Role of Magnesium-Stabilized Amorphous Calcium Carbonate in

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The role of magnesium-stabilized amorphous calcium carbonate in mitigating the extent of carbonation in alkali-activated slag Antoine E. Morandeau, and Claire E. White Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.5b02382 • Publication Date (Web): 09 Sep 2015 Downloaded from http://pubs.acs.org on September 15, 2015

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Chemistry of Materials

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The role of magnesium-stabilized amorphous calcium

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carbonate in mitigating the extent of carbonation in alkali-

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activated slag

4 Antoine E. Morandeaua,b, Claire E. Whitea,*

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a

Department of Civil and Environmental Engineering and the Andlinger Center for Energy and

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the Environment, Princeton University, Princeton NJ 08544, USA b

Ecole Supérieure d'Ingénieurs des Travaux de la Construction de Caen (ESITC-Caen), 1, rue Pierre et Marie Curie, 14610 Epron, France

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* Corresponding author: Phone: +1 609 258 6263, Fax: +1 609 258 2799, Email:

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[email protected]

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Postal address: Department of Civil & Environmental Engineering, E-Quad, Princeton

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University, Princeton NJ 08544, USA

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1 ACS Paragon Plus Environment


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1 Abstract

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Oil well cements have received a significant amount of attention in recent years due to their use

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in high-risk conditions combined with their exposure to extremely aggressive environments.

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Adequate resistance to temperature, pressure and carbonation is necessary to ensure the integrity

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of the well, with conventional cements prone to chemical degradation when exposed to CO2

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molecules. Here, the local atomic structural changes occurring during the accelerated

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carbonation (100 % CO2) of a sustainable cement, alkali-activated slag (AAS) have been

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investigated using in situ X-ray diffraction and pair distribution function analysis. The results

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reveal that the extent of carbonation-induced chemical degradation, which is governed by the

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removal of calcium from the calcium-alumino-silicate-hydrate (C-A-S-H) gel, can be reduced by

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tailoring the precursor chemistry; specifically the magnesium content. High-magnesium AAS

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pastes are seen to form stable magnesium-containing amorphous calcium carbonate phases,

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which prevents the removal of additional calcium from the C-A-S-H gel, thereby halting the

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progress of the carbonation reaction. On the other hand, lower-magnesium AAS pastes form

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amorphous calcium carbonate which is seen to quickly crystallize into calcite/vaterite, along with

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additional decalcification of the C-A-S-H gel. Hence, this behavior can be explained by

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considering (i) the solubility products of the various carbonate polymorphs and (ii) the stability

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of amorphous calcium/magnesium carbonate, where due to the higher solubility of amorphous

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calcium carbonate and associated saturation of solution with respect to calcium, additional C-A-

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S-H gel decalcification cannot proceed when this amorphous phase is present. These results may

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have important implications for the use of new cementitious materials in extremely aggressive

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conditions involving CO2 (e.g., enhanced oil recovery and geological storage of CO2),


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particularly due to the ability to optimize the durability of these materials by controlling the

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precursor (slag) chemistry.

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2 Introduction

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Magnesium is known to play a crucial role in amorphous carbonate phases, in particular in

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biogenic synthesis of calcium carbonate where magnesium stabilizes amorphous calcium

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carbonate (ACC).1–3 It has also been shown recently that magnesium plays an important role in

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some sustainable cements, where it is known that magnesium helps prevent carbonation-induced

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degradation of alkali-activated slag (AAS).4 AAS binders are one type of sustainable

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construction material that has lower CO2 emissions compared to ordinary Portland cement (OPC)

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concrete.5 They are produced by mixing a solid aluminosilicate-rich precursor powder with a

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high pH alkaline activator.6 The alkali-activation reaction leads to a mechanically-hard binder,

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with similar physical properties to OPC.6 However, the implementation of alkali-activated

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materials (AAMs), including AAS-based concrete, in the construction industry is hampered by

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the lack of durability data proving adequate long-term performance. The durability of these

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binders is intimately linked with their ability to resist chemical and mechanical degradation.7

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One form of chemical degradation, carbonation, is a naturally occurring process that decreases

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the pore solution pH of cement paste causing decalcification of the gel binder and precipitation

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of calcium carbonate.8-10 Moreover, accelerated carbonation of cementitious materials is seen to

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occur in industrial processes including the oil and gas industry and geological storage of CO2.11

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Cement is used to encase steel pipes to prevent leakage and blowouts, and therefore it comes into



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direct contact with CO2 when supercritical CO2 is employed. However, conventional cement is

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seen to have limited resistance to carbonation in these highly aggressive conditions due to the

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rapid drop in pore solution pH causing dissolution of portlandite (Ca(OH)2) followed by the

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decalcification of the main binding phase, calcium-silicate-hydrate (C-S-H) gel. Therefore, new

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cementitious materials with superior carbonation resistance (limited degradation of the binder

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phases) would be advantageous to use in these natural and accelerated carbonation situations to

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assist with long-term integrity of the structural components (reinforced concrete for the case of

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natural conditions and wellbores for accelerated carbonation).

