The element distribution maps and profiles obtained with the O primary ion beam produced the following results (Figure 7): (i) B and Li exhibited similar behavior, that is, depletion along a diffusion profile in the hydrated glass and complete depletion in the gel. (ii) The profiles for B (supplied only by the glass) and K (supplied only by the leachate) are completely opposed, delimiting the zone of glass hydration and interdiffusion.

(iii) The mean width of the diffusion profile of boron has been estimated from four measurements: 350 ((150) nm. The latter uncertainty takes into account the standard deviation of the results and the relative error on the calculation. Note that the actual profile does not correspond to the one shown in Figure 7 (see the Materials and Methods). (iv) Mo also follows a diffusive profile, but delayed with respect to the most highly mobile elements. (v) The Zr/Si ratio is constant in the gel. As Zr is immobile in most of the glass alteration processes,27,57 this result shows that there is no Si gradient in the gel. Only the crystallized phases are slightly depleted in Si compared with the pristine glass. (vi) Al and Ca show slight and progressive enrichment between the hydrated glass and crystalline phases. (vii) Na is strongly present in the gel. Its presence is related to the fact that it is also found in the leaching solution, making exchanges relatively likely. NanoSIMS analyses with the Cs+ primary ion beam allow precise examination of the hydrated glass zone because this source provides higher spatial resolution. However, light elements, such as B, Li, and Na, are poorly ionized by this source. Consequently, such analyses took more time and some elements could not be analyzed. We obtained profiles of boron ionized as BO2 . Three element distribution maps for sample 1 yielded a B diffusion profile to a depth of 170 ((30) nm (Figure 8). This value is half the mean value found with the O source, but it remains close, taking into account the uncertainties, and is high 18701



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Figure 7. NanoSIMS analyses on sample 1, performed with the O primary beam. Images (20 μm 20 μm) are normalized to 28Si. The element profiles normalized to 28Si are made from the pristine glass to the gel and to the crystalline phases, and are shown on a logarithmic scale. The small graph shows the 11B and the 39K profiles, normalized to 90Zr, from the pristine glass to the gel on the same zone as the other profiles.

Figure 8. NanoSIMS analysis on sample 1, made with the Cs+ primary beam, showing the BO2 image (10 μm 10 μm), normalized to Si (on the left), and the BO2 profile made from this image (on the right). 18702


The Journal of Physical Chemistry C enough compared with the Cs+ beam size (80 nm) to confirm without ambiguity the diffusive behavior of boron at the glass/ alteration layer interface. Because this analysis is more precise than the one based on the O source, we can conclude that the boron diffusion layer is about 200 nm thick. This result is important as it allows locating, for the first time, the diffusionlimiting step of glass alteration under residual rate conditions. It is shown by this technique that only a small and internal part of the HGL exhibits diffusion. Moreover, contrary to the simple glasses studied by Cailleteau,27 a complex nuclear glass, such as SON68, does not form a dense layer with closed porosity within the gel.

’ DISCUSSION From the Characterization Data to the Most Likely Reaction Pathway. Transmission electron microscopy (TEM) ob-

servation of FIB cross sections shows that the alteration products formed during SON68 glass leaching at a high reaction progress are morphologically layered. From the pristine glass to the bulk solution, three distinct layers have been characterized: a dense hydrated material whose morphology resembles that of the bulk glass, a macroporous gel layer that may include small phyllosilicate crystals, and an external layer of crystalline phases made of 2:1 clay minerals, apatite rich in REE, and some nanoparticles trapped in the clays and in the external part of the gel. This general pattern is qualitatively independent of the position on the glass block, and thus of the glass dissolution rate (at least within the studied range of rates). Only the relative thicknesses of the three layers depend on the sample location. Morphological (TEM), chemical (TEM, NanoSIMS), and structural (Raman) data clearly indicate that the farther the alteration products are from the pristine glass, the more they differ from the glass characteristics. This observation is fully consistent with recent findings by Valle et al., who used isotopes as tracers to show that external phyllosilicates are in equilibrium with the bulk solution, whereas the porous gel layer has an isotopic signature intermediate between those of the glass and solution, meaning that the gel structure inherits some undissolved glass bonds and also undergoes some chemical and structural reorganization.20 Other work confirms that the gel properties are between those of a glass relict structure (dealkalinized glass) and those of an amorphous precipitate of sparingly soluble species.27,58 All the results presented here argue in favor of progressive transformations of the pristine glass into more stable phases (clays, apatite, etc.) with transient stages involving metastable materials (hydrated glass, porous gel). Most of these reactions are irreversible except the formation of crystalline phases (at least phyllosilicates) that should be in equilibrium with the bulk solution. We can finally propose a simplified general reaction pathway that is consistent with all the previous observations:

