# Interaction between carbon steel and low-pH bentonitic cement grout in anoxic, high temperature (80°C) and spatially heterogeneous conditions

J. Goethals, Laurent de Windt, Charles Wittebroodt, A. Abdelouas, Xavier de La Bernardie, Y. Morizet, B. Zajec, V. Detilleux

► To cite this version:

J. Goethals, Laurent de Windt, Charles Wittebroodt, A. Abdelouas, Xavier de La Bernardie, et al. Interaction between carbon steel and low-pH bentonitic cement grout in anoxic, high temperature (80°C) and spatially heterogeneous conditions. Corrosion Science, 2023, 211, pp.110852. 10.1016/j.corsci.2022.110852. irsn-03880498

HAL Id: irsn-03880498

https://asnr.hal.science/irsn-03880498v1

Submitted on 23 Jul 2025

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1 Interaction between carbon steel and low-pH bentonitic
cement grout in anoxic, high temperature (80°C) and spatially
3 heterogeneous conditions.

4

5 J.Goethals<sup>a</sup>, L. De Windt<sup>b</sup>, C. Wittebroodt<sup>c</sup>, A. Abdelouas<sup>a</sup>, X. de la Bernardie<sup>a</sup>, Y. Morizet<sup>d</sup>, B. Zajec<sup>e</sup>

6 and V. Detilleux<sup>f</sup>

7

8 <sup>a</sup> Laboratoire Subatech, UMR 6457, 4 rue Alfred Kastler, La Chantrerie CS 20722, 44307 Nantes CEDEX 3

9 <sup>b</sup> Mines Paristech, Centres de Géosciences, PSL Research University, 35 rue St-Honoré, 77300 Fontainebleau,
10 France

10

11 <sup>c</sup> Institut de Radioprotection et de Sûreté Nucléaire (IRSN)/LETIS, 31 avenue de la division Leclercq, BP 17, 92262

12 Fontenay-aux-Roses Cedex, France

13 <sup>d</sup> Laboratoire de Planétologie et de Géosciences (LPG) UMR 6112, 2 rue de la Houssinière - BP 92208

14 44322 Nantes Cedex 3, France

15 <sup>e</sup>Slovenian National Building and Civil Engineering Institute (ZAG), Dimičeva ulica 12, SI 1000 Ljubljana, Slovenia

16 <sup>f</sup> BEL V, 148 Rue Walcourt, 1070 Anderlecht, Belgium

17 Abstract

18 Anoxic corrosion processes impacting carbon steel samples for both intimate and imperfect metal/
19 low-pH bentonitic cement grout (BCG) contacts were studied at 80°C. SEM, XRD and μ-Raman
20 spectroscopy analyses supported by geochemical modelling allowed a full characterization of the
21 corrosion mechanism. Steel in direct contact with BCG suffered from localized corrosion (pits
22 measuring up to 160 μm in depth after 150 days of exposure) principally triggered by the presence of
23 HS-. Depending on silicate and sulphide ingress from the BCG towards the interface, neoformed
24 magnetite may be converted to iron silicates and/or iron sulphides.

25 **Keywords:** Anaerobic conditions, galvanic coupling, localized corrosion, low-pH cement, sulphides.

26 1 Introduction

27

28 In the scope of deep geologic disposal of nuclear radioactive waste, the National French Agency for
29 Nuclear Waste Management (Andra) envisions to dispose High Level Waste (HLW) in a deep
30 geological facility located 500 m below the surface, in the Callovo-Oxfordian claystone (COx)
formation (Cigeo project) [1]. According to the multi-barrier concept associated to such geological disposal [2], HLW would be trapped in a glass matrix quenched in a stainless steel container. This package would be then encapsulated into a cylindrical carbon steel overpack and inserted in a carbon steel casing tube within an horizontal micro-tunnel drilled in the highly impermeable argillaceous host rock. Ceramic pads would be fixed to the canisters to facilitate insertion in the steel casing and to allow the reversibility of nuclear waste disposal during the operation period (approximately 100 years). A bentonitic cement-based grout (BCG) would be injected between the steel casing and the host rock, in the volume left by the coring operation. A 2D scheme corresponding to this multi-barrier concept is displayed in Fig. 1 a. The selected low-pH BCG aims to (i) neutralize the acidity coming from pyrite oxidation in the COx [3] while preserving glass from detrimental highly alkaline media, (ii) passivate the steel casing and (iii) limit oxygen diffusion and microbial activity [4–8]. This BCG material has a low viscosity related to its very high water to cement (W/C) ratio and to lubricating phyllosilicate constituents. This latter property is of crucial importance to allow the filling of the technological gaps existing between the metallic casing and the host-rock, since the grout would be injected over a distance of 100 m [8]. Due to the long distance of injection, it is possible that part of the steel casing would not be fully covered by the BCG, inducing spatial heterogeneities that would be progressively filled by the host-rock porewater.

After the sealing of the HLW disposal cell, oxygen will be rapidly consumed [9] and the heat produced by radioactive decay would induce an increase of the temperature up to 90°C in the disposal cell [2]. The corrosion of low-carbon steels under anoxic and high temperature conditions representative of nuclear waste geological disposal has been widely studied. These studies deal with carbon steel/bentonite and carbon steel/clay rock interfaces (e.g. [10–19]) and, to a lesser extent, with carbon steel/high-pH cement materials interfaces (e.g. [20–22]). However, only few corrosion studies have been conducted with steel in direct contact with low-pH cement-based materials [4–7,23]. To this purpose, the BACUCE *in situ* experiments (Fig. 1 b) have been designed by the French Institute of Radioprotection and Nuclear Safety (IRSN) in the Tournemire underground research laboratory (URL). The URL is located in a tunnel that crosses a 250 m thick argillaceous formation of Toarcian and Domerian ages at Tournemire (Aveyron, France). The Upper Toarcian argillite presents geochemical and mineralogical properties relatively similar to the COx formation ones [24]. In complement to these field experiments, laboratory mock-up experiments have been carried out to better constraint the processes that would occur *in situ*. Both *in situ* and mock-up experiments have been performed with the same low-pH BCG under anoxic conditions, a constant temperature of 80°C and fully water-saturated conditions. The aim of these experiments is to consistently assess (i) the geochemical, mineralogical and microstructural change of the BCG, (ii) the corrosion rate and
products of low carbon steel in direct contact with BCG or in contact with the solution filling the
voids of the spatial heterogeneities. The present paper deals with the mock-up experiments
conducted within a period of one year. A combination of characterization tools (SEM, XRD and µ-
Raman spectroscopy) associated with geochemical modelling were employed to determine the
corrosion mechanism. The corrosion rates were estimated by weight loss and optical measurements.

70

## 2 Materials and Methods

### 2.1 Materials and porewater solution (PWS)

#### 2.1.1 Low carbon steel

TU37b (A37) steel composition is given in Fig. 2 a. This steel was selected for the BACUCE experiments due to its chemical similarity with the alloy selected by ANDRA for the steel casing (i.e. API 5L X65 steel). This steel exhibits a ferrite-pearlite layered structure (Fig. 2 a) and its external mill scale layer has a heterogeneous microstructure and presents “indentations” of several tens of micrometres (Fig. 2 b). This mill scale layer is primarily composed of magnetite (μ-Raman spectroscopy analysis, not shown).

#### 2.1.2 Bentonitic cement-based grout (BCG)

Table 1 gives the composition of the BCG. The cement that has been used is a blended CEM III/A Rombas cement composed of clinker/anhydrite (34 wt.%) and blast furnace slag (66 wt.%).

Compared to the cementitious component used in a similar previous study (i.e. CEM III/C Rombas) [6,7], this cement contains a smaller blast furnace slag fraction (which has a slow hydration kinetic) and, consequently, less sulphide ions ($S^{2-} \approx 0.55 \text{ wt. %}$), prone to induce detrimental effect on steel corrosion processes [5–7]. Bentonite and hydrotalcite have mainly a rheological role. Silica fume reacts with the CEM III/A hydrated phases to prevent any portlandite formation and high pH conditions by the input of silica to form C-S-H phases with a low Ca/Si ratio. The water on cement ratio (W/C = 1.7) is high enough to ensure fluidity during the injection and a good level of hydration.

To prepare the BCG, the different components were ground in an agate mortar and sieved with a 200 µm mesh grid. Then they were weighted in appropriate amounts and thoroughly mixed together. The final step was the addition of an appropriate volume of degassed (under vacuum) tap-water and mixing/stirring, all done inside an anoxic glovebox. XRD analyses associated to each of the components are given in Fig. A-1.
### 2.1.3 Synthetic porewater solution (PWS)

The synthetic solution was assumed to result from the mixing of the hydrated BCG porewater (75 vol.%) and the Upper Toarcian argillite porewater (25 vol.%). This mixed porewater solution will be denoted as PWS throughout the paper and its chemical composition is given in Table 2. It is worth mentioning that SiO₂ was not added in the solution to avoid its reaction with Ca and Mg and the formation of C-S-H and M-S-H, respectively. The solution was prepared at 20°C in an anoxic glovebox with degassed milliQ® water. The chemical composition and the pH of the solution were controlled after stirring the solution at 20°C by means of ionic chromatography and pH electrode, respectively. These measurements were also performed after heating the solution to 80°C in an anoxic oven and no significant change in the chemical composition were observed.