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Recently, carbonation of synthetic calcium-silicate-hydrate (C-S-H) and calcium-alumino-

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silicate-hydrate (C-A-S-H) gel has been studied using 29Si, 13C and 27Al nuclear magnetic

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resonance (NMR).12 C-S-H gel is the main phase present in OPC, and is thought to possess

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similarities with the C-A-S-H gel present in AAS.13 Sevelsted et al.12 demonstrated that

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aluminum has a similar role to silicon during the carbonation process, whereby it is incorporated

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in the amorphous silica phase mainly as tetrahedral Al(–OSi)4 units as a result of decalcification

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of the C-A-S-H gel, with a small amount of aluminum ending up as penta-coordinated.12 In our

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previous investigation, we investigated the accelerated carbonation (100% CO2) of a synthetic C-

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S-H gel using synchrotron-based high energy X-ray diffraction and subsequent pair distribution

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function (PDF) analysis.14 This in situ measurement allowed for the characterization of the exact

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local structural changes that occurred during carbonation of the gel together with the reaction

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kinetics of the various crystalline and non-crystalline phases. The results revealed the emergence

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of several polymorphs of crystalline calcium carbonate (vaterite and calcite) in addition to the

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decalcified C-S-H gel.14 Furthermore, the use of PDF analysis allowed for a more accurate


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determination of the non-crystalline phases that developed during the accelerated carbonation

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process in comparison to the conventional reciprocal space Rietveld analysis approach. This is

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due to the fact that during carbonation, C-S-H gel transforms rapidly into a predominately

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amorphous silica gel, with most of the calcium being leached out from the C-S-H gel and

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incorporated into ACC12,15 prior to quickly transforming into a more stable crystalline

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polymorph.14

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The local atomic structures of different ACCs have been investigated using a range of techniques

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including X-ray total scattering,16–18 reverse Monte Carlo refinement19 and 43Ca solid-state NMR

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spectroscopy together with molecular dynamics (MD) simulations.3,17,20 These studies have

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revealed that the atomic structures of these ACCs exhibit short-range ordering in the PDF data

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similar to the crystalline polymorphs (below 2.5 Å), however, beyond the nearest-neighbors there

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are limited correlations due to the amorphous nature of this class of material, with no long-range

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ordering above 15 Å.16 Moreover, the MD simulations illustrated that a high magnesium and

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water content in ACC reduces the rate of crystallization.3 Magnesium is known to bind water

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more strongly in its solvation shell compared to calcium,21 and given that the pathway to

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crystallization involves the transition of the hydrous ACC to the anhydrous form,22 a larger

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activation barrier for dehydration (as it is the case for magnesium) will slow down the rate of

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crystallization. However, recent research has also revealed that, for anhydrous magnesium-rich

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calcium carbonate systems, there is limited ability for the carbonate groups to form long-range

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ordered structures.23 Hence, a dual effect is possible in reducing the rate of crystallization for

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magnesium-containing calcium carbonate systems, where both the solvation energy of the

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magnesium ions and the inability of the carbonate groups to form periodic structure being



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responsible for the slower kinetics.

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Apart from water and the presence of magnesium ions, there are other important aspects that

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influence the transformation pathway from ACC to the crystalline polymorphs. The solution pH

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is known to strongly influence the type of crystalline polymorph that forms, with vaterite being

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the preferred polymorph at high pH (pH ~ 9.75 to 10) and calcite at lower pH (pH ~ 9 to 9.5).24

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For the carbonation of cementitious materials in air, other parameters will influence the

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polymorphism of the calcium carbonates such as relative humidity and temperature.25,26

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Although the carbonation process in OPC-based cementitious materials is relatively well

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understood, there is limited knowledge on the carbonation mechanisms occurring in AAMs.10

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Magnesium is thought to play an important role in the resistance of AAS to carbonation,4

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however, additional investigation is required in order to assess the suitability of this class of

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sustainable cement for use in highly aggressive carbonation conditions. Part of the uncertainty

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relates to the role of the magnesium-rich layered double hydroxide (LDH) phases that form in

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slag-based AAMs. These structures are made of layered positively charged aluminum and

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magnesium sheets with interlayers of charge-balancing anionic molecules such as carbonates or

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sulfates, and water.27 Since these phases can capture and exchange carbonate molecules over

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short period of time,28 together with the fact that the carbonate-based LDH (hydrotalcite-like

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phase) is the most stable form (compared to others such as the sulfate-based LDH), the

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hydrotalcite-like phase has been identified as a sink for CO2 thereby limiting the progression of

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carbonation in AAS concrete.10

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In order to assess the suitability of sustainable AAS-based cements for use in aggressive


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carbonation conditions, the phase evolution and stability of AAS pastes during carbonation has

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been elucidated using in situ X-ray diffraction. The local and medium-range atomic ordering of

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four types of AAS paste have been measured using synchrotron-based PDF analysis together

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with the changes induced due to accelerated carbonation (100 % CO2). The effects of magnesium

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have been assess via use of different slag sources (differing in magnesium content), while the

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impact of initial solution pH and free silica in solution has been investigated by augmenting the

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activator chemistry (hydroxide and silicate-based solutions have been assessed). Real-space

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analysis of the data has allowed for the identification of all disordered/amorphous phases

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together with any crystalline polymorphs that emerged, and these data have also been used to

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obtained accurate information on the reaction kinetics.

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3 Materials and methods

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3.1 Alkali-activated slag pastes

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The binders studied in this investigation were prepared using two types of ground granulated

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blast furnace slag as the precursor. The main difference between these sources of slag is the

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magnesium content, and therefore they will be designated by the MgO wt. % (see Tab. 1) as

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quantified by X-ray Fluorescence Spectroscopy (XRF, Rigaku Supermini). For the XRF

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measurement, the slags were dried at 105°C until equilibration of the mass, at which point they

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were mixed with a cellulose binder in order to make flat pellets. These pellets were analyzed in

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the XRF under vacuum.

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In order to identify any crystalline phases present in the neat slags, high-resolution X-ray

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diffraction was performed on 11-BM at the Advanced Photon Source, Argonne National

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Laboratory. Data were collected using a wavelength of 0.413368 Å at a temperature of 100K.

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The diffraction patterns and identified phases are shown in Fig. S1 in the Supporting

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Information. The slags are mainly amorphous with Bragg peaks attributed to merwinite

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(Ca3Mg(SiO4)2), åkermanite (Ca2MgSi2O7) and gehlenite (Ca2Al(AlSi)O7). Quartz (SiO2) and

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calcite (CaCO3) are also present in the 7% MgO slag.

