Microstructural changes in alkali-activated slag mortars induced by accelerated carbonation

Cement and Concrete Research 100 (2017) 214–226

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Microstructural changes in alkali-activated slag mortars induced by accelerated carbonation

MARK

Ning Li, Nima Farzadnia, Caijun Shi⁎ College of Civil Engineering, Hunan University, Changsha 410082, China

A R T I C L E I N F O

A B S T R A C T

Keywords: Alkali-activated slag Carbonation Microstructure C-S-H

This study investigated microstructural changes in alkali-activated slag (AAS) mortars after carbonation using XRD, FTIR, DTG, 1H NMR and SEM examinations. The results showed that decalcification of C-S-H was the main reaction in carbonation of AAS. The C-S-H with a low Ca/Si was more vulnerable to decalcification in AAS samples activated by waterglass. Besides, AAS mortars demonstrated a lower CaCO3 formation compared to Portland cement mortars. Calcite and vaterite were the major CaCO3 polymorphs produced by carbonation of AAS precipitated mainly in gel pores and spaces in C-S-H interlayers. Meanwhile, the carbonation also caused a certain volume of contraction in AAS mortars.

1. Introduction

[7,8]. Although it is convinced that carbonation takes place directly in the C-S-H gel in AAS binders, effects of carbonation on microstructure of AAS mortars are not sufficiently clear. Alkali activation of slag with sodium hydroxide/waterglass was recommended by researchers due to production of C-S-H with a lower Ca/Si resulting in formation of a very cohesive and homogeneous structure [9,10]. During carbonation, however, the low Ca/Si may adversely affect the microstructure of AAS as the decalcification process mostly targets the C-S-H. According to He et al. [11], carbonation decreased Ca/Si from an original value of 1.1–1.2 to 1.0 in C-S-H from activation of slag. The decrease in Ca/Si by carbonation was also reported in [12,13] resulting in shrinkage and an increase in the number of gel pores and micropores. Shi et al. [10] also showed a remarkable shrinkage and increased porosity in AAS which were related to the low Ca/Si of C-S-H. They also reported that almost no crystal phase formed by carbonation. However, the effect of activator type and dosage on the carbonation of AAS is still dubious due to dissimilar Ca/Si from C-S-H as well as secondary products generated by different activators. Puertas et al. [8] reported higher accelerated carbonation depths in AAS mortars activated with sodium silicate than that of sodium hydroxide. In their study, C-S-H had a lower Ca/Si (~ 0.8) in waterglass-activated AAS than those activated with NaOH-only solution (Ca/Si ratio ~ 1.2). They concluded that the higher Ca/Si, along with a reduced silicate chain length observed in NaOH-activated slag, favored the formation and precipitation of a higher amount of carbonation products which filled pore spaces and decreased the diffusivity of CO2 within the material. Higher alkalinity of activators also decreased the carbonation rate by modifying the rendered C-S-H [5,14]. Authors stated that the

Carbonation includes complex physical and chemical reactions that can reduce durability of alkali-activated slag (AAS) at a rate higher than that of Portland cement (PC) concrete [1–3]. Many studies have investigated the carbonation and its governing mechanisms in ordinary Portland concrete. Carbonation of PC takes place when CO2 from the atmosphere diffuses into the pore network of the matrix and reacts with calcium hydroxide (CH), calcium silicate hydrate (C-S-H), calcium aluminate hydrate (C-A-H) and ettringite promoting the formation of calcium carbonate polymorphs through a decalcification process [4]. The carbonation can be different in alkali-activated slag as it belongs to a Me2O-MeO-Me2O3-SiO2-H2O system, in which hydration products mainly consist of C-S-H gel with a lower Ca/Si compared to that of PC. Song et al. [5] reported compressive strength loss in carbonated AAS due to formation of a low cohesive silica gel with a soft matrix; and a larger shrinkage as a result of C-S-H decalcification. They showed that the decalcification of C-S-H was the main reaction in the carbonation as the low Ca/Si of the binders reduced the nucleation and crystal growth of CH. Aperador et al. [6] illustrated that during exposure of specimens to CO2, Na+ slowly spoiled AAS matrix by forming more soluble compounds such as natron. With the absence of CH, Ca2 + from the C-SH gel was the principal source of calcium ions in the pore solution to restrain the pH. Considering comparatively lower Ca/Si in AAS than that in PC, it can be concluded that AAS has a lower capacity to maintain pH of the pore solution than PC. Due to destabilization of C-SH in a low pH environment, the Ca2 + from C-S-H in AAS concrete was transformed into Ca(OH)2, and ultimately to CaCO3 on exposure to CO2 ⁎

Corresponding author. E-mail addresses: [email protected] (N. Farzadnia), [email protected] (C. Shi).

http://dx.doi.org/10.1016/j.cemconres.2017.07.008 Received 2 February 2017; Received in revised form 11 July 2017; Accepted 25 July 2017 0008-8846/ © 2017 Published by Elsevier Ltd.

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100

Cumulative volume (%)

increased alkalinity was beneficial in dissolution and poly-condensation of the species from the slag, which accelerated precipitation of C-S-H gel in the system. Song et al. [5] also suggested that an increase in the activator dosage led to a higher formation rate of C-S-H at an accelerated pace, which increased compaction, compressive strength and carbonation resistance of the microstructure. As the Na2O content in the activator was fixed, the carbonation rate of AAS slowed down as the modulus of waterglass increased from 0.75–1.0 [15,16]. This may be due to an increased reaction degree of the available slag caused by the increase in the activator content. Based on the available literature, it can be stated that the carbonation of AAS depends on C-S-H characteristics varying with type and concentration of activators. There is also uncertainty about type of carbonation products as well as effects of secondary products of activation on carbonation when different activator types are used. During the carbonation process, natron and calcium carbonate polymorphs namely; calcite, vaterite and aragonite were identified as the main carbonation products of waterglass-activated slag binders [8]. However, Bernal et al. [15] reported that calcite was the only calcium carbonate polymorph in carbonated waterglass-activated slag. He et al. [11] reported the formation of calcite and vaterite, but no aragonite was detected in the carbonated waterglass-activated slag. Hydrotalcite is a secondary reaction product observed in slag activated with NaOH and waterglass [17,18]. However, reflections corresponding to hydrotalcite were detected after one day in the NaOH-activated slag, whereas such reflections were clearly identified in the NaOH/waterglass-activated slag after 6 months [17,18]. On account of a faster reaction rate in the early stage of NaOHactivated slag, large hydrotalcite crystals formed in NaOH-activated slag although the chemical compositions were similar in both NaOH or waterglass-activated slag materials [17]. Hydrotalcite was also observed during accelerated carbonation of NaOH/waterglass-activated slag using high CO2 concentrations [19]. León et al. [20] found that the formation of double layered hydroxides with a hydrotalcite-type structure increased CO2 absorption. Bernal et al. [21] also reported that the larger formation of hydrotalcite significantly contributed to enhancing performance of AAS binders when exposed to high CO2 concentrations. Several studies focused on the carbonation rate and mechanisms, however, microstructural changes of AAS after carbonation received less attention. The overarching purpose of this study is to investigate the effect of activator type and concentration on microstructural changes of AAS with comparison to PC. The compressive strength, microstructure of the main reaction products and pore structure of these cements before and after carbonation were investigated. Findings of this study can further improve understanding of the characteristics and mechanism of carbonation in AAS.