## 2.2 Experimental set-up

Two types of experiment were designed to study both the corrosion of steel, either in direct contact with BCG or in contact with PWS, and the BCG mineralogical evolution with time at 80°C in anoxic conditions. Experiment 1-R was dedicated to (i) study the corrosion of steel, either in direct contact with BCG, or in the presence of a spatial heterogeneity, (ii) assess the effect of the presence of a mill scale layer and of the steel surface finish on the corrosion rates, (iii) assess the solution composition and pH change with time. For this purpose, TU37b steel rods of 3.1×1.5×1cm³ were cut from a TU37b tube section. Of the six exposed surfaces of the rod, three were ground with 180 SiC grit paper (top, bottom and left sides), one was polished with diamond paste down to 3µm and on the two other sides, the pristine mill scale layers were preserved (Fig. 3 b). BCG was prepared in an anoxic glove box by mixing the component presented in Table 1 with a spoon and was directly poured in the PTFE reactors so that half of the rod height was embedded in BCG (Fig. 3). After this, BCG was let to harden during 30 days in the closed PTFE reactors prior to their partial filling with the PWS (Table 2). Despite the fact that the samples were prepared in an anoxic glove box, one can note that oxic corrosion processes impacted to some extent the unembedded steel surfaces during the initial phase of BCG paste hardening. As it was possible to retrieve the steel rods from the BCG, grinding and polishing operations were repeated to remove the corrosion products and the sample were replaced in the material prior to launch the experiment. Each of the six prepared samples were put in a different PTFE reactor, filled with 18 mL of PWS (Table 2) and transferred in an oven under argon atmosphere at 80°C.

The experiment 2-C was dedicated to investigate (i) the corrosion of steel in the presence of the BCG and (ii) to follow the mineralogical evolution of the BCG with time at 80°C. For this purpose, TU37b steel coupons of 1×1.5×0.2 cm³ were cut so that the 1×1.5 cm² surface was perpendicular to the
rolling direction of a TU37b tube section. BCG disks of (φ) 27 mm × (h) 4 mm were cut from BCG cylinders prepared in a similar manner as the one produced for experiment 1-R (Fig. 3 c). A specific in-house tool was used to machine the disks of the desired dimensions after 32 days of cure. The disks were put at the bottom of PTFE reactors, and the steel coupons were inserted in a PTFE basket located above the disks, in order to avoid the contact of steel with the disks or the PTFE reactor. As for experiment 1-R, the 6 reactors, each containing one coupon and one cement-grout disk, were filled with 6.6 mL of the porewater solution and put in an oven under argon atmosphere at 80°C. For both experiments, samples were periodically retrieved and characterized following the different methods presented in the next sections.

## 2.3 Sample preparation for micro characterization

The steel BCG assembly from experiment 1-R were first transferred in an anoxic glove box and immersed in deaerated ethanol (99%) for 4h to stop corrosion. Then, the samples were dried at least 12 h inside the glove box using silica gel prior to be embedded in Struers® Epoxy resin. After this, slices of the BCG covered and uncovered steel sample were cut perpendicularly to the rod axis with a diamond wire saw and ethanol as lubricant. Samples containing BCG were then immediately embedded with a more viscous resin (ESCIL® araldite) under vacuum, in order to impregnate the BCG material and to strengthen its mechanical resistance. The samples were then polished with SiC grit paper and diamond paste, down to 1µm. As the sample preparation were performed in aerated conditions, the polishing operation was done just before the analyses and the samples were transported to the different analysing facilities in an argon-containing sealed box.

## 2.4 Solution analyses

Porewater solutions coming from the experiment 1-R were analysed to monitor their chemical evolution with time. Periodically, reactors were retrieved from the oven, put in an in-house isothermal container to keep the solutions at 80°C and transferred in an anoxic glove box. For pH measurements, a fraction of the solution was transferred in a thermostated bath at 80°C and measurements were made with a pH electrode (Mettler-Toledo) freshly calibrated with IUPAC pH buffer (9.188 and 12.46). The other fraction of the solutions was filtered with 0.2 µm PTFE Millipore® membrane filter prior to their analyses. Major cations (Na, K, Ca, Mg, Sr,) and anions (Cl, F, SO₄, NO₃) were analysed by means of ionic chromatography with a Metrohm 850 Professional IC. Total inorganic carbon (TIC) was measured by means of a carbon analyser Analytikjena 2100S. ICP-MS analyses of Fe, Al and Si were carried out with a Q (quadrupole)-ICP-153 MS (Thermo Scientific X Series 2).
160 2.5 Solid characterization methods

161 2.5.1 Scanning electron microscopy (SEM)

162 Scanning Electron Microscopy coupled with Energy Dispersive X-ray spectroscopy (EDX) was used to
163 conduct analyses on the embedded polished samples coming from experiment 1-R as well as on one
164 steel coupon coming from experiment 2-C (28 days of interaction with PWS) kept for the surface
165 analyses of the corrosion products. SEM-EDX analyses were performed with a JEOL 5800 LV
166 microscope operating at 15 kV and 0.5 nA. Acquisition time of the EDX measurements was fixed to 60
167 s. Oxygen was measured independently from the other elements due to the presence of non-oxide
168 minerals in the corrosion product layer. The following standards were used for calibration: Al₂O₃ for
169 O, halite (NaCl) for Na, sylvite (KCl) for Cl, periclase (MgO) for Mg, wollastonite (CaSiO₃) for Ca and Si,
170 pyrite (FeS₂) for Fe and S and native copper for Cu. Non-quantitative EDX maps were collected at
171 higher current intensities (1 nA) to optimize the quality of the mapping vs the acquisition time.

172 2.5.2 X-ray diffraction (XRD)

173 In order to evaluate mineralogical evolution of the BCG with time, XRD analyses were conducted
174 directly on the BCG disk surfaces in a Bragg-Brentano geometry. When possible, corrosion products
175 were sampled from the steel coupons, ground in anoxic conditions and dispersed with 99% pure
176 ethanol on glass thin sections prior to their analyses. For both type of analyses, a D5000 Bruker AXS
177 diffractometer equipped with a Cu anticathode (λ = 1.5406 Å) operating at 40kV and 25 mA was
178 used. XRD scans were collected using an anti-air scatter screen in the range 2θ = 5° to 70° with a step
179 size of 0.01° and 3 sec/step. These scans were first treated with the DIFFRACPLUS EVA software
180 (Bruker) and the crystallographic open database (COD) was then used to attribute diffractions peaks
181 to the corresponding mineralogical phases.

182 2.5.3 µ-Raman spectroscopy

183 Corrosion products observed at both TU37b steel/BCG interface (experiment 1-R) and immersed
184 TU37b steel coupon surface (experiment 2-C) were characterized by confocal µ-Raman spectroscopy.
185 For the latter analyses, the steel coupons were transferred from the glove box to an in-house anoxic
186 cell avoiding further oxidation of the sample during analyses. Raman analyses were conducted with a
187 Jobin-Yvon Labram© HR800 Evolution spectrometer equipped with a 1200 lines/mm grating, leading
188 to a spectral resolution of ≈ 1-2 cm⁻¹. The Raman excitation was applied using a solid-state 532 nm
189 wavelength laser filtered at 10% of the maximal power (100 mW). An output energy of 1.5 mW was
190 measured with a power meter at the end of the microscope. The laser beam was focused using a
191 100× Olympus objective for experiment 1-R and a 50x long working distance Olympus objective for
192 experiment 2-C, leading to theoretical spots of approximately 0.90 µm and 2 µm, respectively, and
output energies of 2.4 x10⁹ Watt/m² and 4.8 x 10⁸ Watt/m², respectively. The acquisition time on a given sample was typically 20 s with 3 accumulations per measurement channel. The spectra were collected in the range 100 to 1400 cm⁻¹. Calibration was performed on a silicon standard using the 520.5 cm⁻¹ band prior to measurements. The spectra were acquired using the Labspec 6© software package and, when possible, the RRUFF database [25] was used for phase identifications. As corrosion product layers were often heterogeneous, a minimum of 20 Raman spectra were collected on each sample to ensure a good representativeness of the corrosion product layers.