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Table 1: Oxide composition (wt. %) of the two sources of slag used in this investigation, as

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determined using XRF. Slag source

CaO

SiO2

Al2O3

MgO

SO3

(Ca+Mg)/Si

Mg/Ca molar

molar ratio

ratio

13% MgO

36.1

34.5

10.5

12.7

2.7

1.67

0.49

7% MgO

42.5

34.5

11.7

7.3

1.7

1.64

0.24

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The slag powders were activated using sodium hydroxide (denoted H-activated) and sodium

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silicate (S-activated) solutions. These solutions were synthesized by dissolving the anhydrous

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alkali powder/pellets (NaOH and Na2SiO3, reagent grade, Sigma-Aldrich) in deionized water. To

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ensure that the silicate oligomers in the S-activating solution had reached an equilibrated state,

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this solution was mixed for at least 24 hrs prior to synthesis of the paste using a hot plate (~

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40 °C) and a magnetic stirrer bar. For each paste, the Na2O wt. % was 7 (i.e., 7 g of Na2O per

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100g of slag) and the water to slag ratio was 0.44. Pastes were synthesized by manual mixing in 8 ACS Paragon Plus Environment

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50 mL centrifuge tubes for 2 minutes in the laboratory with natural levels of CO2. After mixing,

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the samples were left to cure in the sealed plastic tubes for 24 hrs prior to measurements.

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3.2 X-ray data collection

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For the in situ X-ray PDF measurements, each paste sample was ground just before the

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experiment using a mortar and pestle, then loaded into a 1 mm diameter polyimide capillary and

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sealed using porous glass wool in order to enable the gases (Ar and CO2) to flow through the

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powdered sample. The capillary was then mounted and aligned in the gas cell29 on 11-ID-B at the

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Advanced Photon Source, Argonne National Laboratory, under ambient temperature conditions.

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Each sample was analyzed using a wavelength of 0.2114 Å and a two-dimensional image plate

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detector.30 The detector was located 167 mm from the sample. Data collection commenced under

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a flow of Ar for the first 2 min to obtain data on the initial (non-carbonated) AAS. Then the gas

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was switched to CO2 (100%), and data sets were acquired every 10 seconds to 2 minutes for a

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duration of at least 2 hours after which no significant changes in the local structure were

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detected. A non-carbonated control sample was characterized before and after each experiment to

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make sure that the observed changes were due to carbonation and not from the alkali-activation

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reaction (see Fig. S2 in the Supporting Information). Samples were also measured 24 hrs after

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the in situ carbonation measurements to assess the stability of the amorphous carbonate phases

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that formed. The data conversion from 2D to 1D was carried out using the program Fit2D with

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CeO2 as the calibration material.31,32

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3.3 Data analysis

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Real space analysis was carried out using the PDF data. The PDF, otherwise known as G(r), is

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obtained by taking a sine Fourier transform of the measured total scattering function, S(Q), as

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shown in Eq. (1) , where Q is the momentum transfer given in Eq. (2).33

2 G (r )= π

Q max

∫

Q [S (Q )− 1 ]sin (Qr )dQ

(1)

Q min

195 Q=

4 π sin θ λ

(2)

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Standard data reduction procedures were followed to obtain the PDF using PDFgetX2,34 with a

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Qmax of 20 Å-1. The instrument parameters were elucidated by measuring a nickel standard

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(Sigma-Aldrich) and performing a refinement using the PDFgui software.35 The refined

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instrument parameters were Qbroad = 0.0196 Å-1 and Qdamp = 0.0347 Å-1.

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To quantify the extent of reaction in real space, a methodology similar to that outlined in our

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previous articles was used14,36 where the initial and final PDF data sets from an in situ diffraction

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experiment were employed as the start and end points for quantification. The intermediate data

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sets were then fitted over an r range of 0.1 < r < 50.0 Å using a linear combination of the initial

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and final data sets (Eq. (3)), where α is a fractional number between 0 and 1. G ( r )calc = (1− α ) initial+ (α) final

(3)

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In this investigation, the best fit has been obtained by minimizing Σ[(G(r)calc – G(r)exp)2] with

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respect to α for each intermediate data set (G(r)exp), giving the α value (or extent of carbonation

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reaction), where α = 0 corresponds to the initial data set, and α = 1 corresponds to the final data

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set (fully carbonated). The linear combination method was carried out using Matlab, where α was

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varied in increments of 0.001, and the value of α that gave the best fit (smallest difference

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between simulated and experimental data) being the value reported in the Results and Discussion

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section.

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4 Results and discussion

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4.1 Reciprocal space analysis

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The initial (non-carbonated) X-ray total scattering functions, S(Q), for the alkali-activated slag

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pastes are presented in Fig. 1. These data reveal that there are two distinct phases in the non-

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carbonated pastes; (i) a hydrotalcite-like phase, crystalline magnesium/aluminum LDH (COD

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#900927237), and (ii) alkali-C-A-S-H gel (denoted C-(N)-A-S-H, similar to ICSD #18704138 and

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C-S-H in our previous work14), the main binder phase in AAS (predominately amorphous). The

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dominant hydrotalcite-related peak at Q = 0.783 Å-1 is attributed to the interlayer spacing of the

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LDH structure and reveals that the 003 reflection has a d-spacing of 8.02 Å, which is typical of

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carbonate based LDH39 and therefore indicates that some CO2 is already in the paste. It is clear in

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Fig. 1 that after a relatively short period of curing (24 hrs), the S-activated slags have lower

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amounts of the hydrotalcite-like phase compared to H-activated, as seen by the lower intensity

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hydrotalcite-related peaks in the S-activated slags. Hence, the higher initial pH of the activating

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solution (13.8 for H-activated versus 13.5 for S-activated, obtained using a pH meter 11 ACS Paragon Plus Environment


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immediately after initial mixing) is seen to favor the extent of LDH formation, which is in

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agreement with previous literature on AAS pastes.40 Furthermore, the availability of free silica in

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solution for the S-activated pastes is thought to influence (enhance) the incorporation of

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aluminum in the C-(N)-A-S-H/N-A-S-H gel during the initial stages of the reaction, leading to

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less aluminum being available to form the LDH.