Process 1 results both from water diffusion through the glass and from leaching of the weakly bonded atoms, such as alkali metals and boron, based on the reasonable assumption that the

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boron diffusion front evidenced by NanoSIMS is close to that of water. This process conserves the glassy structure with only slight changes, as shown by Raman spectroscopy. This is qualitatively confirmed by ab initio calculations that indicate that some restructuring is possible, but details are not accessible to the interpretation of the Raman spectra due to the complexity of the glass.59 The boron gradient evidenced by NanoSIMS indicates that the ion-exchange reaction is a slow process, but faster than the hydrolysis of the silicate network in these leaching conditions. A diffusion coefficient can be derived from such profiles (see below). Process 2 corresponds to the irreversible transformation of the hydrated glass into a porous gel. NanoSIMS and TEM chemical analyses performed on sample 1 indicate that the concentrations of the main formers (Si, Al, Zr) remain stable along the gel profile. This is consistent with in situ reorganization, which, at a very high reaction progress, allows local precipitation of phyllosilicates, as evidenced on old natural glass samples.23 The transformation of the hydrated glass into porous gel also leads to the dissolution of the sparingly soluble elements that form the gel. Process 3 corresponds to the precipitation of crystalline phases. As a consequence of this process, gel-forming elements are consumed in stable phases. This has potentially several consequences: it would certainly modify the pore water composition, it would also modify the gel composition (e.g., zeolite precipitation is known to act as a Si and Al sink, leading to resumption of glass alteration60,61), and it may impact the stability of the hydrated glass. It has been shown that the growth of phyllosilicates at the surface of SON68 glass was limited by the availability of some transition metals present in small amounts in the glass, such as Ni and Zn.62 TEM analyses (not shown here) suggest that these elements are not highly retained in the amorphous alteration products. Moreover, Mg provided by the synthetic groundwater is known to favor process 3. Because phyllosilicates are open structures (very low protectiveness compared with the hydrated glass), one can imagine that their precipitation tends to sustain glass alteration. Rate-Limiting Step(s). Determining the step(s) controlling the glass dissolution rate is of primary importance in order to predict the long-term rate in geological disposal conditions. Assuming that the previous phenomena correspond to the most likely scenario of nuclear glass evolution in contact with water, one can wonder if one or more reactions control the overall dissolution phenomenon. First, consider that glass cannot be in equilibrium with the bulk solution even if the latter has reached apparent saturation.63,64 This means that the low dissolution rates observed in this test (3 10 3 r0 for sample 1 and 6 10 2 r0 for sample 2, where r0 is the forward rate) cannot be due only to an affinity effect (thermodynamic equilibrium between glass and solution). Second, a chemical gradient for boron is evidenced in the internal portion of the HGL. If diffusion is rate-limiting, it is necessarily in this area, as no chemical gradients and large pores have been found elsewhere in the alteration film. In this case, water diffusion through the pristine glass and/or diffusion of mobile elements replaced by water molecules or protons may be a rate-limiting step. This hypothesis would necessarily imply a glass dissolution rate decreasing with the inverse of the square root of time (at least decreasing with time). However, the rate remained very constant during the 26-year experiment.34 Consequently, this process cannot be retained as the sole phenomenon 18703