### 2.5.4 Micro X-ray tomography (μ-XCT)

μ-XCT scans were obtained with an Easy Tom XL scanner (RX Solutions, France) equipped with a tungsten X-ray source and a 1516x1920 pixels CCD camera. Parallelepiped samples of BCG coming from experiment 2-C were scanned with 100 kV and 100 μA source settings and without filter. The scanned samples were enclosed in a 5x4x3 mm³ volume and a total of 1440 images were acquired per sample. Reconstruction was handled by the RX Solutions Xact software, producing 3D images with a 3.8 μm voxel size. The voxel values can then be compared between samples because the geometric, source, detector and reconstruction parameters remained unchanged. Post-processing of the μ-XCT scans was performed using Dragonfly 2020.2, Object Research System (ORS) Inc, Montreal, Canada, 2020 (www.theobjectscom/dragonfly). For each sample, a cylindrical or polygonal region of 12 ± 3 mm³ was selected. Thresholds were manually defined to segment the different component of the BCG and to allow the computation of volume rates.

## 2.6 Weight loss and optical measurements of the corrosion rates

In order to assess the evolution of the corrosion rates with time, weight loss measurements were performed on steel coupons periodically retrieved from experiment 2-C, following the ASTM recommendation [26]. Briefly, steel coupons were taken-off from the reactors in an anoxic glove-box and immersed in 99% pure de-aerated (under vacuum) ethanol for few seconds to stop corrosion. Then the coupons were dried, transferred outside the glove box, and immersed in the following successive baths: 15% HCL – 5% Hexamine solution, 5% Na₂CO₃ and Isopropanol (99%) with intermediate rinsing in deionized water between each bath. After this process, the coupons were weighted with an accuracy of 0.1 mg. This operation was repeated until the weight stabilization of the coupons. The corrosion rates were calculated from the following equation:

$$ \text{Corrosion rate (μm} \cdot \text{y}^{-1}) = \frac{3650 \times \text{weight loss (mg)}}{\text{Density} \left(\frac{g}{\text{cm}^3}\right) \times \text{Area} (\text{cm}^2) \times \text{Time (days)}} $$

with a density of 7.85 (g/cm³) for steel. In order to assess uncertainty on these measurements, a replicate weight loss measurement was performed for the sample retrieved at 92 days.
Corrosion damage was also assessed by optical measurements on cross-sectioned samples coming from experiment 1-R using a metallographic microscope (Kern, Germany). For these analyses, micrographs were collected along the whole length of the steel rods for each type of exposed surface (3µm mirror polished, 180 grid, external mill scale and internal mill scale). Measurements of the corrosion thickness were realized every 50 µm using ImageJ software, leading to more than 200 measurements per surface type. For the mill-scale exposed surface, the initial mean mill-scale thickness was previously determined by measuring it in a pristine sample with a similar protocol. This thickness was considered as the zero value and subtracted to the measured value of the corroded samples. Uncertainty on the measurements can be established from the measurements on two different cross sections for one sample immersed in PWS (porewater solution) for 91 days.

## 2.7 Geochemical modelling and thermodynamic databases

Pourbaix diagrams and thermodynamic equilibrium reaction paths were calculated with the geochemical module CHESS of the reactive transport code HYTEC [27] to support the discussion of the experimental results. The truncated-Davies activity correction model was considered in the reaction-path modelling. Similar results were obtained while considering either the thermodynamic databases Thermoddem, version 10 [28], or ThermoChimie, version 10 [29]. The Thermoddem database was, however, selected since the thermodynamic constants are accurately defined at 80°C, whereas ThermoChimie is assumed to be operational below 80°C only. Cronstedtite made an exception since the pure-iron composition, Fe₄SiO₅(OH)₄, of ThermoChimie was found to be simpler and more relevant for the present calculations than the mixed aluminium-iron, Fe₃SiAlO₅(OH)₄, composition of Thermoddem. Table A-1 gives the formation constants of the solid phases (mostly iron corrosion products) considered in the modelling.

# 3 Results

## 3.1 Evolution of the BCG mineralogy and porewater chemistry

After 32 days of curing time, XRD analyses of BCG disks indicated the presence of AFt (aluminate-ferrite trisulfate/ettringite), AFm (aluminate-ferrite monosulfate) and gypsum that mostly came from hydration of the cementitious phases. Hydrotalcite, cristobalite, calcite, montmorillonite and quartz were also identified and were related to the constituents of Table 1. The presence of amorphous SiO₂ from the silica fume and C-(A)-S-H was also pointed out by the presence of bulge between 10 and 35° 2θ and between 28.5 and 30.5 °2θ, respectively (Fig. A-2). The lack of portlandite (CH) was a characteristic of such low-pH cement where silica fume enhanced the formation of C-S-H instead of CH. The main change in time of the bulk mineralogy is the full depletion of ettringite and gypsum (or anhydrite) precipitation (Fig. A-2).
Mineralogical and microstructural evolution of the BCG in contact with PWS at 80°C was also assessed by µ-XCT. Indeed an increasing interest was devoted to this technic for cementitious material characterization [30] as it allows to differentiate cement phases in function of their relative density, no matter their structural state (i.e. crystalline or amorphous), while preserving the microstructural integrity of the sample. Fig. 4 shows µ-XCT slice sections of densified silica fume particles (DSFp) together with BCG sections immersed for different period of time in PWS at 80°C.
Four different components can be differentiated in the pristine sample (i.e. cured for 32 days). They are denoted as Porosity, Blast Furnace particles (BFp), Densified Silica Fume particles (DSFp) and Matrix. The BFp component is constituted of bright dense particles measuring several tens of micrometers that mostly come from the blast slag furnace component added in the CEM III-A. Some of these bright particles might also come from the bentonite (i.e., pyroxenes, feldspar, zeolite, quartz, cristobalite, pyrite) or from the crystalline fraction of the silica fume component (i.e., cristobalite, quartz and moissanite). The volume fraction of this component was determined to be approximately 2 % in the pristine material and it dropped to 1% for all the other samples immersed in PWS at 80°C. The DSFp consist of large aggregates of spherical nanoparticles of SiO₂ [31]. These particles are cracked and totally dissolved in some areas in the pristine sample (t = 0 day) which reflect their porous character and the availability of silicon in solution to form C-(A)-S-H from SF (Silica Fume). Also, these particles are darker than the unhydrated DSFp analysed in similar conditions and thus present a lower density. This latter observation suggests that hydration or partial dissolution of these particles occurred at 20°C. After 28 days of immersion in the PWS at 80°C, the mean grey level of these particles increased, corresponding to an increase of their mean density.
Moreover, a density zonation occurred in most of these particles, indicating a porosity and/or a chemical variation in the different layers of these particles. These observations can be attributed to the enrichment in the calcium content of the DSFp which is due to the calcium release from Ettringite degradation. This suggested formation of C-(A)-S-H inside these particles with a zonation coming from the dissolution of some layers prior to the immersion of the cementitious disks in the PWS at 80°C. Then, the grey level of the DSFp started to decrease with time, especially in the core of the particles, indicating a decrease in the core density of the DSFp. This result suggested a destabilization of C-(A)-S-H in the DSFp.

Fig. 5 shows the evolution of both pH and chemistry of PWS sampled from experiment 1-R. The initial pH at 80°C was 9.6 and only slight variations of pH occurred over 350 days. Within the first 14 days of contact with the BCG, the concentration of all the measured elements started to increase in the PWS, with a noticeable increase of SO₄²⁻ and Ca²⁺ due to ettringite destabilization followed by some decrease after 200 days due to anhydrite and gypsum precipitation. The concentration of Cl⁻, which
292
293

could have a detrimental effect on corrosion, increased from 66 to 183 mg/L, maybe due to the dissolution of Friedel's salts contained in the CEMIII/A component.

294

## 3.2 Corrosion of steel coupon immersed in PWS

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Raman spectroscopy was performed on the surface of the steel coupons coming from experiment 2-C after different periods of interaction with the PWS and the BCG disks at 80°C. Typical Raman spectra collected from the coupons exposed for 28, 91 and 300 days are given in Fig. 6. After 28 days of immersion, spectra characteristic of partially oxidized mackinawite ($Fe^{2+}_{1-3x}Fe^{3+}_{2x}S$) with bands located at 251, 307 and 348 cm$^{-1}$ and of $\alpha$-$S_8$ with bands at 167 cm$^{-1}$ and 230 cm$^{-1}$ [32,33] were observed (Fig. 6 a). The presence of calcite ($CaCO_3$) was also depicted with characteristic bands located at 280, 711 and 1086 cm$^{-1}$ (Fig. 6 b). Those phases were also detected in the sample immersed for 91 days (Fig. 6 c, d). Mackinawite is often encountered as a corrosion product of steel in contact with argillaceous medium and can have either biotic [15,16,34,35] or abiotic [5-7,18,35] origin. Mackinawite can oxidize without the presence of $O_2$ in acidic conditions and in the present alkaline conditions, it is likely that oxidizing species such as residual $O_2$ or polysulfides were involved in the oxidation of mackinawite [33].