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Figure 1. X-ray total scattering functions, S(Q), of alkali-activated slag pastes prior to

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carbonation.

238 239

As seen in Fig. 2, after 2 hrs of accelerated carbonation the total scattering functions reveal that

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there are large differences in the type of solid crystalline phases that precipitate. The main peak

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attributed to the C-(N)-A-S-H gel at ~ 2.06 Å-1 (from here on denoted as C-A-S-H for simplicity)

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decreases in intensity for all pastes, illustrating that the C-A-S-H gel undergoes decalcification as

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a result of being exposed to high concentrations of CO2, however, the secondary C-A-S-H gel 12 ACS Paragon Plus Environment

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Bragg peaks appear to remain unchanged in the diffraction patterns. Calcite is seen to precipitate

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in both pastes synthesized using 7% MgO slag, with vaterite also being present in the H-

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activated paste at the end of the carbonation measurement. Although vaterite is not present in the

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diffraction pattern for S-activated 7% MgO slag in Fig. 2 (after 2 hrs of being exposed to CO2), it

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was identified as an intermediate phase during the in situ measurement (see Fig. S3 in the

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Supporting Information). Fig. 2 shows that the intensities of the calcite Bragg peaks are much

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greater in the S-activated 7% MgO paste compared to the equivalent H-activated sample. The

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rapid transformation of vaterite to calcite in the S-activated 7% MgO paste is likely attributed to

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either a higher gas flow rate during the experiment and/or finer particle size/looser particle

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packing of the powder in the capillary. For the samples made using the 13% MgO slag, it is

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interesting to note that during the 2 hrs of carbonation these pastes do not form any crystalline

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calcium carbonate polymorphs, and even after 24 hrs no delayed crystallization was observed

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(see Fig. S4 in the Supporting Information).



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Figure 2. X-ray total scattering functions, S(Q), of alkali-activated slag pastes after accelerated

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carbonation.

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Apart from changes to the C-A-S-H gel as a result of carbonation, it has also been reported that

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the hydrotalcite-like phase plays an important role in mitigating the extent of carbonation in AAS

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pastes4 via the transformation of the LDH to huntite (CaMg3(CO3)4). As seen in Fig. 2, the LDH

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is still present in the H-activated 13% MgO slag and the intensity of the Bragg peaks do not

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appear to change (increase or decrease) as carbonation takes place (Fig. S5 in the Supporting

266

Information). The d-spacing of the 003 reflection is located at ~ 7.75 Å, which is in agreement

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with literature for carbonate-based hydrotalcite-like phases.41,42 Furthermore, there is only a

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small shift in the main peak position for the hydrotalcite-like phase (from a Q value of 0.807 Å-1 14 ACS Paragon Plus Environment

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to 0.812 Å-1 corresponding to a d-spacing shift from of 7.79 to 7.74 Å) which indicates that there

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is limited exchange of anions in the layered structure. Hence, this LDH probably contains mostly

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carbonates and some water molecules since sulfate-based LDHs have larger interlayer spacings

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(003 reflection d-spacing ~ 8.6 Å) and hydroxide-based LDHs have smaller interlayer spacings

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(003 reflection d-spacing ~ 7.6 Å).41 Interestingly, Radha and Navrotsky reported a shift of 0.14

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Å (increase in the d-spacing from 7.61 to 7.75 Å) due to the sorption of CO2 in a degassed LDH

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sample.42 Therefore, it is possible that some of the anions originally in the hydrotalcite-like phase

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in AAS are sulfates, and as carbonation progresses the sulfates are replaced by carbonates

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leading to a slight contraction of the d-spacing approaching the value reported by Radha and

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Navrotsky (7.75 Å).42 The H-activated 7% MgO slag paste behaves differently since the intensity

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of the main hydrotalcite-related peak is seen to decrease indicating that the LDH is reacting with

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the low pH solution created by the dissolution of CO2 molecules in the pore liquid43 (see also

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Fig. S6 in the Supporting Information). Hence, for the case of the H-activated 13% MgO sample,

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carbonation leads to some decalcification of the C-A-S-H gel, but does not significantly affect

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the hydrotalcite-like phase. On the other hand, for the lower MgO sample (H-activated 7% MgO)

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both the C-A-S-H gel and the hydrotalcite-like phase are affected.

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4.2 Pair distribution function analysis

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4.2.1 Non-carbonated pastes

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The X-ray PDFs for the four types of AAS paste studied in this investigation are displayed in

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Fig. 3, where both the non-carbonated and carbonated results are given over an r range of 0.7 < r

290

< 30 Å (Fig. 3a). Also provided in this figure are the short range correlations present below 5 Å 15 ACS Paragon Plus Environment


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291

(Fig. 3b). Prior to carbonation, the two types of H-activated slag (7 and 13% MgO) have more

292

intense atom-atom correlations at longer range compared to the S-activated pastes due to the

293

higher hydrotalcite-like phase content. On the other hand, given that C-A-S-H gel in AAS is

294

predominately amorphous13 and that the hydrotalcite-like phase content is negligible in the S-

295

activated slags, a rapid decrease in the ordering is observed at ~ 15 Å for these samples, which is

296

consistent with previous X-ray PDF data on these systems.13

297 298


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(a)

299



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(b)

300 301

Figure 3. X-ray pair distribution functions of alkali-activated slag pastes before and after

302

accelerated carbonation over an r range of (a) 0.7 < r < 30.0 Å and (b) 0.7 < r < 5.0 Å.