The Journal of Physical Chemistry C controlling the glass dissolution rate. If we now consider that the diffusion process is coupled with dissolution of the hydrated glass layer, then it becomes possible to derive a constant rate. The rate of process 2 should be strongly dependent on the chemical composition of the fluid at the reaction front. The fluid composition at this position may depend, in turn, on the chemical constraints of the bulk solution and probably on the confinement constraints in nanopores. As shown by McGrail and co-workers,9 the ion-exchange phenomenon (reaction involved in process 1) locally increases the pH, leading to gradients in chemistry in the alteration layer and acting as a driving force sustaining the hydrated glass dissolution. These authors also suggest that such a phenomenon could have a major effect on the long-term glass durability, mainly because the rate of ion exchange is faster relative to glass matrix dissolution once silica saturation conditions are reached. In our study, the existence of a thin diffusion zone within a thick hydrated glass is in agreement with this idea. As a result of all our observations, it is not possible to identify a single reaction that controls the glass dissolution process as it appears that the overall rate results from strong couplings between chemical reactions and the transport of aqueous species within the alteration layer. Only a model implementing these mechanisms would be able to predict the rate as a function of the glass composition, intrinsic properties, and chemical boundary conditions. Two semiempirical models, GM200165 and GRAAL,24,66 are currently able to simulate such processes, even simplified, with quite good results when applied to laboratory experiments1,67 or archeological analogs.68 What Is New? Compared with borosilicate alteration layers previously observed,26,42,69,70 the present study points out two major differences: the internal dense layer has never been clearly evidenced and characterized for this type of complex glass, and no external dense layer, such as the one found by Cailleteau et al. within a 5-oxide borosilicate alteration layer, has been observed in our study.27 To our knowledge, there is only one reference in the literature to this dense layer formed at a high reaction progress on nuclear glass.71 After 12.2 years of leaching at 90 C, a SON68 specimen exhibited an internal layer about 200 nm thick, called “unaltered glass” by the authors. Unfortunately, the microstructure and the properties of this material were not investigated in greater detail. This layer apparently corresponds to the hydrated glass in our study. Its lower thickness is likely due to the higher pH and a shorter duration. Both the GM2001 and the GRAAL models already take into account a layer in which water transport is limited. The existence of this layer at a high reaction progress was assumed based on indirect evidence at a lower reaction progress.72 74 Here, we show that this expected layer does exist and behaves like a diffusion barrier for mobile species. Assuming that boron behaves similarly to water, the apparent diffusion coefficient of water through the passivating layer can be derived from the NanoSIMS profile with the analytical version of the GRAAL model.66 In this case, it is equal to 2 10 21 m2/s. This value is in good agreement with data obtained by other indirect methods,14,74,75 showing that the glassy structure considerably slows down water motion compared with classical porous media. This is a key point regarding the long-term durability of this kind of material. Even though the resolution of Raman spectroscopy does not allow accurate structural characterization of the hydrated layer, our findings indicate that no major reorganization occurs within this layer. Such a layer, called a “passivating reactive interphase” by

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Frugier et al 3 24 in the GRAAL paradigm, is transformed more or less rapidly into a porous gel depending on the solution chemistry. The gel, in turn, partly dissolves, and crystalline phases precipitate. All these reactions take place at a near constant volume (isovolumetric reactions).34,76 Because of the limited number of oxides (SiO2, B2O3, Na2O, CaO, ZrO2) and the absence of Al in the glass studied by Cailleteau et al., no crystalline phases precipitated during leaching.27 Moreover, the amount of sparingly soluble elements was low compared to nuclear glass. These reasons can explain why silicon is more reactive and mobile in the alteration layer in the case of simple glass, leading to clogging of the porosity in the external part of the gel. This mechanism was expected by the authors to be responsible for the rate drop. In our opinion, this phenomenon is not a prerequisite for the rate drop, but rather a supplementary process improving the glass durability. This conclusion is supported by observations of simple glass alteration layers that do not necessarily form a nonporous (or closed porosity) layer within the alteration film. Consequently, processes 1 and 2 are sufficient for the rate of borosilicate glass to drop if silica saturation is reached in solution.

’ CONCLUSION We have shown that a detailed characterization of the alteration film obtained from an experiment of long duration conducted on a surrogate nuclear glass allows us to better understand the reaction pathways and the relative importance of the different mechanisms involved on the glass dissolution rate. Modeling these reactions remains difficult as most of the alteration products are transient and nonstoichiometric phases. Moreover, it has been shown that the leaching solution near the reactive interface is confined in nanopores and the composition of such a solution is probably influenced by ion exchange and in situ reactions of hydrolysis and condensation. The physics and chemistry of such a confined solution remain largely unknown even if recent studies suggest that water properties are far from those of the bulk.77 79 Consequently, in the future, we must pay greater attention to the phenomena occurring within the liquid and solid phase at the interface between the bulk glass and the bulk solution. ’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]

’ ACKNOWLEDGMENT We would like to thank the Museum National d’Histoire Naturelle for the NanoSIMS analyses (the national NanoSIMS facility at the Museum d’Histoire Naturelle was established by funds from the CNRS, the Ile de France region, the Ministry of Higher Education, and the Museum itself). We are also very grateful to Anne Michelin for her assistance with the Raman microspectroscopy. Finally, we would like to thank the two anonymous referees for their stimulating and very helpful comments. ’ REFERENCES (1) A critical evaluation of the dissolution mechanisms of high-level waste glasses in conditions of relevance for geological disposal (GLAMOR); European Commission report, EUR 23097, 2007. 18704


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Nuclear Glass Durability: New Insight into Alteration Layer Properties

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