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316

On the last sample, immersed for 300 days, vibrational modes arising from partially oxidized mackinawite and $\alpha$-$S_8$ were observed as well as two other bands located at 339 and 376 cm$^{-1}$. These narrow bands had been attributed to mackinawite by Sherar et al. [36] and to unknown iron sulphides by Robineau et al. [7]. However, the first authors were not able to attribute all bands to mackinawite and the second mentioned that this Raman signal was very similar to the signature of pyrite or marcasite which are both polymorphs of $FeS_2$. It is well known that mackinawite is a metastable iron sulphide that can destabilize to form pyrite, with intermediate greigite formation, at temperature as low as 60°Cand for a wide range of pH from 3 to 12. Such transformations are possible in the presence of oxidizing species or when the surface of the FeS precursor is pre-oxidized [37-39].

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324

Fig. 7 shows the XRD analyses of the corrosion products retrieved at the surface of the steel coupons. After 28 days, magnetite was the only corrosion product revealed by XRD and poor hkl reflection corresponding to calcite was observed. The large bulge located between 12 and 38°2θ corresponds to the glass thin section used as sample holder. After 300 days, magnetite was still the major corrosion product observed by XRD. However new hkl reflections appeared with d values of 7.1 and 3.55 Å, which can be attributed to Fe-silicates such as berthierine, odinite, cronstedtite or greenalite. However, the poor (hkl) reflections did not allow discriminating between all these layered T-O phyllosilicates of the kaolinite-serpentine group. These Fe-silicates are commonly found at clay/iron
325 interfaces at high temperature [10,40–42]. Fe-sulphides were not observed by XRD, although μ-
326 Raman spectroscopy indicated that they were present in large amount at the surface of the coupons.
327 It suggests that those phases were either poorly crystallized, or only present in low amount in the
328 outermost corrosion products (CPs) layer.

329 The surface of one 28-day coupon from experiment 2-C was scanned by SEM (Fig. 8). EDX analyses of
330 some large dark tubercles showed main contents in Fe, S, Si and O and secondary contents in Na, Al
331 and Ca (Fig. 8 a,b). This is in agreement with the formation of partially oxidized mackinawite
332 observed by μ-RS and magnetite observed by XRD. Moreover, octahedral crystals containing Ca and
333 O were also present in large amount at the surface of the coupon (Fig. 8 c,d). These grains were
334 identified as calcite (already observed by XRD and μ-RS), and EDX analyses indicated a slight
335 incorporation of Fe (less than 1 at.%) into this phase. Eventually, platy crystals containing mostly Fe,
336 Si and O were locally observed forming circular cluster inside bigger dark spots of CPs (Fig. 8 e,f).
337 These crystals also incorporated Na, Mg, Al, Ca, and S in lower amount. Such Fe-silicates formation
338 was also identified by XRD but only after 300 days of immersion in PWS. The quantity of Fe-silicate
339 produced in 28 days was indeed probably too low to be detected by XRD. EDX analysis of this phase
340 showed that it could not be odinite or berthierine due to its low Al content. The Fe/Si ratio of this
341 phase was more representative of cronstedtite, but the rosette-like morphology of the crystals was
342 more characteristic of greenalite [41].

343 3.3 Corrosion of steel in contact with BCG

344 3.3.1 Macroscopic assessment of corrosion damage

345 Fig. 9 shows cross sectional views of the steel rod in direct contact with BCG after different periods of
346 interaction with the PWS at 80°C (experiment 1-R). As for experiment 2-C, macroscopic observations
347 showed that localized corrosion occurred. Highly damaged areas were preferentially localized on
348 both mill-scale exposed surface and 180-grit size exposed surfaces (i.e. surfaces presenting defects)
349 of the rod. Furthermore, one can note that these surfaces were generally much more corroded than
350 the rod surfaces only exposed to PWS, i.e., simulating the spatial heterogeneity (data not shown). For
351 instance, no corrosion was macroscopically observed for samples immersed for 147 and 224 days.
352 For this reason, most of the characterizations were conducted in the area in direct contact with BCG.
353 Moreover, the cross-sectional view of the steel in direct contact with BCG showed an important
354 transformation of the BCG medium, up to 5 mm into the bulk after 350 days of exposure.

355 3.3.2 Nature and corrosion sequence of the dense product layer (DPL)

356 SEM-EDX and μ-Raman analyses were conducted at the interface between the BCG and the steel on
357 cross-sectioned samples from experiment 1-R. Fig. 10 reveals typical SEM-EDX mapping of corroded
areas obtained on the different rod sample surfaces for different periods of interaction with BCG.
After 28 days, EDX mappings performed on the 3 µm rod surface revealed the formation of an
uniform corrosion layer, measuring between 10 and 20 µm, mainly composed of Fe, S and O.
However, areas presenting hints of localized corrosion were macroscopically observed on the mill
scale surfaces (Fig. 9 and Fig. A-3). In this experiment, localized corrosion was clearly depicted after
91 days of interaction, particularly on the mill-scale layer exposed surfaces (Fig. A-3) with the
presence of pits already observed in experiment 2-C. Such highly damaged zones were more
abundant at the longest periods of interaction but also more extended laterally, suggesting their
coalescence with time. These dense corrosion product layers (DPL) are very similar in their structure
and chemistry for the samples that interacted for more than 147 days with saturated BCG. Additional
EDX mapping and EDX analyses (Fig.11 and Table 3) were performed in parallel of µ-Raman analyses
(Fig.12) on mill scale surface of a rod sample directly exposed to BCG for 147 days. Results showed
the formation of a dense corrosion product layer (DPL) composed of several layers that are
particularly well visible in this sample (Fig.11). The innermost layer was constituted of Fe, O plus
small amounts of S (Table 3, analysis 1). Raman analyses were conducted on this layer. A low laser
beam energy was necessary to avoid the degradation of its constitutive phases into hematite (4.8 x
$10^7$ Watt/m$^2$), compensated with a reasonable longer acquisition time. A very low signal to noise
ratio was obtained despite this compensation. Fig. 12 a displays the Raman spectrum representative
of this layer. Only a broad band near 680 cm$^{-1}$ was evidenced, indicating that this phase was certainly
very poorly crystallized. Since this phase is mainly composed of Fe and O, it can correspond to
several types of iron oxides or oxyhydroxides. Typically, this spectrum can be attributed to poorly
crystallized magnetite, maghemite or ferrihydrite [43], but the latter phase is unstable and transform
into goethite under alkaline conditions [44]. The Fe/O ratio of about 0.42 rather suggests that this
layer is mainly constituted of oxyhydroxides. Moreover, in other location of the innermost layer,
Raman analyses revealed the presence of goethite (Fig. A-4), in agreement with such assumption.
The presence of low amount of S suggests that oxyhydroxides are mixed with FeS in this layer. Such
Raman spectra were also obtained for the Fe, O, S containing layer observed in the 28-day sample
(data not shown).

The second layer was large and alternated dense products, composed of Fe and O with an Fe/O
atomic ratio of approximately 0.7, and more porous mixtures of Fe, O containing CPs with scarce Fe-S
compounds. The morphology of some Fe, O micrometric crystals displayed habits from the cubic
symmetry (Fig. 11 b). The Raman analyses showed the main occurrence of magnetite, with bands
located at 322, 501 and 665 cm$^{-1}$. The inclusion of residual particles from the innermost layer in this
magnetite layer suggests that it has been partially replaced by magnetite. The outermost layer was a
thin fringe composed of Fe, Si, O and S. The cutting of the magnetite layer by this Fe, Si, O, S fringe
suggested that it was replacing magnetite. In some part of this fringe, Raman spectra typical of
disordered hematite were obtained (not shown). As mackinawite was observed in EXP 2-C, it is
probable that mackinawite was present in this layer and destabilized under Raman beam. This
degradation phenomenon was observed by Genchev and Erbe (2016), on mackinawite layers that
were analysed ex situ without being protected by a thin water layer [45]. Octahedral shaped crystals
containing less S were also observed in the vicinity of this layer (Fig. 11 b and Table 3 - analyse 5). The
μ-RS with a characteristic band at 330 cm⁻¹ did not allow to identify this phase (Fig. 12 c). Eventually,
Fe-sulphides were occasionally observed between the outermost layer and the BCG. However, these
sulphides were not observed in the last sample after 350 days. EDX showed that this phase was
mostly composed of Fe and S, with possible incorporation of Si and O. Raman signal of this phase
showed bands located at 315, 322, 368 and 376 cm⁻¹ (Fig. 12 d), similarly to the 300-days sample in
experiment 2-C. In some areas, the Fe-sulphides also contained Cu up to 1.3 at.% (Table 3). Cu-
sulphides precipitating at the outermost layer of the DPL were also discussed in a recent study on
steel corrosion in an argillaceous medium at 90°C [18].

In addition to the DPL sequence, a magnetite layer was also observed in the vicinity of the BCG (Fig. 10). This layer could be either present instead of, or in addition to, the magnetite layer adhering to the first layer. In a similar manner, this outer layer of magnetite was partially replaced by a Fe, Si, S, O phase or mixture of phases, and unknown iron sulphides precipitate in the interstitial voids between this fringe and the BCG.