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304

To compare the similarities and differences in the local atomic structure for the non-carbonated

305

pastes, the PDF data have been plotted on the same axis, as shown in Fig. 4. This figure clearly

306

shows that the pastes have systematic differences which align with the chemical composition of

307

the initial slags and activating solutions. In particular, the correlation labelled Si-O in the figure

308

at ~1.64 Å (which also contains Al(IV)-O (four coordinated aluminum)) is more intense for the

309

S-activated slags due to the presence of free silica in the activator. For the samples synthesized

310

using 13% MgO slags the Mg-O/ Al(VI)-O correlations at 2.02 Å is higher in intensity due to the

311

higher magnesium content in the slag, while the correlation at 2.35 Å attributed to Ca-O is lower

312

intensity since the higher magnesium is associated with a lower calcium content. The peak at ~

313

3.1 Å is attributed to Si-Si and Si-Al in C-A-S-H gel and to O-O and Mg-Al in the hydrotalcite-

314

like phase.39 This peak appears slightly more intense in the two pastes synthesized using 13%

315

MgO slag and therefore shows that the Mg-Al correlation is apparent in this peak. On the

316

contrary, the ~3.63 Å peak (associated with Ca-Si in C-A-S-H gel)44 is more intense for the

317

pastes synthesized using 7% MgO slag due to the higher calcium content. Hence, these results

318

illustrates that the C-A-S-H gel in the 7% MgO slag pastes have a slightly higher initial Ca/Si

319

ratio, which means that there is more calcium in the C-A-S-H gel available for carbonation,15 in

320

agreement with the XRF results presented in Tab. 1.



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321 322

Figure 4. Comparison of the X-ray pair distribution functions of alkali-activated slag pastes

323

before accelerated carbonation over an r range of 0.7 < r < 5.0 Å.

324

325

4.2.2 Carbonated pastes

326

As the carbonation reaction proceeds there are changes in all atom-atom correlations as

327

illustrated in Fig. 3. Specifically, the C-O correlation (1.26 Å) increases, while the Ca-Si

328

correlation (~3.63 Å) decreases for all samples due to the decalcification of the C-A-S-H gel,

329

which consists of calcium ions being pulled out from the gel structure with subsequent formation

330

of various calcium carbonate polymorphs.14 The decalcification process appears to be less

331

extensive for the 13% MgO slag samples, as seen in Fig. 3b where the intensity of the 3.63 Å Ca-

332

Si correlation is not as strongly affected by carbonation. The two Ca-O correlations at ~2.35 and 20 ACS Paragon Plus Environment

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333

~4.21 Å are seen to increase for all pastes due to carbonation (Fig. 3b), however, these

334

correlations undergo larger changes in the 7% MgO AAS pastes (Fig. 5), which is due to

335

extensive precipitation of crystalline calcium carbonate polymorphs (calcite and vaterite as

336

identified by their diffraction peaks in reciprocal space, see Fig. 2 and Figs. S3, S5 and S6 in the

337

Supporting Information). Furthermore, the formation of crystalline carbonate phases also leads to

338

the emergence of distinct atom-atom correlations at higher r spacings in the PDFs (above ~10

339

Å). The main difference between the carbonated pastes according to MgO content for the short

340

range correlations (< 5Å) is the emergence of a peak at ~3.25 Å in the 7% MgO pastes that does

341

not exist in the 13% MgO samples (see Figs. 3 and 5). This peak is attributed to the formation of

342

vaterite and/or calcite and is mainly associated with the Ca-C and Ca-O correlations in calcite

343

(see Fig. S7 in the Supporting Information for more details).16 As was pointed out by Michel et

344

al. in their study on synthetic ACCs,16 and by Cobourne et al.17 on magnesium stabilized ACC,

345

the peaks at ~2.3 (Ca-O) and ~4.2 Å (Ca-Ca correlation in the Ca-rich framework of the ACC)

346

are present in these amorphous carbonates, however, there is no atom-atom correlation present at

347

~3.2 Å. Hence, these literature results align with the PDF data in Fig. 5, where for 13% MgO

348

pastes there is no distinct correlation at this r spacing (3.2 Å), and therefore these high

349

magnesium slag pastes are seen to form ACC as a result of the accelerated carbonation process.



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350 351

Figure 5. Comparison of the X-ray pair distribution functions of alkali-activated slag pastes after

352

accelerated carbonation over an r range of 0.7 < r < 5.0 Å.

353 354

These amorphous carbonates in the high magnesium pastes are stable not only during the

355

accelerated carbonation measurements, but also 24 hrs after the experiment, as determined via

356

PDF analysis of one of the carbonated sample (H-activated 13% MgO slag, see Fig. S4 in the

357

Supporting Information). As mentioned in the Introduction, crystallization of ACC and AMC is

358

highly correlated with the ability of the (originally hydrated) amorphous phase to dehydrate,22

359

and since these hydrated phases consist of nanoporous networks saturated with water45 the

360

connectivity of these networks and strength of the cation solvation shells play a crucial role in

361

the crystallization process. Hence, since it is seen that the high magnesium slag pastes (13%

362

MgO) form stable ACC as a result of the carbonation process, the magnesium ions are playing a 22 ACS Paragon Plus Environment

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363

stabilizing role in this phase, preventing the amorphous phase from undergoing crystallization.