80 60 40 20 0 1

10

100

Particle size (µm) Fig. 1. Particle size distribution of GGBS and cement.

10

20

30

40

2

50

60

70

80

(degree)

Fig. 2. XRD pattern of GGBS.

Fig. 3. SEM picture of GGBS particles.

phase. The GGBS particles had angular shape as shown in Fig. 3. River sand with a maximum particle size of 2.36 mm, a fineness modulus of 2.75, and apparent density of 2530 kg/m3 was used to prepare mortar samples. In this study, sodium hydroxide and sodium silicate solution were used as alkaline activators. Sodium hydroxide (NaOH) in pellet-form was an industrial-grade with purity of 99 ± 1%. Also, sodium silicate solution used was an industrial grade sodium with a chemical composition of 8.3% Na2O, 26.5% SiO2 and 65.2% H2O. The alkaline activator consisted of sodium silicate and sodium hydroxide with a molar modulus (SiO2/Na2O molar ratio) of 0 (NaOH-only solution), 0.5, 1.0, 1.5 prepared 24 h prior to use.

2. Materials and methods 2.1. Materials A vitreous ground granulated blast furnace slag from a local steel company was used. Portland cement PI 42.5 in compliance with Chinese standard GB 175-2007 was also used as a reference. The chemical composition of slag and cement was determined by X-ray fluorescence (XRF) and shown in Table 1. The particle size distribution was determined by laser particle size analyzer as shown in Fig. 1. Fig. 2 shows the XRD pattern of the GGBS indicating its dominant amorphous

2.2. Mixture proportions and sample preparation

Table 1 Chemical composition of blast-furnace slag and cement, wt%.

Slag Cement

GGBS Cement

SiO2

Al2O3

CaO

MgO

K2O

Fe2O3

Na2O

SO3

LOI

33.81 21.90

14.78 4.81

38.81 65.15

7.09 1.95

0.44 –

0.36 3.41

0.26 0.65

2.49 0.51

1.40 1.62

The slag was activated by mixing with the alkaline activator at a water-to-binder ratio of 0.47 by mass of slag. For all groups, Na2O in activator was kept constant at 4% of mass of slag, while SiO2 content was changed to reach the proposed modulus. The sand to binder ratio of 2.25 was used. The mixture proportions of AAS and PC mortars are 215

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shown in Table 2. The required water in the experiment was mixed with sodium silicate solution prepared in advance, then cooled to room temperature. The activator solution was placed in a stirring pot and then the weighed slag was added. After immediate stirring at low speed for 30s, the sand was added during second 30s. A high-speed mixing was applied in the third 30s followed by a stop of 30s to scrap the mortar off the blade and walls into the pot. The mortar was then stirred for another 60s before casting into 40 × 40 × 40 mm molds according to the Chinese standard GB 17671-1999 (similar to the European standard UNE-EN 197-1). After casting, the samples were kept in the room at ambient temperature range of 20 ± 2 °C and humidity higher than 90% for 24 h. Later, the demolded samples were kept in a steam curing container at 80 °C and cured for 48 h. Then, they were put into a carbon dioxide curing chamber with a temperature of 20 ± 2 °C and humidity of 65 ± 5%. The AAS and PC mortar samples were carbonated in the curing chamber with carbon dioxide concentration of 20 ± 0.2% in compliance with the Chinese standard GB/T 50082-2009.

(d) Differential thermograms analysis (DTG). The samples in powder form (about 10 mg per each sample) were heated from 20 to 1200 °C in a nitrogen atmosphere at rate of 10 °C/min in a TGA instrument (Netzsch STA 409PC). (e) 1H Nuclear magnetic resonance (1H NMR). NMR test was conducted on 40 × 40 × 40 mm mortar samples before and after carbonation. First, the samples were placed in an oven at 60 °C for 24 h. Once the samples cooled to room temperature (20 °C), a vacuum saturation device was used to saturate the mortar samples. The vacuum pressure value was 0.1 MPa and the pump down time was 4 h. Then, the samples were kept in distilled water for 24 h before repeating the process for the NMR relaxation measurement. Prior to NMR relaxation measurement, the samples were dried by wiping off moisture from the surface and wrapped with a preservative film. Comparison analysis was done by the MacroMR12-150H-I analysis system. Porosity and pore size distribution analysis of the samples were conducted at a resonant frequency of 12.798 MHz, a magnet strength of 0.3 T, and a magnet temperature of 32 °C. For the NMR measurements of water saturated mortar, calibrations were made using standard scale samples to convert the signal strength to porosity. The SIRT inversion algorithm was used to calculate the T2 spectrum of each sample by using the echo attenuation signal obtained from the sample after the saturated water. T2 spectrum reflects the distribution of pore sizes. Also, scanning image analysis was performed using the NMR and analyzed by ImageJ software. (f) Scanning electron microscopy (SEM). A scanning electron microscope (HITACHI S-4800) was used to trace the morphology changes in the microstructure of the matrix. After compressive strength test, freshly broken samples were collected then dried at 60 °C for 24 h and coated with gold before observation using ION SPUTTER E1045.