### 3.3.3 Characterization of the transformed medium (TM).

Fig. 13 shows SEM-EDX mapping of large halos spreading in the BCG. They are located in the vicinity of highly corroded areas on the mill-scale side for 147 and 224 days. Such halos of transformed medium (TM) were frequently observed in the 3 last samples (147, 224 and 350 days) from experiment 1-R, and more particularly where mill-scale was exposed to the BCG and/or at the edges of the steel rod. These halos were denser and less porous than the rest of the BCG. They were constituted of three envelopes. The innermost envelop of the 147-days sample was mainly enriched in Fe and O, the second one in S, Fe, Si and O, and the outermost one in Fe, Si and O. This sequence was similar to the one observed in the DPL in the vicinity of the steel, except that Fe-sulphides were also observed in the first layer of the DPL. Such a sequence would demonstrate that some Fe did not precipitate at the surface of the steel but migrated in the highly porous BCG. The subsequent precipitation of Fe-oxides occurred inside the TM, probably as magnetite and/or FeO(OH) due to
425
426

local variations in pH. Eventually, the diffusion of HS⁻ and of Si coming from the BCG promoted the transformation of the Fe-oxides into Fe-sulphides and Fe-silicates.

427
428
429
430
431
432
433
434
435
436
437
438
439

Some halos only presented a dense outer envelope composed of Fe, S, Si and O, no matter the duration (e.g., 224-day sample in Fig. 12). This might be linked to the amount of Fe release in the solution and to the availability of HS⁻ and Si in the surrounding BCG, which was quite heterogeneous. The presence of DSFp relics were also visible inside the dense halos. However, these relic grains were Fe-enriched compared to the DSFp localized in the surrounding environment. Moreover, TM was sometimes characterized by denser areas enriched in Si, O, and to a lesser extent Ca and Fe. C-S-H with low Ca/Si ratio could have precipitated in the vicinity of the steel in the TM and up-taken Fe. The EDX mapping confirms the evolution of DSFp with time as observed by µ-XCT. The DSFp surrounding the TM were enriched in Ca during the first 28 days (Fig. 11), followed by Ca depletion in the cores of DSFp (Fig. 10 and Fig. 13). Eventually, it is worth mentioning that Ca-O, and Ca-O-S containing phases were also present in the TM (Fig. 10, 224-days sample). They can be attributed to portlandite and gypsum, respectively, phases also identified by XRD at the surface of the BCG disk in experiment 2-C (not shown).

440

## 3.4 Steel corrosion rates

441
442
443
444
445
446
447
448
449
450
451
452
453
454
455
456
457

All the samples coming from both experiments 1-R and 2-C showed evidences of localized corrosion processes associated to the presence of numerous pits whose depth could reach 160 µm (or 122.5 µm considering the maximum thickness of the mill scale layer) in the 147 days sample (Fig. 14 a). To assess the extension of the localized corrosion phenomenon, the dispersion on the optical corrosion thickness measurement is reported in Fig. 14 b. The histogram shows that most of the measurements (>60 %) are in the range [0-10] µm for all samples, which principally corresponds to areas unaffected by corrosion, confirming the localized corrosion phenomenon. However, microscope observations clearly indicated an increase of the number of damaged zones with time, rather than an increase of the pit depth (Fig. 14 a and Fig. A-3). Indeed, a large increase of the number of pits with depth limited to 40 µm is clearly depicted after 350 days of interaction, whereas the number of measurements of pits >40 µm is not increasing anymore (Fig. 14 b). This observation suggests that the coalescence of these pits would eventually lead to a generalized corrosion process.

For the sake of comparison, the mean corrosion rates obtained by optical measurements for EXP 1-R and by weight loss measurements for EXP 2-C are presented in Fig. 15. A first increase in the corrosion rate is depicted in the first 224 days for EXP 1-R and in the first 93 days for EXP 2-C up to values reaching 10.8 µm/y and 8 µm/y, respectively. Then, the mean corrosion rates seem to stabilize, or even to decrease, for both experiments.
458 **4 Discussion**

459 **4.1 Initiation of steel corrosion and effect of the mill-scale**

460 In both experiments 1-R and 2-C, localized corrosion started to occur on surfaces presenting defects,
461 namely, at the edges of the coupons/rod, on the rough surfaces (ground with 180 grit) and on the
462 mill-scale layer exposed surfaces. This can be explained by the formation of corrosion cells induced
463 by the presence of an electron conductive layer (here magnetite), acting as the cathodic area, and
464 the ingress of HS⁻ from the BCG that reacts with iron to form FeS in the anodic area (step 1 Fig. 16).
465 This process has been observed in various environment (oxic, anoxic) [5,6] and well described in the
466 study of Robineau et al. (2020) [6]. On the mill-scale exposed surfaces, the magnetite layer acting as
467 a cathode is already present at the beginning of the experiment. In this case, either (i) structurally
468 defective area corresponding to zone uncovered by magnetite or (ii) chemically depassivated area
469 can act as anode. Typically, HS⁻ might also react with magnetite to form mackinawite even at
470 moderate alkaline pH (i.e. 9-10) provided that the concentration of dissolved sulphides is high
471 enough (Fig. A-5).

472
473 Localized corrosion process also occurred in other types of surfaces (i.e. 180 grit ground surfaces).
474 Passivation of these surfaces could have occurred prior to the HS⁻ ingress from the grout to the steel
475 surface. Indeed, nanometric layers of magnetite can form either by the Schikorr reaction or by a solid
476 state reaction implying Fe diffusion from the steel and O diffusion toward the steel bulk, through O
477 vacancies diffusion toward the surface [21]. For this type of surface, local depassivation of the
478 magnetite layer and formation of galvanic coupling would preferentially occur in the defect-
479 presenting areas [17]. Such nanometric layers of iron oxide were not observed, nor at the surface of
480 the rods, neither at the surface of the coupons. However, a thin layer of magnetite oxide was
481 observed on preliminary short-term experiments conducted in similar conditions, but without BCG
482 (not shown). This layer was possibly destabilized by the presence of HS⁻ coming from the cement.

483 It is also well known that Cl⁻ ions have a prominent role in the depassivation of steel, depending on
484 parameters such as pH, O₂ availability and temperature. Particularly, depassivation is associated to a
485 [Cl⁻]/[OH⁻] threshold value. In experiment 1-R, the Cl⁻ concentration increased to a value of 5.10⁻³
486 mol/L when the BCG interacted with the PWS at 80°C. In parallel, the increase of the temperature
487 lead to a decrease of pH and a concentration in OH⁻ of 4.10⁻⁵ mol/L (corresponding to a [Cl⁻]/[OH⁻]
488 ratio of 125). Compared to the literature data, this high ratio can be responsible for the depassivation
489 of steel [46].
The presence of residual O₂ in the BCG might also be responsible for the initiation of localized corrosion. Indeed, in some areas, the first layer of corrosion products was mainly composed of iron hydroxides. At the beginning of the experiment, residual O₂ coming from the cement could have diffused through this material and been progressively consumed by the corrosion process to form Fe-hydroxides. The ingress of O₂ to the surface of the steel was transport-limited. Sherar et al. [47,48] showed that localized corrosion can occur at trace level of O₂, with the formation of locally corroded areas induced by the formation of a corrosion cell with magnetite layer acting as cathodic zone.

Eventually, it was shown that localized corrosion could be triggered by the presence of porous and heterogeneous mill scale layer composed of both magnetite (Fe₃O₄) and hematite/maghemite (Fe₂O₃). For instance, Qin et al. (2004) demonstrated that a decoupling of anodic and cathodic reactions could occur with iron oxidation and Fe₂O₃ reductive dissolution in anodic and cathodic areas, respectively [49]. In the present experiments, it is possible that Fe₂O₃ was present in some areas (not detected by µ-Raman) and led to the initiation of localized corrosion, even prior to the ingress of HS⁻.

## 4.2 Transformation of magnetite to iron silicates/iron sulphides

A large amount of magnetite was observed in both experiments. In experiment 1-R, the innermost layer was principally composed of Fe and O with a Fe/O ratio of approximately 0.42. The combination of this low Fe/O ratio together with the presence of goethite revealed by Raman in some area of the innermost layer suggested that this layer was principally composed of iron oxyhydroxide. Indeed, higher potential due to the presence of residual oxygen might have triggered the formation of goethite (Fig. 17). This phase was progressively replaced by magnetite in the DPL. As proposed by Usman et al. [50], this reaction can occur with an input of Fe²⁺ and OH⁻ which are produced in the anodic part and cathodic part of the corrosion cell, respectively. The reaction writes as:

$$2\text{Fe}^{3+}\text{OOH} + \text{Fe}^{2+} + 2\text{OH}^{-} \rightarrow \text{Fe}^{2+}\text{Fe}^{3+}_{2}\text{O}_{4} + 2\text{H}_{2}\text{O} \quad (1)$$

Its partial recrystallisation might have occured in the innermost layer, explaining the poor crystallinity of this phase as observed by Raman.