364

On the other hand, for the lower magnesium slag (7% MgO), the ACC that forms is not stable,

365

and quickly transforms into the crystalline polymorphs, due to the lack of available magnesium

366

that is incorporated into the ACC structure. The hydrotalcite-like phase is seen to partially

367

dissolve during the carbonation reaction in the H-activated 7% MgO paste (Fig. 1), and therefore

368

magnesium is being released into solution. However, LDH dissolution cannot be the sole source

369

of magnesium for stabilizing ACC in the H-activated 13% MgO paste, since the S-activated 13%

370

MgO paste (limited amount of LDH) is also seen to form a stable amorphous phase during

371

carbonation. Further analysis is required to ascertain all the phase(s) containing magnesium in

372

these alkali-activated slag pastes.

373 374

It is interesting to note that for the higher magnesium pastes, the carbonation process does not

375

appear to decalcify the C-A-S-H gel as extensively, as seen in Fig. 5 by the higher intensity Ca-

376

Si correlation at ~ 3.63 Å for these carbonated samples. Hence, although the 7% MgO slag

377

samples originally possess more intense Ca-Si correlations, as evident in Fig. 4, the carbonation

378

process removes a significant amount of calcium from the gel. Fig. 6 displays PDFs for the

379

various pastes obtained during the carbonation reaction, at the stage when the 7% MgO pastes

380

are still amorphous (no calcite or vaterite has precipitated). This figure reveals that even for the

381

lower magnesium slag pastes, ACC is present during the initial stages of the carbonation process,

382

as evident by the similarities between the low and high magnesium pastes after 3 min of reaction

383

(S- activated) and 35 min (H-activated). Hence, the results reveal that the rate limiting step for

384

decalcification of the C-A-S-H gel under extremely aggressive carbonation conditions (100%

385

CO2) is the crystallization of calcium carbonate. If ACC is stabilized by magnesium, this 23 ACS Paragon Plus Environment


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386

crystallization process does not occur within the time frame analyzed in this investigation, and

387

therefore additional decalcification does not proceed.

388 389

Figure 6. X-ray pair distribution functions of alkali-activated slag pastes obtained during the

390

carbonation reaction, at the denoted times when all samples are amorphous (no calcite or vaterite

391

has precipitated in the 7% MgO pastes at this stage).


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393


4.3 Extent of reaction

394

The real space quantification of the extent of reaction as described in the Materials and Methods

395

section is presented in Fig. 7, revealing the rate of carbonation reaction for all samples (S- and

396

H-activated, 7 and 13% MgO slag). Fig. 7 confirms the existence of a two-step carbonation

397

process for the H-activated 7% MgO slag paste. The plateau in the extent of reaction data for this

398

sample (at ~ 37 min) coincides with the commencement of crystallization of calcite/vaterite (see

399

Fig S6 in Supporting Information). Therefore, prior to 37 min the C-A-S-H gel is undergoing

400

decalcification, with the Ca2+ ions leached out of the gel being incorporated into ACC (as seen in

401

Fig. 6). For the other three pastes, the mechanisms are less clear (i.e., one or two step) and the

402

curves look like imperfect exponential-shaped functions. However, it is likely that the S-

403

activated 7% MgO slag paste does contain a two-step mechanism and the reason it is not visible

404

in Fig. 7 is due to the rate at which carbonation occurs (attributed to a faster CO2 flow rate and/or

405

looser particle packing in capillary). In Fig. 6, it is clear that there is an amorphous intermediate

406

structure after 3 min of carbonation for this sample (S-activated 7% MgO slag), which

407

subsequently crystallizes at later times. For all the four types of AAS paste, the “final” endpoint

408

is reached between 20 and 80 minutes after the beginning of the accelerated carbonation

409

reaction. This means that this 100% CO2 accelerated method is extremely aggressive (as

410

expected) compared to regular accelerated carbonation methods using 1% CO2 (which need

411

weeks to reach steady-state10).



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412 413

Figure 7. Extent of carbonation reaction for the alkali-activated slag pastes as a function of time.

414

415

4.4 Carbonation mechanism

416

Based on the results highlighted in the previous sections, there is substantial evidence regarding

417

the nature of accelerated carbonation (100%) in AAS paste. At such high CO2 partial pressures

418

there is limited influence from the hydrotalcite-like phase, as seen by the unaffected Bragg peak

419

in the H-activated 13% MgO paste after carbonation (see Fig. S5 in the Supporting Information).

420

The LDH does appear to be affected by carbonation in the lower magnesium system (H-activated

421

7% MgO paste), via a reduction in intensity on the low-Q peak (see Fig. S6 in the Supporting

422

Information). However, as was seen in the PDF data, the C-A-S-H gel still decalcifies at a similar

423

rate to the S-activated paste, and therefore the hydrotalcite-like phase does not have a significant


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424

influence on the carbonation reaction in this sample. It remains unclear as to the reason why the

425

hydrotalcite-like phase in the high magnesium system is not affected by carbonation (Bragg peak

426

does not shift or change in intensity) while the lower magnesium analogue does change

427

substantially. Possible reasons include differences in solution pH due to dissolution of CO2 and

428

the type of anions originally present in the hydrotalcite-like phase interlayer spacing.

429 430

An overview of the mechanism of carbonation according to magnesium content of the slag

431

precursor is given in Fig. 8. Irrespective of the magnesium content in the slag or the type of

432

activator used, all AAS pastes are seen to undergo decalcification of the C-A-S-H gel and

433

formation of ACC during the initial stages of carbonation (Fig. 8). This process is best analyzed

434

via the atom-atom correlation present in the PDF data at 3.63 Å, which is attributed to the Ca-Si

435

correlation. As decalcification of the C-A-S-H gel proceeds, this correlation reduces in intensity.