2.3. Test procedure

3. Results and discussion

To ensure full carbonation of samples, the carbonation depth was measured every 3 days by splitting each sample in half along the center line. Then, the inner surface of samples was sprayed with 1% phenolphthalein alcohol solution. The procedure continued until no color change was observed on the split surface as a sign of carbonation. In order to elaborate the microstructural changes in AAS mortars and to compare them with Portland cement mortars, the following tests were conducted:

3.1. Effect of carbonation on compressive strength

Table 2 Mixture proportion of alkali-activated slag and cement mortar. Mortar

Composition

Modulus of activators

Water/ binder

Sand/ binder

PC WG-0 WG-0.5

Cement Slag + NaOH Slag + NaOH + sodium silicate Slag + NaOH + sodium silicate Slag + NaOH + sodium silicate

– 0 0.5

0.47 0.47 0.47

2.25 2.25 2.25

1.0

0.47

2.25

1.5

0.47

2.25

WG-1.0 WG-1.5

Compressive strength of mortars with different modulus of waterglass before and after carbonation was measured and compared with that of Portland cement mortar as shown in Fig. 4. It can be seen that the carbonation increased the compressive strength of Portland cement mortar by up to 13.9%. The results are consistent with previous studies [24–27]. Carbonation of PC was mainly related to the chemical reaction of Ca(OH)2 and C-S-H with CO2 and formation of CaCO3 particles. CaCO3 particles are generally micro-sized crystals which can fill pores in the matrix, reduce the porosity; and hence improve the compressive strength [24–27]. On the contrary, the carbonation had an adverse effect on AAS mortars, which experienced 13.6%, 26.9%, 26.5%, and 25.8% decrease

(a) Compressive strength test. Uncarbonated and carbonated samples (40 × 40 × 40 mm) were tested under compression at a loading rate of 2.4 kN/s and an average of results from at least three samples was reported. (b) X-ray diffraction (XRD). XRD analysis was performed on the samples before and after carbonation. The collected powders were dried in an oven at 60 °C for 24 h, then sieved through a 325 μm sieve. The applied pre-conditioning was for a standardisation purpose to reach a known moisture state in samples. Since the drying process may affect microstructural results [22], careful sampling was performed according to [23]. The fine powders were analyzed using Philips X-ray diffractometer with CuKα radiation. The samples were step-scanned from 3 to 70° (2θ) at a rate of 5°/min (with step of 0.004°). (c) Fourier transform infrared spectroscopy (FTIR). The same powder for XRD test was used for FT-IR measurement using a ThermoScientific IS10 FT-IR workstation and by employing the conventional KBr disc method. Approximately 1 mg of the sample powder was ground together with 100 mg of IR-grade KBr for 5 min and pressed into a thin disc. Each sample was tested at a resolution of 2 cm− 1 with 32 scans. The blank KBr pellet was tested at the same time as reference.

Compressive strength (MPa)

80 Uncarbonated Carbonated

70 60 50 40 30 20 10 0 PC

WG-0

WG-0.5

WG-1.0

WG-1.5

Fig. 4. Compressive strength of AAS and PC mortars before and after carbonation.

216

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carbonate was mainly in the form of hexagonal calcite and vaterite. However, the intensity of peaks associated to vaterite was less in waterglass activated slag than that of its counterpart. The alteration of polymorphs in PC and AAS was reported to be as a result of relative humidity, Ca/Si, the Mg content, and carbonation conditions [29–33]. There are three common crystalline forms of calcium carbonate: calcite, aragonite and vaterite. Calcite crystals are relatively complete with good stability while vaterite is of poor crystallinity and low stability [29,34]. Stability of aragonite falls between the calcite and vaterite, and generally crystal polymorphic transformation occurs at higher pressure [29]. During the exposure of PC to 20% CO2 until complete carbonation (in a period of 45 days), vaterite may transform into more thermodynamicly stable phases of calcite and aragonite. This is consistent with previous studies [35,36], which showed that the polymorphism of calcium carbonate may vary in a mineral system. However, at a constant relative humidity, Black et al. [31] showed that calcite and vaterite were major polymorphs of CaCO3 in carbonation of C-S-H with Ca/Si between 0.67 and 1.33. This can well explain the dominant formation of calcite and vaterite in carbonated AAS shown in Fig. 5(b) as the Ca/Si (~ 1.2) is reportedly lower than that of PC [8]. Consumption of Mg through formation of hydrotalcite (Fig. 5) may also result in precipitation of calcite rather than aragonite as the Mg/Ca lowers in the pore solution. A Mg/Ca of 2 was suggested as a boundary between formation of aragonite and calcite [32]. Carbonation condition is another influential factor in formation of polymorphs. Bernal et al. [33] showed that calcite, vaterite, and aragonite coexisted in AAS under natural carbonation (0.03–0.04% CO2) after a duration of 3 years. On the contrary, calcite and vaterite were the dominant polymorphs of CaCO3 when the duration prolonged to 7 years [9]. It is in agreement with a study by Bernal et al. [19] that reported metastable vaterite and aragonite would change into calcite after a long-term carbonation. Calcite and vaterite were also the dominant polymorphs of CaCO3 in accelerated carbonation of AAS [11]. Meanwhile, a quick carbonation rate in AAS (9 days to complete carbonation) may also result in changes in transformation of the polymorphs. As can be seen from Fig. 5, the intensity of calcite was higher in PC samples comparing to that of AAS. The intensity was 10.71% higher than major peak of calcite at around 2θ of 29° in AAS samples. The different residual compressive strength of mortars can be partially explained by formation of different polymorphs and contents of CaCO3 in the matrix. In this study bicarbonates (such as nahcolite) were not detected, although precipitation of bicarbonates was reported previously in AAS exposed to accelerated carbonation [33]. This may be related to the sampling procedure prior to XRD test. Bicarbonates are prone to thermal decomposition, albeit, 60 °C as in sampling preparation could not change crystal phase of carbonates [37,38].