Direct precipitation of a magnetite fringe was also observed in the vicinity of the BCG in some samples. In that case, magnetite was produced after aqueous release of Fe followed by the Schikorr reaction:

$$\text{Fe} + 2\text{H}_{2}\text{O} \rightarrow \text{Fe(OH)}_{2} + \text{H}_{2} \quad (2)$$
$$3\text{Fe(OH)}_2 \rightarrow \text{Fe}_3\text{O}_4 + 2\text{H}_2\text{O} + \text{H}_2 \quad (3)$$

523

Magnetite precipitation started after the consumption of a first source of HS⁻ (step 2 in Fig. 16), in agreement with the Pourbaix diagram without HS⁻ (Fig. 17 a). When magnetite formed in anodic zone, OH⁻ ions were produced in the cathodic ones, enhancing the dissolution of the hydrated silica fume in the vicinity of the corroded area. The presence of a highly soluble silica source promoted the transformation of magnetite into iron silicates (step 3 of Fig. 16), as confirmed by the Pourbaix diagram of Fig. 17 c. Fe-silicates were observed in the DPL in both experiments. Furthermore, the experiment 1-R clearly evidenced that Fe-sulphides, combined with Fe-silicates, were replacing magnetite. The formation of a second Fe-sulphides containing layer, might be due to (i) the presence of a second source of sulphides in the BCG or (ii) to a change in the mechanism of sulphide release. Indeed, it is probable that the sulphides came from CEM III/A and bentonite that contained oldhamite (CaS) and pyrite (FeS₂), respectively. In blast furnace slag, CaS is present either in the form of free dentritic crystals or in inclusion in mellilite [51,52]. Thus, free crystals of CaS were readily available as a first source of sulphides in our experiment and then a second source would be available after the partial dissolution of mellilite crystals during the hydration of the BCG. Indeed, parts of the BFp reacted during the experiment as discussed in section 3.1, inducing potential sulphide release from the less reactive crystalline fraction of the BFp.

540

Iron sulphides were also observed in the outermost region of the DPL, these unidentified sulphides would have formed after a longer time of reaction. They might come from the dissolution of pyrite that is more stable than CaS. For instance, it is possible that H₂ release from corrosion triggered pyrite reduction via the following reaction:

$$\text{FeS}_2 + \text{H}_2 \rightarrow \text{Fe}^{2+} + 2\text{HS}^- \quad (4)$$

Such pyrite reduction was observed by Schlegel et al, who studied the corrosion of steel at contact with COx at 90°C and in anoxic conditions [18]. Regarding the formation of Fe-silicates, XRD and SEM-EDX analyses performed on steel coupons from experiment 2-C indicated that they belonged to the kaolinite-serpentine group. Based on their chemical formulas, as well as their structural and morphological features, they could correspond to greenalite and/or cronstedtite. Such Fe-silicates are often found associated to magnetite, iron sulphides and iron carbonates in primitive chondrites [53] or at the iron/clay [19,41] or iron/glass [54] interfaces. Prerequisites for the formation of cronstedtite are the presence of metallic iron, neutral or alkaline pH and a temperature between 50 and 120 °C [41]. It has recently been shown that the nature of serpentine-like Fe-silicates strongly
relied upon the precursors involved in their formation [55], with berthierine
((Fe²⁺,Fe³⁺,Al)₃(Si,Al)₂O₅(OH)₄) most likely formed from iron and kaolinite reaction, and cronstedtite
formed from iron and quartz reaction. Hydrated silica fume was the major source of silicon in the
present study, which explains the low Al content of the Fe-phyllosilicates.

Transformation of magnetite into Fe-silicates was also evaluated in the TM. The periodic post-mortem analyses of the TM evolution showed a periodic migration of the mineralogical front between Fe-oxides, Fe-sulphides and potentially Fe-silicates.

Thermodynamic equilibrium calculations were performed in order to simulate the observed
transformation of magnetite into iron silicates and iron sulphides in the DPL. The starting point
assumed a full hydration of the bentonite/cement-grout components and a porosity of 0.75 (i.e. 75 %
of water and 25% of solid phases in volume). Such a large porosity is due to the very high
Water/Cement ratio and is in agreement with water porosity measurements performed on this
material (not shown)". Furthermore, the aluminous, sulphated and carbonated phases were not
considered. That is to say that the full hydrated grout was initially composed of C-S-H with a Ca/Si of
0.8 (0.26 kg/L of porous media PM) and some remaining silica fume (0.11 kg/L of PM) modelled as
cristobalite. The initial sulphide content was fixed to 2 wt.% (brought by CEM III/A and bentonite)
and modelled as pyrite (0.019kg/L of PM). Temperature was set to 80°C and the initial pH to 9.5. The
effect of steel corrosion under anoxic conditions was tackled through the diffusion of an
alkaline/hydroxide Fe(II) plume modelled as the progressive addition (i.e. a titration-like model) of
Fe(OH)₂ moles that were free to react with the hydrated grout assumed to be located very close to
the steel interface.

Fig. 18 a shows that, in Stage I, pyrite is progressively transformed into mackinawite in parallel
("simultaneously") to the neoformation of cronstedtite (greenalite was not included in the model).

However, the amount of iron silicate is much higher than the amount iron sulphide and cronstedtite
keep precipitating in Stage II. Magnetite starts to precipitate (Stage III) only once the silica source
(cristobalite) is fully exhausted in the thermodynamic calculation. It is worth mentioning that the pH
remains all along the mineralogical sequence in a narrow range close to the values measured
experimentally (Fig. 18 b)). The redox potential Eh shifts towards very reducing conditions at the
boundary between elemental iron and mackinawite/cronstedtite equilibrium (Fig. 17 d, Fig. 18 b).

The second set of graphs (Fig. 18 c, d) splits the global reaction into a subset of increased chemical
complexity. Stage I-a considered only the oxyhydroxide system, then stage I-b broadened it to the
oxyhydroxide-sulphur system, and stages II and III included the full oxyhydroxide-sulphur-silicate
system. The calculation mainly indicates that cronstedtite can form at the expense of magnetite,
provided a Fe(OH)₂ source is supplied by steel corrosion. This phase would form until the whole
consumption of the Si source, which would be most probably hydrated silica fume in these experiments (modelled as cristobalite in the calculations). These results are in agreement with the observed corrosion product sequence in the DPL and suggest that the magnetite layer would not be totally destabilized by Fe-silicates from the drop in the corrosion rate (step 4 in Fig. 16). Note that, the second ingress of sulphides is not considered in this titration like model for the sake of simplicity. However, this model showed that Fe-sulphides and Fe-silicates can form simultaneously as observed experimentally.

## 4.3 Effect of spatial heterogeneity

In experiment 1-R, the surfaces exposed to PWS were less corroded than the surfaces directly exposed to BCG. This result is consistent with a diffusion-controlled process. Indeed HS⁻ ions released from the BCG were able to rapidly react at the surface of the steel prior to diffuse in the solution (electrolyte). Moreover, the presence of a restricted volume of water between the steel and the BCG locally affect the pH and the chemistry of PWS. This local change might have contributed in the extension of corrosion. However, it seems that the distance between the steel and the BCG has no effect on the observed corrosion mechanism. Indeed, similar corrosion products were observed in experiment 2-C in which steel was separated from the BCG by a large volume of solution. Nevertheless, a significant higher mean corrosion rate was initially observed in experiment 2-C despite the absence of mill-scale or rough surface on the coupons. This large difference can be attributed to the volume ratio between the PWS and the BCG. Equal Vol~pws~/Vol~cement~ (about 2.85) were used in both experiments, but the S~steel~/Vol~pws~ ratio was twice higher in experiment 1-C. Therefore, the steel exposed surface was also twice higher in experiment 2-C for the same quantity of released HS⁻.

## 4.4 Fe incorporation/sorption in C-S-H

The formation of Fe-silicates occurred in the DPL in experiments 1-R and 2-C. However, it is not clear whether the Fe-Si-O enriched envelope present in the TM was primarily composed of Fe-silicates or not. This layer was indeed enriched in Ca in some zones. It has recently been shown that Fe³⁺ can substitute Ca²⁺ in the interlayer of C-S-H of high Ca/Si ratio (1.2 to 1.5) or can form a separate secondary phase (or clusters) composed of Fe-Ca and Si at the surface of C-S-H of low Ca/Si ratio (0.8). A high Kd value of 700 (± 150) m³/kg was reported for Fe³⁺ whatever the Ca/Si ratio [56]. Furthermore, it has also been demonstrated that Fe²⁺ can substitute Ca²⁺ in the interlayer of C-S-H. The incorporated amount of Fe²⁺ does not depend on the Ca/Si ratio nor the pH, but it is very limited compared to Fe³⁺, with Kd values about 0.1 m³/kg [57]. In the present study, the preferential dissolution of silica fume close to the cathodic zone combined with the Ca release by ettringite
622 dissolution should have promoted the formation of low Ca/Si C-S-H. Furthermore, the aqueous
623 complex Fe(OH)₂ was the main species under the anoxic conditions, although Fe²⁺ can oxidize very
624 rapidly at low alkaline pH in the presence of trace amount of O₂ [57]. Therefore, incorporation of Fe²⁺
625 (and also possibly Fe³⁺) in C-S-H was likely but further investigation would be necessary to validate
626 such a hypothesis.