436

Decalcification occurs due to the change in solubility of the C-A-S-H gel attributed to a

437

reduction in the pore solution pH (CO2 molecules acidify the solution). For the pastes

438

synthesized using the high magnesium slag (13% MgO), the Ca-Si correlation decreases in

439

intensity less compared to the low magnesium pastes (7% MgO), although a similar decrease in

440

intensity is seen prior to crystallization of calcite and/or vaterite (see Fig. 6). These results imply

441

that the extent of decalcification of the C-A-S-H gel is controlled by the solubility of the reaction

442

products, since once ACC crystallizes (as is the case for the low magnesium pastes (7% MgO)),

443

additional decalcification occurs (mechanism displayed in Fig. 8). The solubility products for

444

calcite, vaterite and ACC have been reported in the literature (Ksp of 10-8.48, 10-7.91 and 10-6.39,

445

respectively),46,47 although a recent investigation on ACC in dilute solutions reported lower

446

solubility values for ACC compared to the investigation by Brečević and Nielsen (Ksp values



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

447

ranging from ~ 10-7.64 to 10-7.48),48 citing that the structural ordering of ACC depends on the

448

solution concentration, with highly concentrated solutions leads to the formation of ACC that is

449

less stable and exhibits a lower degree of short-range ordering. Given the aggressive carbonation

450

environment used in this current investigation, the higher solubility product for ACC (10-6.39) is

451

likely to be closer aligned with the experimental conditions. However, given the lack of data

452

available in the literature on the solubility product of magnesium-stabilized ACC, it is uncertain

453

by how much this value will change relative to the pure ACC.

454

455 456

Figure 8. Schematic of the accelerated carbonation reaction of alkali-activated slag. The

457

mechanism of carbonation depends on the magnesium content of the slag precursor, with high

458

magnesium slag pastes leading to the formation of stable Mg-ACC. For low magnesium slag

459

pastes, the initially formed ACC quickly transitions to the crystalline polymorphs (calcite and/or

460

vaterite). This transition, and associated differences in solubility of the calcium-rich phase(s)

461

(ACC versus calcite/vaterite) leads to additional decalcification of the C-A-S-H gel.

462 463 464

For the low magnesium AAS pastes (7% MgO), the metastable ACC is seen to quickly transform

465

into crystalline polymorphs. Initially, when ACC is still present, the amount of Ca2+ ions and


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466

carbonate molecules in solution will be relatively high due to the high solubility of this phase

467

(10-6.39) and the continual supply of CO2 in the gas stream. Hence, if the solubility of ACC is

468

higher than that of C-A-S-H gel (at a given pH value), leading to a significant amount of Ca2+ in

469

solution, then the system will reach a point at which C-A-S-H gel decalcification halts since the

470

solution is saturated with Ca2+ ions. This behavior appears to have happened for the AAS pastes

471

synthesized using 13% MgO slag (Fig. 5), where decalcification of the C-A-S-H gel (reduction

472

in Ca-Si correlation) is not as extensive as for the 7% MgO pastes. On the other hand, if calcite

473

and vaterite begin to form at the expense of ACC, more Ca2+ ions and carbonate molecules will

474

be removed from the solution due to the relatively low solubilities of these crystalline calcium

475

carbonate phases compared to ACC. If the amount of Ca2+ ions in solution falls below the

476

solubility level of the C-A-S-H gel, decalcification of the gel will recommence. This behavior is

477

seen to occur for the 7% MgO pastes in Fig. 5.

478 479

Typically the stability of ACC is relatively low, with crystallization occurring in a matter of

480

minutes to hours depending on the synthesis method and environmental conditions.21 However,

481

stable ACC has been synthesized using stabilizing agents such as citrate, magnesium and

482

silica.21,49,50 Given the two sources of slag have the same silica content, it is highly likely that

483

magnesium is stabilizing the ACC phase in the high MgO pastes (13%) due to the higher Mg/Ca

484

molar ratio in the starting slag (Table 1). Hence, the extent of carbonation of AAS pastes in

485

extreme conditions (100% CO2) is controlled by the stability of ACC, with magnesium leading to

486

a more stable amorphous carbonate phase, and therefore inhibiting the progress of carbonation.

487

Over the time scales investigated in this study (24 hrs) the magnesium-stabilized ACC is seen to

488

be resistant to crystallization, with a previous investigation reporting that synthetic amorphous



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

489

magnesium carbonate (AMC) crystallizes on the order of months, not days (as is the case for

490

ACC).21

491

492

4.5 Broader implications

493

The results presented in this article have potential implications for the use of cement/concrete in

494

highly concentrated CO2 environments, including in the oil and gas industry and in carbon

495

sequestration, where the steel pipes are surrounded by cement to prevent leaks and blowouts.

496

These cements are exposed to supercritical CO2, and therefore suffer extensively from

497

carbonation.51 If the extent of the decalcification of the C-S-H/C-A-S-H gel can be mitigated in

498

these extremely aggressive conditions, specifically by preventing the precipitation of calcite

499

and/or other crystalline polymorphs, the integrity of the cement may be prolonged. Hence, AAS

500

pastes synthesized using high magnesium slags (e.g., 13% MgO) may be ideal for use in such

501

industries, and further research is required to fully assess the suitability of these alternative

502

cements for such applications.