in strength when the modulus of waterglass was 0, 0.5, 1, and 1.5, respectively. Previously, two hypothetical mechanisms were proposed to explain the strength loss. One related the strength loss to a lower rate of CaCO3 crystallization during carbonation process to fill the pore structure. Availability of trace amounts of CH and ettringite as sources of Ca2 + for CaCO3 formation was regarded as one main reason [5,28]. Besides, CaCO3 formation rate can be suppressed by low Ca/Si of C-S-H in hardened alkali-activated slag. The other mechanism stated that, in spite of a low Ca/Si of C-S-H gel with high durability, once the decalcification occurred the C-S-H gel damaged more quickly and cohesion of gel decreased [8]. As for the activator type and modulus of waterglass, results were not indicative of any relationship between the change in modulus of waterglass and strength loss of samples after carbonation, nonetheless, the samples activated with NaOH-only solution (WG-0) were less affected than that of waterglass-activated slag. The decrease in compressive strength of waterglass-activated slag was also reported in a study by Puertas [8] et al. However, they showed that the carbonation increased strength of samples activated with NaOH-only solution which was related to precipitation of greater amounts of calcium carbonate in the pores, decline in total porosity and average pore size; and consequently, an increase in compressive strength. In this study, although no increase was observed in the strength of NaOH-activated samples (WG-0) after carbonation, the strength loss was comparatively low. 3.2. Effect of carbonation on reaction products 3.2.1. XRD analyses Fig. 5 illustrates the XRD patterns of hydrated Portland cement and activated slag before and after carbonation. XRD patterns of uncarbonated Portland cement samples showed eminent peaks of Portlandite at 2θs of 18°, 28°, 34°, 47°, 51° and 54°. The characteristic peaks of alite and belite were also traced at 2θs from 29° to 44°. As for the alkali-activated slag samples, a wide dispersion ring (hump) was observed around 2θ of 30° indicating the C-S-H gel in an amorphous form as a dominant hydration product of activated slag as well as presence of some unreacted slag. In addition, a peak associated to hydrotalcite (at 2θ of 12°) was also traced for NaOH-activated slag. The formation of double layered hydroxides of the hydrotalcite appeared to act as an internal CO2 sorbent, which could reduce the susceptibility of AAS to carbonation [21]. Fig. 5(b) illustrates the XRD patterns of PC and AAS after carbonation. As can be seen, carbonation affected PC and AAS in different fashions. XRD pattern of PC is illustrative of disappearance of alite and belite peaks and crystal transformation of CH to CaCO3. Carbonation of PC generated CaCO3 with a hexagonal system in the form of calcite and a orthorhombic system in the form of aragonite, albeit, calcite was observed in a greater proportion. As for AAS samples, the calcium =portlandite, =alite =belite =hydrotalcite

a)

b) PC

PC

WG-0

WG-0

WG-1.0

WG-1.0

10

=calcite =aragonite =vaterite

20

30

40

2 (degree)

50

60

10

20

30

40

2 (degree) 217

50

60

Fig. 5. XRD patterns of PC and AAS before (a) and after carbonation (b).

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1639

875

3643

970

1483

454

3416

1401

3131

713 855 451

1630

1082

3439

875

3134 Uncarbonated Carbonated 4000

3500

875 cm− 1. It should be noted that bands at 450 cm− 1 are associated with the SieO band in SiO2. The first broad and large absorption band was traced at 959 cm− 1 in AAS samples activated with NaOH-only solution (WG-0) which can be related to C-S-H [15]. This peak appeared at 970 cm− 1 in PC samples that signifies the dissimilar C-S-H formed in cement hydration with different Ca/Si. The alteration of Ca/Si was also observed when waterglass was used in which the increase in the waterglass modulus shifted the band from 959 to 968 cm− 1. The higher compressive strength of samples activated with NaOH/waterglass can be well explained by the decrease in the Ca/Si. However, it should be noted that an excessive increase in the waterglass modulus may cause decrease in liquid alkalinity; the remnant monomer [SiO4] cannot participate in the reaction, and hence the reaction extent would be incomplete [45]. The data show that the band associated with C-S-H underwent a distinctive change after carbonation. For carbonated samples, the bands in NaOH-activated samples (WG-0) moved to 1026 cm− 1, as calcium in the C-S-H gel reacted with CO2 and the calcium content in the C-S-H decreased to produce silica gel [5,46]. The increase of waterglass modulus shifted the peaks from 1031 to 1018 cm− 1. The comparison of Figs. 6 and 7 shows that the C-S-H gel structures of AAS with different activators were different from that of PC. Conclusively, the decalcification of C-S-H was the main source of carbonation in AAS samples. The stretching vibrational peaks of O-C-O in carbonate ([CO3]2 −) [34] were traced at a hump peak around 1436–1441 cm− 1. This can be attributed to the overlapped stretching vibrational bands of calcite at 1430 cm− 1 and vaterite at 1490 cm− 1 and 1420 cm− 1 as reported by Sato et al. [47]. Other growing peaks were also detected at 875 and 713 cm− 1 corresponding to the bending vibration peak of OeCeO in carbonate [CO3]2 − [34]. The band at 713 is the characteristic peak of calcite while 875 cm− 1 represents the overlapped peaks of vaterite and calcite. After carbonation, bands at 875 cm− 1 and 713 cm− 1 became wide and obvious, showing formation of more carbonated products. Vaterite has reportedly a distinct in-plane bending band at 750 cm− 1 [47], however, FTIR technique was not able to detect the band. This may indicate that the in-plane band of [CO3]2 − in carbonated AAS was of a low content. Fig. 8 illustrates the change of main bands in AAS and PC over the range of 800–1300 cm− 1, which can show information on four types of vibration of SieO (Q1, Q2, Q3 and Q4) in C-S-H gel [40,42]. The band at around 810 cm− 1 is due to SieO stretching vibration at Q1 site as the main element in unhydrated cement paste [39,48]. As the hydration was relatively sufficient, the degree of polymerization of silicate increased [49], so the peak in this study shifted to 840 cm− 1. The main band at around 970 cm− 1 is assigned to SieO stretching vibrations at Q2 site [39,48]. The Q2 tetrahedron indicates a wide range of C-S-H (Ca/Si ≈ 2:1) and is the most common tetrahedron in well hydrated cement paste. This band showed the highest peak in PC samples as the Ca/Si is the highest in C-S-H from hydration of cement [39,48]. On the contrary, the Q2 site decreased and shifted to 959–968 cm− 1 in AAS showing the different density of C-S-H. Carbonation changed the Q2 site of samples in different fashions. Decalcification of C-S-H caused a different Ca/Si, so the Q2 site shifted to higher wavenumber after carbonation. The two other bands at 1089 cm− 1 and 1135 cm− 1 are typical bands of Q3 and Q4 sites assigned to C-S-H with higher SiO2 content [50,51]. The Q3 site shows carbonated C-S-H, while content of Q4 site represents incomplete hydration. Q4 is relatively low as bridging oxygens in SieO bands are mostly replaced by Ca in well-hydrated cement paste. The shift of Q3 to larger wavenumbers (1089 cm− 1) may indicate the transformation of C-S-H to silica gel which was also reported by Rostami et al. [52]. In addition, the band at 1036 cm− 1 is corresponding to high concentrations of sodium bound onto silicate-based gel [50]. It can be seen from Fig. 8 that the C-S-H formed in PC and AAS are dissimilar. The peak of main site (Q2) in C-S-H gel of un‑carbonated PC