627 5 Conclusion

628 This time-resolved study provides significant insights on the corrosion of low carbon steel in the
629 presence of low-pH BCG designed for nuclear waste disposal but has also general implications for
630 steel corrosion under slightly alkaline conditions. The presence of HS⁻ and of residual O₂ in the BCG
631 promoted the formation of a galvanic coupling between cathodic zones principally composed of
632 magnetite and anodic ones where iron sulphides and iron oxihydroxydes formed. The structurally
633 irregular mill scale layer principally composed of magnetite enhanced the galvanic coupling, leading
634 to large corrosion damages with corrosion pits measuring up to 160 µm in depth after 147 days of
635 exposure. The formation of large amounts of magnetite indicated that passivation of the steel did not
636 occur, even after the consumption of aggressive species (e.g., HS⁻) coming from the BCG. However,
637 the mean corrosion rate dropped after one year indicating that the dense corrosion product layer
638 eventually acted as a protective barrier and limited the extension of pit depth. To the best of our
639 knowledge, the further transformation of magnetite into Fe-silicates (cronstedtite or greenalite) and
640 Fe-sulphides was evidenced for the first time in such a low-pH cement medium. Fe-silicate was
641 promoted by the dissolution of hydrated silica fume particles from the BCG, dissolution itself induced
642 by the release of OH⁻ at the cathodic site of the corrosion cells. This study also highlights the fact that
643 due to the diffusion-controlled nature of the corrosion processes, spatial heterogeneity leading to
644 imperfect contact at BCG/steel interface would not induce additional degradation of the metallic
645 surface.

646

647 Acknowledgments.

648 Research funding by the ACED/EURAD H2020 program under grant agreement No 847593 are
649 gratefully acknowledged. We would like to express our sincere gratitude to Nicolas Bessaguet and
650 Veronique Baty (Subatech) for their help in the solution analyses, Erwan Le Menn (LPG) for help with
651 Raman measurements and Nicolas Stephant (IMN) for his assist in SEM-EDX measurements. We
652 sincerely acknowledge the three anonymous reviewers for their fruitful comments. This manuscript
653 is dedicated to J-C Goethals.
654 **Data availability**

655 The raw/processed data will be available upon request.

656

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837

838

## Figure captions

839

840

**Figure 1:** (a) Schematic cross section of the French multi-barrier concept of a disposal cell for high-level nuclear waste [55]; (b) cross section of the BACUCE *in situ* experiment.

841

842

843

844

845

**Figure 2:** (a) Microstructure parallel to the rolling direction and composition of the TU37b steel measured by Optical emission spectrometry with a SPECTRO MAXx Ametek®; (b) microstructure of the external mill-scale layer. The microstructure of the steel was revealed after polishing the steel surface down to 0.25 µm with silicon colloids following chemical etching with a 2% HNO₃ Nital solution.

846

847

**Figure 3:** (a) Set-up of the experiment 1-R with (b) a schematic 3D-view of the grout-steel assembly put in the PTFE reactor; (c) set-up of the experiment 2-C.

848

849

850

**Figure 4:** µ-XCT slices sections of (a) the densified silica fume particles (DSFp) together with grout cement disk sections coming from experiment 2-C after (b) 0, (c) 28, (d) 93 and (e) 300 days of interaction with PWS at 80°C.

851

**Figure 5:** Change with time of (a) the pH and (b) the chemistry of the PWS in experiment 1-R.

852

853

854

**Figure 6:** Raman spectra of the corrosion products collected in anoxic cell, at the surface of the steel coupons from experiment 2-C (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).

855

856

**Figure 7:** XRD patterns of the corrosion products collected at the surface of the steel coupons from experiment 2-C after 28 and 300 days of interaction with PWS at 80°C.

857

858

859

**Figure 8:** SEM micrographs of the corrosion products observed at the surface of a coupon from experiment 2-C after 28 days of interaction with PWS at 80°C (a,c,e) along with EDX analyses of the corrosion products denoted by the red cross symbols (b,d,f).

860

861

862

863

**Figure 9:** Optical micrographs of the cross sectioned samples from experiment 1-R for (a) 28, (b) 93, (c) 147 and (d) 350 days of interaction with BCG at 80°C. Observed area on the cross sections are indicated by red rectangles (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
864

**Figure 10:** SEM-EDX chemical maps of the BCG/steel interface for progressing times at different location of the steel rod (3µm surface finish, 180 grit surface finish and external mill scale layer surfaces). The relative amounts of elements are given by a gray-scale value.

865

866

867

**Figure 11:** SEM pictures of a highly corroded area observed on a sample from experiment 1-R after 147 days of interaction with BCG at 80°C, with a focus on (a) the DPL and (b) the outermost region of the DPL. EDX and µ-Raman analyses marked by red spots are reported in Table 3.

868

869

870

**Figure 12:** Raman spectra corresponding to the area marked in Fig. 11.

871

**Figure 13:** SEM-EDX mapping of the transformed medium in samples from experiment 1-R after 147 and 224 days of interaction between BCG and steel at 80°C (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).

872

873

874

**Figure 14:** Optical measurements of the corroded layer for EXP 1-R with (a) Maximum corroded thickness obtained. The values correspond to the maximum reported thickness observed on the external mill scale layer at direct contact with the BCG. The "0 value" corresponds to the maximum thickness reported for the external mill-scale layer in a pristine steel sample. (b) Statistical histogram of the optical measurements of the corrosion thickness measured on the external mill-scale layer for samples coming from EXP 1-R. More than 250 measurements spaced 50µm apart were performed on each sample. The thickness of the mill-scale layer was not subtracted from the measurement.

875

876

877

878

879

880

881

**Figure 15:** Mean corrosion rates of carbon steel obtained from optical measurements of the corroded layer for experiment 1-R (black squares) and from weight loss measurements for experiment 2-C (red circles).

882

883

884

**Figure 16:** Schematic representation of the sequence of corrosion processes and BCG evolution deduced from the experiments 1-R and 2-C.

885

886

**Figure 17:** Pourbaix diagrams of iron for increasing chemical complexity calculated with CHESS at 80°C. (a) oxyhydroxide system, activity $Fe^{2+} = 10^{-5}$; (b) oxyhydroxide-sulfur system, activity $Fe^{2+} = 10^{-5}$, total S = $2 \times 10^{-4}$; (c) oxyhydroxide-silicate system, activity $Fe^{2+} = 10^{-5}$, total Si = $2 \times 10^{-3}$; (d) full system, activity $Fe^{2+} = 10^{-5}$, activity $HSiO_3^-$ = $2 \times 10^{-3}$, total S = $2 \times 10^{-4}$. The blue zone stands for the range of pH measured in the cell solution at 80°C.

887

888

889

890

891

**Figure 18:** Simplified geochemical modelling with CHESS at 80°C of the diffusion of an $Fe(OH)_2$ plume induced by steel corrosion inside the cement/grout material. (a, b) full oxyhydroxide-sulfur-silicate system from the beginning; (c, d) increased chemical complexity: oxyhydroxide system (stage I-a),

892

893
894 then oxyhydroxide-sulfur system (stage I-b), and finally the oxyhydroxide-sulfur-silicate system.

895 (stages II and III).

896 Table captions

897 **Table 1:** Composition of the bentonitic cement-based grout (BCG).

898 **Table 2:** Initial chemistry of the synthetic solution used in experiments 1-R and 2-C. Values are given in mg/L.

899 **Table 3:** EDX analyses localized on the red spots of Fig. 11. The values are given in atomic percentage and normalized to 100%.