503

504

5 Conclusion

505

In this paper, we investigated the accelerated carbonation (100 % CO2) for four types of alkali-

506

activated slag using in situ X-ray diffraction and pair distribution function analysis. For highly

507

aggressive accelerated carbonation conditions, such as those present during enhanced oil

508

recovery and geological storage of CO2, a high magnesium content slag (13% wt. MgO) limits

509

the extent of decalcification of the C-A-S-H gel, as seen by analysis of the Ca-Si correlation (at ~


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510

3.6 Å) in the PDFs. Furthermore, a high MgO content favors the formation of amorphous

511

calcium carbonate over calcite or vaterite. This implies that due to the relatively high solubility

512

of ACC (compared to calcite and vaterite), further decalcification of the C-A-S-H gel does not

513

occur since the solution is already fully saturated with Ca2+ ions and carbonate molecules. The

514

extent of carbonation is controlled by the stability of ACC, with magnesium leading to a more

515

stable amorphous carbonate phase, and therefore inhibiting the progress of carbonation. The

516

hydrotalcite-like phase does not appear to play a significant role as a CO2 sorbent in these

517

cements, but it may due to the 100% CO2 accelerated conditions. The results presented in this

518

paper have potential implications for the use of cement/concrete in highly concentrated CO2

519

environments, including in the oil and gas industry and in carbon sequestration, where AAS

520

pastes synthesized using high magnesium slags may be ideal due to the ability to limit the extent

521

of decalcification, however, further research is required to fully assess the suitability of these

522

alternative cements for such applications.

523 524

Acknowledgments

525

This work was financially supported by the Princeton E-ffiliates Partnership award, Andlinger

526

Center for Energy and the Environment (Princeton University). We gratefully acknowledge the

527

assistance and guidance provided by Karena Chapman and Kevin Beyer for their help with the

528

gas cell. The 11-ID-B and 11-BM beam lines are located at the Advanced Photon Source, an

529

Office of Science User Facility operated for the U.S. DOE Office of Science by Argonne

530

National Laboratory, under U.S. DOE Contract No. DE-AC02-06CH11357.

531 532

Supporting Information Available



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

533

High-resolution X-ray diffraction patterns of the two slag precursors used in this investigation.

534

Evolution of the X-ray diffraction patterns of S-activated 7% MgO slag and H-activated 7% and

535

13% MgO slags. As a function of the extent of carbonation. X-ray pair distribution functions for

536

the non-carbonated control sample (H-activated 7% MgO slag). X-ray pair distribution functions

537

for the carbonated control sample (H-activated 13% MgO slag). X-ray pair distribution functions

538

of calcite, vaterite and amorphous calcium carbonate, with associated major atom-atom

539

correlations labelled. This material is available free of charge via the Internet at

540

http://pubs.acs.org.

541 542

References

543

(1) Loste, E.; Wilson, R. M.; Seshadri, R.; Meldrum, F. C. The Role of Magnesium in Stabilising

544

Amorphous Calcium Carbonate and Controlling Calcite Morphologies. J. Cryst. Growth

545

2003, 254, 206–218.

546

(2) Politi, Y.; Batchelor, D. R.; Zaslansky, P.; Chmelka, B. F.; Weaver, J. C.; Sagi, I.; Weiner, S.;

547

Addadi, L. Role of Magnesium Ion in the Stabilization of Biogenic Amorphous Calcium

548

Carbonate: A Structure−Function Investigation. Chem. Mater. 2010, 22, 161–166.

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(3) Singer, J. W.; Yazaydin, A. Ö.; Kirkpatrick, R. J.; Bowers, G. M. Structure and

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Transformation of Amorphous Calcium Carbonate: A Solid-State 43Ca NMR and

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Computational Molecular Dynamics Investigation. Chem. Mater. 2012, 24, 1828–1836.

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(4) Bernal, S. A.; Nicolas, R. S.; Myers, R. J.; Gutiérrez, R. M. de; Puertas, F.; van Deventer, J.

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S. J.; Provis, J. L. MgO Content of Slag Controls Phase Evolution and Structural Changes

554

Induced by Accelerated Carbonation in Alkali-Activated Binders. Cem. Concr. Res. 2014,

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57, 33–43.


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(5) Duxson, P.; Provis, J. L.; Lukey, G. C.; van Deventer, J. S. J. The Role of Inorganic Polymer

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Technology in the Development of “Green Concrete.” Cem. Concr. Res. 2007, 37, 1590–

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(6) Provis, J. L.; van Deventer, J. S. J. Alkali Activated Materials - State-of-the-Art Report, RILEM TC 224-AAM. Springer/RILEM, Dordrecht, 2014. (7) Bernal, S. A.; Provis, J. L. Durability of Alkali-Activated Materials: Progress and Perspectives. J. Am. Ceram. Soc. 2014, 97, 997–1008. (8) Morandeau, A.; Thiéry, M.; Dangla, P. Investigation of the Carbonation Mechanism of CH

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and CSH in Terms of Kinetics, Microstructure Changes and Moisture Properties. Cem.

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Concr. Res. 2014, 56, 153–170.

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(9) Hyvert, N.; Sellier, A.; Duprat, F.; Rougeau, P.; Francisco, P. Dependency of C–S–H

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Carbonation Rate on CO2 Pressure to Explain Transition from Accelerated Tests to Natural

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Carbonation. Cem. Concr. Res. 2010, 40, 1582–1589.

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(10) Bernal, S. A.; Provis, J. L.; Brice, D. G.; Kilcullen, A.; Duxson, P.; van Deventer, J. S. J.

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Accelerated Carbonation Testing of Alkali-Activated Binders Significantly Underestimates

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Service Life: The Role of Pore Solution Chemistry. Cem. Concr. Res. 2012, 42, 1317–

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1326.

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(11) Shen, J.; Dangla, P.; Thiery, M. Reactive Transport Modeling of CO2 through Cementitious

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Materials under CO2 Geological Storage Conditions. Int. J. Greenh. Gas Control 2013, 18,

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75–87.

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(12) Sevelsted, T. F.; Skibsted, J. Carbonation of C–S–H and C–A–S–H Samples Studied by 13C, 27

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Graphical Abstract

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Role of Magnesium-Stabilized Amorphous Calcium Carbonate in

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