3000

1402 2500

2000

1500

1000

500

-1

Wavenumber (cm ) Fig. 6. Fourier-transform infrared spectra of PC before and after carbonation.

3.2.2. FTIR analyses To further elaborate the phase change in carbonated PC and AAS samples, FTIR technique was implemented with a focus on bands of OeH, HeOeH, SieOeT and CeO, corresponding to CH, chemically combined water, C-S-H, and CaCO3, respectively. Fig. 6 shows the FTIR spectra of PC before and after carbonation in the range of 400–4000 cm− 1. The peaks of key interest for PC samples were bands related to CH and C-S-H which are theoretically affected by carbonation. As can be seen, there is an absorption peak at about 3643 cm− 1 that represents the stretching vibration of OeH in Portlandite [5]; which disappeared after the carbonation indicating consumption of CH during carbonation. The other peak associated with CH and chemically combined water in hydration product is the bending vibration peak of H-OH at about 1639 cm− 1 [39] that transformed to a smaller peak shifted to 1630 cm− 1. The peaks related to C-S-H were also affected by carbonation, accordingly. The band around 970 cm− 1 represents the stretching vibration of SieOeT in C-S-H, in which T represents tetrahedral silicon or aluminum [40]. After carbonation, the peak shifted from 970 cm− 1 to 1082 cm− 1, indicating decalcification of C-S-H and formation of highly polymerized and elongated form of amorphous silicate gel [41,42]. Consequently, the changes and transformation of bands of CH and C-S-H altered peaks related to CaCO3. As illustrated, the absorption peaks of 1483 cm− 1 and 875 cm− 1 became wider and higher and new peaks appeared at 855 cm− 1 and 713 cm− 1 after carbonation. Reportedly, bands at about 1483 cm− 1, 875 cm− 1, 855 cm− 1, 713 cm− 1 are typical characteristics of [CO3]2 − [34]. This specifies that the amount of [CO3]2 − increased in carbonated samples. Bands at 1483 cm− 1 represent [CO3]2 − stretching vibration [34], whereas 875 cm− 1, 855 cm− 1, 713 cm− 1 are caused by [CO3]2 − bending vibration [34]. According to Hidalgo et al. [43], the corresponding absorption peaks of calcite from carbonated cementitious materials are 1410 cm− 1, 874 cm− 1, and 710 cm− 1. The difference between calcite and aragonite lies in the peaks at 875 cm− 1 and 855 cm− 1.The band at 875 cm− 1 is the characteristic peak of calcite while 855 cm− 1 represents aragonite [34,44]. The results showed that the major products of carbonation reaction were calcite and aragonite. The coexistence of calcite and aragonite was also detected in XRD. Fig. 7 shows the FTIR spectra of the AAS samples before and after carbonation. The major carbonation induced changes were traced at 500–2000 cm− 1. No absorption peak was observed at the vicinity of 3643 cm− 1, which signifies formation of no Ca(OH)2 (the illustration of spectra is limited to the active wavenumbers). However, the bending vibration peak of the H-OH bond was observed at 1641 cm− 1 which can be related to chemically combined water in the hydration products formed after activation [15]. This peak can't be linked to CH as XRD did not trace any CH in the AAS matrix. The major bonds targeted in the carbonation of AAS matrix were 959–968 cm-1, 1018–1031 cm− 1 and 218

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Fig. 7. Fourier-transform infrared spectra of AAS before (a) and after carbonation (b).

WG-0

WG-0

713

1638

1641 1441 WG-0.5

877

WG-0.5

1026 875

1441

451 713

959 1638

876

1641 1443

444

447

1031 875

449 1440 WG-1.0

962

WG1.0

713

1638

1024 875

451

1439

1641 1441 WG-1.5

451

WG-1.5

963

1641 1443

a) 2000

877

876

713 1638

452

b)

968 1500

1000

Wavenumber (cm-1)

1018

500

2000

875

448

1436

1500

1000

500

Wavenumber (cm-1)

was also showed by increase of Q3 in Fig. 8. The results are also consistent with the experimental findings of Song et al. [5]. Fig. 10 shows the DTG curves of the AAS samples before and after carbonation. From Fig. 10(a), it can be seen that the characteristic peaks of Ca(OH)2 were not found in the un‑carbonated AAS samples. The results are in agreement with the data from the XRD and FTIR indicating formation of no Ca(OH)2 in the AAS samples. The major mass loss mainly occurred at the range of 100–220 °C due to evaporation of molecular water in C-S-H and pores [5,19]. At 170 °C, there was an obvious endothermic peak in NaOH-activated slag related to dehydration of hydrotalcite [19,55]. This phenomenon also resulted in a mass loss at 350–400 °C [19]. All the un‑carbonated AAS samples showed obvious CaCO3 endothermic decomposition peaks at 600–680 °C, albeit, it was not detected in XRD. However, the OeCeO bands were traced in FTIR showing that certain carbonation reaction happened during early curing. This may correspond to amorphous calcium carbonate, which was also characterterized previously in [41]. Furthermore, it was observed that CaCO3 content in NaOH-activated slag was of minimum amount which can be related to the available hydrotalcite in the matrix. Previously, Bernal et al. [21] reported that the higher content of MgO in slag enhanced the carbonation resistance of mortar due to formation of hydrotalcite in the matrix. The modulus of waterglass was also an influential factor in the early carbonation of samples after exposure to CO2 in the atmosphere. As illustrated, WG-1.0 slag had the lowest content of CaCO3 (0.59%) compared to that of WG-0.5 (1.25%) and WG-1.5 (1.12%). It can be mainly attributed to low porosity and well-distributed pore size in samples resulting in high strength and resistance to carbonation. After carbonation, noticeable changes occurred in DTG curves of AAS samples. The initial peak (at 20–220 °C) showed an extensive reduction in size, which can be indicative of decalcification of C-S-H in the matrix. The other major change was related to formation of hydrotalcite at 177 °C in all AAS samples. Comparing to un‑carbonated samples it can be concluded that the presence of waterglass delayed the formation of hydrotalcite in the matrix. It may be explained by lower alkalinity of the liquid phase of waterglass activated slag than NaOHonly solution, even though Na2O was fixed at 4% mass of slag. As mentioned earlier, formation of hydrotalcite may hinder crystallization of CaCO3 and decrease its enhancing effect on strength as observed in Fig. 4. The other change induced by carbonation was partial disappearance