902
<figure><img src="image_2.png" /></figure>

<figure><img src="image_3.png" /></figure>
<figure><img src="image_4.png" /></figure>

<figure><img src="image_5.png" /></figure>
a)

<figure><img src="image_6.png" /></figure>

b)

<figure><img src="image_7.png" /></figure>

c)

<figure><img src="image_8.png" /></figure>
<figure><img src="image_9.png" /></figure>

<figure><img src="image_10.png" /></figure>

<figure><img src="image_11.png" /></figure>

<figure><img src="image_12.png" /></figure>

<figure><img src="image_13.png" /></figure>
<figure><img src="image_14.png" /></figure>

<figure><img src="image_15.png" /></figure>
28 days

a)

<figure><img src="image_16.png" /></figure>

b)

<figure><img src="image_17.png" /></figure>

91 days

c)

<figure><img src="image_18.png" /></figure>

d)

<figure><img src="image_19.png" /></figure>

300 days

e)

<figure><img src="image_20.png" /></figure>
<figure><img src="image_21.png" /></figure>

Mg: Magnetite

Cc: Calcite
<figure><img src="image_22.png" /></figure>

a)

<figure><img src="image_23.png" /></figure>

c)

<figure><img src="image_24.png" /></figure>

e)

<figure><img src="image_25.png" /></figure>

b)

<table><tr><td>O</td><td>59.4</td></tr><tr><td>Na</td><td>0.7</td></tr><tr><td>Al</td><td>0.1</td></tr><tr><td>Si</td><td>4.3</td></tr><tr><td>S</td><td>7.6</td></tr><tr><td>Ca</td><td>0.1</td></tr><tr><td>Fe</td><td>27.7</td></tr></table>

<figure><img src="image_26.png" /></figure>

d)

<table><tr><td>O</td><td>50.3</td></tr><tr><td>Na</td><td>0.7</td></tr><tr><td>Mg</td><td>0.3</td></tr><tr><td>Al</td><td>0.2</td></tr><tr><td>Si</td><td>0.8</td></tr><tr><td>Si</td><td>0.2</td></tr><tr><td>Ca</td><td>46.7</td></tr><tr><td>Fe</td><td>0.8</td></tr></table>

<figure><img src="image_27.png" /></figure>

f)

<table><tr><td>O</td><td>56.1</td></tr><tr><td>Na</td><td>2.4</td></tr><tr><td>Mg</td><td>1.2</td></tr><tr><td>Al</td><td>0.5</td></tr><tr><td>Si</td><td>13.3</td></tr><tr><td>S</td><td>0.8</td></tr><tr><td>K</td><td>0.8</td></tr><tr><td>Ca</td><td>1.6</td></tr><tr><td>Fe</td><td>23.5</td></tr></table>
<figure><img src="image_28.png" /></figure>

Cement-grout

<figure><img src="image_29.png" /><figcaption>a)</figcaption></figure>

Cement-grout

<figure><img src="image_30.png" /><figcaption>c)</figcaption></figure>

steel

steel

<figure><img src="image_31.png" /></figure>

Cement-grout

<figure><img src="image_32.png" /><figcaption>d)</figcaption></figure>

steel

steel
<figure><img src="image_33.png" /></figure>

<figure><img src="image_34.png" /></figure>

<figure><img src="image_35.png" /></figure>

<figure><img src="image_36.png" /></figure>
<figure><img src="image_37.png" /></figure>

<figure><img src="image_38.png" /></figure>

<figure><img src="image_39.png" /></figure>

<figure><img src="image_40.png" /></figure>
<figure><img src="image_41.png" /><figcaption>224 days</figcaption></figure>

<figure><img src="image_42.png" /></figure>
<figure><img src="image_43.png" /></figure>

<figure><img src="image_44.png" /></figure>
<figure><img src="image_45.png" /></figure>
<figure><img src="image_46.png" /></figure>

* Destabilization of CaS coming from CEM III/A leads to the release of HS⁻ in porewater solution (PWS).

* Presence of mill-scale defective layer (magnetite)/Local depassivation of the protective layer (magnetite).

* Formation of a corrosion cell with OH⁻ produced in the cathodic area (magnetite layer) and FeS (mackinawite)/FeO(OH) produced in the anodic area (steel surface).

* Iron starts to diffuse in the low-pH bentonitic cement grout (BCG) and precipitate in the form of iron oxide, leading to partial pore clogging of the BCG.

* At 80°C, the densified silica fume particles (DSFp) are transformed into C-S-H (Ca come from AFt dissolution).

<figure><img src="image_47.png" /></figure>

* HS⁻ in the PWS are consumed by the corrosion, leading to the formation of magnetite. both at the steel surface and at the vicinity of the BCG.

* The formation of magnetite and the hydration of the unreacted cementitious phases (UC) are favoured by the release of OH⁻ in the cathodic area.

* Impoverishment of the Ca content of C-S-H formed in the DSFp.

<figure><img src="image_48.png" /></figure>

* Si and a second HS⁻ ingress from the BCG destabilize the magnetite layer present at the vicinity of the grout. Mixed iron silicates (probably cronstedtite or greenalite) and sulphides replace magnetite both in the dense product layer (DPL) and in the transformed medium (TM).

* Large source of silicon is available from the hydrated DSFp.

* Magnetite can still form at the vicinity of the steel in the area where localized corrosion was initiated due to the limited input of Si and HS⁻ in this area.

<figure><img src="image_49.png" /></figure>

* Destabilization of pyrite (FeS₂) induces a third release of sulphides (HS⁻ and or polysulphides) that diffuses from the BCG and lead to the formation of unknown iron sulphides between the steel and the BCG. Eventually, it is probable that the magnetite formation and the pore clogging in the transformed medium (TM) act as a diffusion barrier for species such as HS⁻. Deepening of the pre-formed pit is limited.
<figure><img src="image_50.png" /></figure>

<figure><img src="image_51.png" /></figure>

<figure><img src="image_52.png" /></figure>

<figure><img src="image_53.png" /></figure>
<figure><img src="image_54.png" /></figure>

pH

<figure><img src="image_55.png" /></figure>

Eh [V]

<figure><img src="image_56.png" /></figure>

pH

<figure><img src="image_57.png" /></figure>

Eh [V]
<table><thead><tr><td>Constituents</td><td>Weight (g)</td><td>Provider</td></tr></thead><tbody><tr><td>CEM III/A 52,5L CE PM-ES CP1 NF</td><td>600</td><td>Rombas</td></tr><tr><td>Silica Fume S95DM</td><td>600</td><td>Condensil</td></tr><tr><td>Bentonite BENTOSUND A100</td><td>160</td><td>Laviosa</td></tr><tr><td>Hydrotalcite (Mg<sub>6</sub>Al<sub>2</sub>(CO<sub>3</sub>)OH<sub>16.4</sub>H<sub>2</sub>O</td><td>40</td><td>/</td></tr><tr><td>Tap water</td><td>2400 (W/C = 1.7)</td><td>/</td></tr></tbody></table>
<table><thead><tr><th></th><th>Temp = 20°C</th><th>Temp = 80°C</th></tr></thead><tbody><tr><td>pH</td><td>11.1</td><td>9.6</td></tr><tr><td>Na<sup>+</sup></td><td>6.55×10<sup>-3</sup></td><td>7.00×10<sup>-3</sup></td></tr><tr><td>K<sup>+</sup></td><td>7.2×10<sup>-4</sup></td><td>7.7×10<sup>-4</sup></td></tr><tr><td>Mg<sup>2+</sup></td><td>9×10<sup>-5</sup></td><td>6×10<sup>-5</sup></td></tr><tr><td>Ca<sup>2+</sup></td><td>1.7×10<sup>-4</sup></td><td>1.5×10<sup>-3</sup></td></tr><tr><td>Cl<sup>-</sup></td><td>1.75×10<sup>-3</sup></td><td>1.87×10<sup>-3</sup></td></tr><tr><td>SO<sub>4</sub><sup>2-</sup></td><td>2.33×10<sup>-3</sup></td><td>2.45×10<sup>-3</sup></td></tr><tr><td>NO<sub>3</sub><sup>-</sup></td><td>&lt;8×10<sup>-6</sup></td><td>&lt;8×10<sup>-6</sup></td></tr></tbody></table>
<table><thead><tr><th></th><th>O</th><th>Na</th><th>Mg</th><th>Al</th><th>Si</th><th>S</th><th>Cl</th><th>K</th><th>Ca</th><th>Fe</th><th>Cu</th></tr></thead><tbody><tr><td>1</td><td>67.0</td><td>0.7</td><td>0.3</td><td>0.4</td><td>1.0</td><td>2.1</td><td>0.3</td><td>0.2</td><td>0.2</td><td>28.1</td><td>0.0</td></tr><tr><td>2</td><td>56.1</td><td>0.2</td><td>0.2</td><td>0.2</td><td>0.4</td><td>0.7</td><td>0.1</td><td>0.0</td><td>0.1</td><td>41.9</td><td>0.0</td></tr><tr><td>3</td><td>60.7</td><td>0.4</td><td>0.3</td><td>0.2</td><td>0.3</td><td>0.5</td><td>0.1</td><td>0.1</td><td>0.1</td><td>37.4</td><td>0.0</td></tr><tr><td>4</td><td>62.2</td><td>0.5</td><td>0.3</td><td>0.2</td><td>10.1</td><td>8.1</td><td>0.2</td><td>0.1</td><td>1.1</td><td>15.8</td><td>1.3</td></tr><tr><td>5</td><td>63.7</td><td>0.6</td><td>0.2</td><td>0.3</td><td>9.4</td><td>5.5</td><td>0.2</td><td>0.2</td><td>0.9</td><td>19.1</td><td>0.1</td></tr><tr><td>6</td><td>55.2</td><td>0.5</td><td>0.3</td><td>0.4</td><td>9.1</td><td>12.0</td><td>0.3</td><td>0.2</td><td>1.2</td><td>20.7</td><td>0.2</td></tr></tbody></table>