sample was the dominating peak. After exposure to full carbonation, peaks corresponding to Q2 decreased, while Q3 and Q4 increased. This signifies the decalcification of C-S-H gel to produce CaCO3 and silica gel in PC samples after carbonation. It was noted that the total gel content was basically unchanged. In AAS samples, however, the peaks of Q2 and Q3 reduced after carbonation largely, which caused the total gel content to decrease at different extent, especially, in WG-1.0. The data show that the gel content loss was the least in NaOH-activated slag comparing to WG-0.5, WG-1.0, and WG-1.5. All these may show more vulnerability of the C-S-H in AAS samples to decalcification especially with presence of waterglass. The results can well explain the reduction of compressive strength of AAS after carbonation. 3.2.3. DTG analyses Fig. 9 shows the DTG curves of the PC before and after carbonation. Three endothermic peaks are detectable in cement matrix at 100–200 °C, 400–500 °C and 760–950 °C corresponding to dehydration of C-S-H, CH and thermal decomposition of CaCO3, respectively [53]. The first two peaks were traced at 20 to 220 °C centered at 143 °C and 410 to 480 °C centered at 438 °C showing 4.78% and 4.60% mass loss related to degradation of C-S-H and CH, respectively. The characteristic endothermic peak of CaCO3 was also observed at 668 °C with a low mass loss of 2.01%. The trace amount of CaCO3 may be caused by exposure of samples to CO2 during DTG testing process. However, accelerated carbonation caused significant increase in the CaCO3 associated peak (at 721 °C) to 13.37%, concomitant with reduction of peak of C-S-H to 2.77% and complete disappearance of CH peak. The DTG curve of PC showed that the Ca2 + from CH was the main source to generate CaCO3. However, the complete carbonation also decalcified CS-H gel resulting in formation of CaCO3 and silica gel. The results show that the generated CaCO3 decomposed at lower temperatures (721 °C), which may be attributed to the presence of calcite as the main polymorph of CaCO3 from both carbonation of CH and tansformation of aragonite. Thiery et al. [53] showed that the CaCO3 ensuing from carbonation of CH is in calcite form as it decomposes at higher temperatures of 760–950 °C. The calcite can also be from transformation of aragonite at about 460 °C [54], which decomposes within the temperature range of 680–780 °C due its less perfect crystalline state and high dispersity within the matrix [41]. Furthermore, the observed mass loss at 500–650 °C indicates the loss of chemically bound water from the silica gel. The formation of silica gel from decalcification of C-S-H 219

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Fig. 8. Change of the SieO bands of PC and AAS before and after carbonation.

0.30

PC

Uncarbonated Carbonated

875

0.25 970

0.20

1036 1089 1135

0.15 855

0.10 0.05

840

0.00 800

900

1000

1100

1200

1300

Wavenumber (cm-1)

0.30

962

WG-0

Uncarbonated Carbonated 1036

0.40

Uncarbonated Carbonated

959

WG-0.5

0.35

0.25

1036

0.30

1089

0.20

1135

875

1089 1135

0.25 0.20

0.15

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875 840

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-1

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WG-1.0

Uncarbonated Carbonated

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968

WG-1.5

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1300

Uncarbonated Carbonated

1036 1089

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1135

0.20 875

0.15

875

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840

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Wavenumber (cm )

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˚C

˚C ˚C

2.77%

˚C

˚C

13.37% ˚C

Uncarbonated Carbonated

200

1100

1200

1300

of the peak traced at 250–400 °C in un‑carbonated samples. Some studies related this peak to the dehydration of C-S-H gel and other calcium aluminate hydrate products [19,56,57]. Therefore, the disappearance of the associated peak may also show the decalcification of C-S-H. The most important difference between the DTG curves was decomposition of the carbonate-rich phase (300–850 °C). The characteristic peak at 690 °C showed decomposition of CaCO3 which occurred at a lower temperature comparing to PC samples (721 °C) referring to the coexistence of calcite and vaterite as major polymorphs of CaCO3. According to Sauman et al. [54], carbonation of C-S-H may lead to formation of vaterite, which decomposes at lower temperatures between 500 and 700 °C as compared to calcite due to its imperfectly crystallized formation or finer crystal structure. In this study, the formation of vaterite was detected in AAS samples by observing patterns corresponding to endothermic decomposition peaks at 500–600 °C, centered at 590 °C [58]. The coexistence of calcite and vaterite was also observed in the results of XRD (Fig. 5). The DTG results also showed that the mass loss related to CaCO3 was 4.64%, 3.17%, 4.44%, and 4.16% in NaOH-activated samples, WG-0.5, WG-1.0 and WG-1.5, respectively. This may show lower rate of calcite formation in AAS samples comparing to that in PC. It can explain the lower compressive strength of slag activated samples after carbonation.

2.01%

4.60%

4.78%

900

Wavenumber (cm-1)

-1

˚C

400

600

800

˚C Fig. 9. Differential thermograms of PC before and after carbonation.

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Fig. 10. Differential thermograms of AAS before (a) and after carbonation (b).

˚C WG-0 WG-0

0.12%

4.43%

4.64%

5.43% WG-0.5

˚C

WG-0.5

3.15%

3.17%

1.25% 5.30%

WG-1.0 0.59%

WG-1.0

2.26% 4.44% WG-1.5

6.69%

1.83% WG-1.5

200

a)

˚C

5.48%

400

600

4.16%

˚C

1.12%

800

b)

˚C

200

400

˚C

600

800

˚C indication of removal of chemically bound water in 1H test in order to obtain an accurate porosity. The data show that a larger signal with greater peak area was observed for PC samples after carbonation, whereas the peaks weakened significantly in AAS samples. The mechanism below can be used to explain the finding. With formation of CaCO3 in the pore structure of PC matrix, closed pores were formed which hindered the drying of moisture in the capillaries by which larger drying signals were recorded for PC samples. “Pore blocking” effect also decreased the porosity of PC samples from 14.70% to 9.91% after carbonation. On the contrary, the carbonation increased the porosity of AAS samples by 28.46%, 41.04%, 50.79%, and 60.20% when modulus of waterglass set to 0, 0.5, 1.0, and 1.5, respectively. A remarkable reduction in the dry sample signals was also observed which can again be related to a lower formation rate of CaCO3 crystal. As can be seen, the increase in the modulus of waterglass decreased the signal peak area and lowered the “pore blocking” effect of the carbonation process. The results confirmed the higher formation rate of CaCO3 in NaOHactivated AAS than that of samples with waterglass as was shown in the chemical phase analysis. The higher strength of PC samples and strength decrease of AAS mortars can also be well explained by the effect of carbonation on pore structure of samples.

3.3. Effect of carbonation on pore structure 3.3.1. Porosity As the formation of carbonation products is an expansive reaction, porosity analysis was carried out by 1H NMR measurement on water saturated concrete. A built-in computing method was used by which the sampling data were inversed as T2 (surface transverse relaxation of 1H protons in porous media) distribution. In this test, volume change, dry sample signal, and porosity were targeted to investigate the changes induced by carbonation in pore structure of AAS and PC samples. Table 3 illustrates the obtained results from the samples before and after carbonation. As can be seen, major differences were observed in volume, dry sample signal, and porosity of PC and AAS before and after carbonation. The results are indicative of shrinking effect of carbonation process on all samples, albeit, the shrinking rate was lower in PC samples. This can be mainly related to the expansive transformation of Ca(OH)2 to CaCO3. It was reported that the formation of CaCO3 increased the volume by 11.8% [12,25]. On the other hand, the increase in the modulus of waterglass increased the shrinkage rate by 7.43%, 7.67%, 9.98% and 11.53% when modulus of 0, 0.5, 1, and 1.5 were used in AAS samples. This may be corresponding to hindered crystallization of CaCO3 by formation of C-S-H with lower Ca/Si as the waterglass modulus increased. Furthermore, formation of hydrotalcite during the carbonation process could also have an interrupting effect on CaCO3 crystallization. The other distinctive behavior of PC and AAS samples was in the dry sample signals. In this test, the dry sample signal was used as an

3.3.2. Pore structure Nuclear Magnetic Resonance Imaging (NMRI) of saturated water samples was carried out in order to visualize the pore structure of PC and AAS before and after carbonation. A two-dimensional imaging of the cross-section in two axial directions along the specimen was

Table 3 Porosity analysis of PC and AAS mortars before (a) and after carbonation (b). Sample

Volume/ml

Dry sample signal area

Dry porosity/%

Saturated peak area

Saturated porosity/%

Porosity (saturated-dry)/%

a) PC b) PC a) WG-0 b) WG-0 a) WG-0.5 b) WG-0.5 a) WG-1.0 b) WG-1.0 a) WG-1.5 b) WG-1.5

69.19 67.09 71.56 66.24 72.74 67.16 71.41 64.15 71.74 63.47

6229.12 6459.87 11,071.75 1495.31 14,476.41 1662.77 16,167.72 1691.17 17,323.88 1712.90

3.43 3.69 6.06 0.70 7.86 0.78 8.97 0.85 9.58 0.87

31,271.38 22,827.10 35,407.89 30,437.47 36,475.36 30,304.38 37,324.68 30,344.34 36,992.39 29,587.31

18.13 13.60 19.87 18.44 20.14 18.10 21.00 18.99 20.72 18.71

14.70 9.91 13.81 17.74 12.28 17.32 12.03 18.14 11.13 17.83

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a) PC

a) WG-0

a) WG-0.5

a) WG-1.0

a) WG-1.5

b) PC

b) WG-0

b) WG-0.5

b) WG-1.0

b) WG-1.5

Fig. 11. NMRI images of mortars before (a) and after carbonation (b).

The PC had the smallest span while activation of slag especially with higher modulus of waterglass (up to 1) shifted pore size distribution to micro pores with stronger signal intensities. After carbonation, the peak and the width of the first peak decreased, especially in the PC samples. On the contrary, the peak and width of the second peak increased, while little effect was observed on the third peak. The large decrease in signal intensity of micro pores from 720 to 380 showed decreased porosity in PC and densification of the matrix. This indicates that the CaCO3 particles precipitated mainly into the micro pores including gel pore and C-S-H interlayers. At the same time, the decalcification caused by carbonation destroyed the structure and increased the content of mesopores, but little effect on the macro-pores was detected in the matrix. The decrease in the signal intensity of PC micropores was higher than that of AAS. As modulus of waterglass increased the signal intensity of micropores declined. However, signal intensity of mesopores in waterglass activated slag was higher than that of slag activated with NaOH-only solution. The results can show the rate of calcium carbonate crystallization and degradation of C-S-H. Therefore, the AAS had a lower rate of calcium carbonate formation than PC. Besides, the C-S-H in waterglass activated slag seemed to be more vulnerable than that of samples activated with NaOH-only solution. The results are in agreement with compressive strength results, XRD, FTIR and DTG.

obtained, accordingly. In the image, the color spots are the semaphores of the water molecules. The warmer color of spots in the image specifies the pore size and the distribution rate. Using this feature, NMRI can directly illustrate the pore size distribution inside the matrix. Fig. 11 illustrates the variation in porosity of samples before and after exposure to CO2. Comparison of the images is illustrative of densification of PC sample by carbonation, despite its higher porosity before carbonation when compared with AAS samples. Higher porosity of AAS with modulus of 1.5 can also be observed in the images which is consistent with the results from compressive strength. This can be mainly related to incomplete reactions in matrix. As illustrated, carbonation increased the porosity of AAS samples activated with waterglass to a greater extent than that of samples activated with NaOH-only solution. A higher CaCO3 content was formed in NaOH-activated slag as observed in the results from DTG test. This can contribute to a more porous matrix as a result of continuous decalcification of C-S-H gel [8]. For a better elaboration of pore structure of samples, the gray value data were also obtained and illustrated in Fig. 12. The pore size and pore distribution were also calculated based on T2 spectrum of NMR. According to NMR principle [59,60], the distribution of T2 spectrum of NMR is relative to the size and distribution of the pores. The position of the peak is related to the pore size while its signal intensity represents the number of pores associated with that size. Generally, the 1 ms T2 relaxation time corresponds to a pore size of 24 nm in cement paste [61,62]. Fig. 13 shows the T2 spectra of different types of mortar samples before and after carbonation. As can be seen, the T2 spectrum of all samples were distributed over a range of 0.01 to 10,000 ms, indicative of 3 peaks corresponding to micro-, meso-, macro pores. The graphs show that the water in samples was distributed over a wide range of pore sizes. For PC samples, the signal intensity and area of the first peak in the range of 0.01–10 ms were the largest indicating that the water in the sample was mainly distributed in a continuous micro sized range. The signal intensity and area of the second (10100 ms) and third peaks (100–10,000 ms) were relatively small. It shows that the water in the sample was spread in meso and macro sized pores. However, a significant difference was observed between PC and AAS in the first peak which specifies the dissimilarity of two matrixes at micro level. The pore size distribution at larger scale was almost identical in AAS and PC samples. The span of the pore distribution in each matrix is due to the difference in the density of reaction products.

3.4. SEM observation Fig. 14 shows SEM pictures of PC and AAS mortar samples before and after carbonation. SEM pictures well elaborate differences between C-S-H formed from hydration of cement and activation of slag using NaOH-only solution and NaOH/waterglass. The different morphology of C-S-H is reportedly attributed to different Ca/Si of C-S-H gel. Chen et al. [63] reported 30 kinds of C-S-H gel based on Ca/Si. For instance, CSH (I) was an incomplete form of 1.4 nm tobermorite in which Ca/Si varied between 0.7 and 1.5; CSH (II) was an incomplete form of jennite with Ca/Si near to 2; CSH (I) and CSH (II) were amorphous [34]. Previously in a study by Puertas et al. [8], it was reported that the Ca/Si of C-S-H gel of waterglass-activated slag was ~0.8 and was close to 1.4 nm tobermorite of amorphous C-S-H gel. The Ca/Si of C-S-H gel of NaOH-activated slag was ~1.2, and Ca/Si was larger than that of waterglass-activated hardened matrix. 222

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Fig. 12. Gray value of NMRI images of mortars before (a) and after carbonation (b).

a) PC

b) PC

a) WG-0

b) WG-0

b) WG-0.5

a) WG-0.5

a) WG-1.0

a) WG-1.5

b) WG-1.0

b) WG-1.5

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Fig. 13. T2 spectrum of PC and AAS mortars before and after carbonation.

800

PC

Signal amplitude

700

Uncarbonated Carbonated

600 500 400 300 200 100 0 0.01

0.1

1

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100

1000 10000

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After carbonation, microstructure of PC and AAS mortar samples also showed different characteristics. As illustrated, CaCO3 was formed with white round small particles in PC samples, filled into the pores and densified the matrix; however, the decalcification of C-S-H gel in AAS mortar samples resulted in more dispersed porosity in the matrix. The CaCO3 crystals were rather with ruled and oriented growth. Obvious hexagonal crystals of almost 0.5 μm were formed as illustrated in the image. Comparing SEM pictures of PC and AAS after carbonation, it is apparent that there are differences in the size and morphology of CaCO3 particles formed by carbonation AAS and PC.

•

4. Conclusion

•

The microstructural changes of alkali-activated slag with different modulus of waterglass after carbonation were investigated and compared with that of Portland cement mortar. The results showed dissimilar effects of carbonation on microstructure of samples.

• Carbonation had an adverse effect on compressive strength of alkali-

•

100

800

WG-1.0

700

0 0.01

10

T2 (ms)

T2 (ms)

activated slag mortars. However, the samples activated with NaOHonly solution were more resistant to carbonation in terms of residual strength. The strength loss of samples activated with NaOH/waterglass (≈ 26%) was twice more than sample activated with NaOHonly solution (13.6%). On the contrary, the compressive strength of PC mortar samples increased after carbonation by up to 14%. The increase in the modulus of waterglass shifted the C-S-H to a phase with lower Ca/Si, which resulted in a higher compressive

strength of samples comparing to that of samples activated with NaOH-only solution. With presence of no CH, the decalcification of C-S-H was the main carbonation reaction in AAS samples. The change of the main bands of C-S-H showed that the Q2 reduced after exposure to carbonation, which indicated the disintegration of the C-S-H gel. More vulnerability of C-S-H to decalcification in AAS samples especially with presence of waterglass was observed. Carbonation of PC samples generated CaCO3 in the form of calcite and aragonite. On the other hand, the calcium carbonate formed in AAS was mainly in the form of calcite and vaterite. The reason was the low Ca/Si in the AAS binder as well as formation of hydrotalcite, which consumed the Mg in the pore solution to a certain amount and hindered crystallization of aragonite. A lower rate of calcite formation was detected in AAS samples. CaCO3 particles precipitated mainly in the micropores including gel pore and spaces in C-S-H interlayers. However, little effect was observed on the macropores, which caused a certain volume contraction. Restricted crystallization of CaCO3 led to decrease the compressive strength in carbonated AAS samples which was more underscored in samples activated with a combination of NaOH and waterglass.

Acknowledgement This research is financially supported by the National Science Foundation of China under project numbers of 51638008 and 51461135001. 224

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Fig. 14. SEM pictures of PC and AAS mortars before and after carbonation.

(a) Pictures of PC before (a) and after carbonation (b).

(b) Picture of WG-0 before (c) and after carbonation (d).

(c) Picture of WG-1.0 before (e) and after carbonation (f).

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