# INNOVATION IN CEMENTS FOR SUSTAINABILITY

EDITED BY : John L. Provis, Maria Juenger and Miroslav Komljenovic PUBLISHED IN : Frontiers in Materials

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ISSN 1664-8714 ISBN 978-2-88963-346-3 DOI 10.3389/978-2-88963-346-3

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# INNOVATION IN CEMENTS FOR SUSTAINABILITY

Topic Editors:

John L. Provis, University of Sheffield, United Kingdom Maria Juenger, University of Texas at Austin, United States Miroslav Komljenovic, University of Belgrade, Serbia

Citation: Provis, J. L., Juenger, M., Komljenovic, M., eds. (2020). Innovation in Cements for Sustainability. Lausanne: Frontiers Media SA. doi: 10.3389/978-2-88963-346-3

# Table of Contents


*Residue-Slag-Cement Binding Material System* Rongjie Song, Qingxin Zhao, Jinrui Zhang and Jizhong Liu

## Editorial: Innovation in Cements for Sustainability

#### John L. Provis <sup>1</sup> \*, Miroslav Komljenovic<sup>2</sup> and Maria C. G. Juenger <sup>3</sup>

*<sup>1</sup> Department of Materials Science and Engineering, University of Sheffield, Sheffield, United Kingdom, <sup>2</sup> Department of Materials Science, Institute for Multidisciplinary Research, University of Belgrade, Belgrade, Serbia, <sup>3</sup> Department of Civil, Architectural and Environmental Engineering, University of Texas at Austin, Austin, TX, United States*

Keywords: cement, sustainability, construction, concrete, alkali-activated (AA) cement

#### **Editorial on the Research Topic**

### **Innovation in Cements for Sustainability**

Sustainability is trying to balance economic growth, environmental protection and social progress with future technologies, while supporting innovation, and not compromising the way of life for present and future generations. The idea of sustainability should bring people together to foster the biosphere for the benefit of all living beings on our planet Earth.

It is with great pleasure that we present this Research Topic on Innovation in Cements for Sustainability. The collection of 10 articles in this Research Topic represent some very exciting advances in how we design, use, and analyse cementitious materials, and we expect that this work will lead to further advances along the pathway toward sustainability for the construction materials sector.

The work presented here related to Portland cement includes a pioneering study of tricalcium silicate hydration under microgravity conditions (Moraes Neves et al.), and important work on reducing Portland cement production temperatures using increased iron content in the clinker (Elakneswaran et al.). Portland-blended cements are highlighted through work on improving metakaolin dispersion in mortars (Geng et al.), and on the use of nano-alumina to improve the strength and microstructure of high-volume blended cements (Shaikh and Hosan). Hybrid alkaline-belite cements, which involve aspects of both conventional cement and alkali-activation chemistry, are also presented (Sánchez-Herrero et al.).

The development of innovative approaches to the activation of aluminosilicates to produce non-Portland cements is also addressed, using phosphate activators (Wang et al.) and also alkaline waste soda residues (Song et al.) as pathways to produce cements with novel characteristics. The combination of mechanical and chemical activation of fly ash was demonstrated as a way to unlock the potential to produce high-performing binders from ashes of various quality levels (Fernández-Jiménez et al.). The use of new precursors to produce alkali-activated binders has also shown exciting results in materials designed for key applications: heat-resistant porous cementitious materials produced from a combination of wastes from different industries (Bajare et al.), and also innovative mixed-valence ferrosilicate/ferrisilicate binders from synthetic iron-rich slags (Peys et al.).

The Guest Editorial team is confident that this collection of papers will be of benefit to the broader community in research and application of cementitious materials.

### Edited and reviewed by:

*Zuhua Zhang, University of Southern Queensland, Australia*

#### \*Correspondence: *John L. Provis j.provis@sheffield.ac.uk*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *24 October 2019* Accepted: *06 November 2019* Published: *19 November 2019*

#### Citation:

*Provis JL, Komljenovic M and Juenger MCG (2019) Editorial: Innovation in Cements for Sustainability. Front. Mater. 6:298. doi: 10.3389/fmats.2019.00298*

**4**

## AUTHOR CONTRIBUTIONS

All authors participated in drafting and approving this Editorial.

**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Provis, Komljenovic and Juenger. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

## Mechanical-Chemical Activation of Coal Fly Ashes: An Effective Way for Recycling and Make Cementitious Materials

### Ana Fernández-Jiménez, Ines Garcia-Lodeiro\*, Olga Maltseva and Angel Palomo

Department of Materials, Eduardo Torroja Institute, Consejo Superior de Investigaciones Cientificas (CSIC), Madrid, Spain

This paper is wholly committed to resource efficiency through the valorization of waste from other industries, and more specifically fly ash as a raw material to produce concrete-like geopolymers. In particular, this study aimed to determine the effect of the physical and chemical characteristics of recycled coal fly ash used to manufacture alkaline cement on reaction kinetics and product microstructure and performance. The ash was mixed with 8 M NaOH and cured for 20 h at 85◦C and RH > 90% to form a compact paste, after which mechanical strength was determined and the reaction rate was calculated using isothermal conduction calorimetry. The findings showed that vitreous content (SiO2/Al2O3), reactive and (especially) fineness play a very important role in both the development of cement mechanical strength and the composition and structure of the reaction products formed.

## Edited by:

Maria Juenger, University of Texas at Austin, United States

#### Reviewed by:

Prinya Chindaprasirt, Khon Kaen University, Thailand Jae Eun Oh, Ulsan National Institute of Science and Technology, South Korea

\*Correspondence:

Ines Garcia-Lodeiro iglodeiro@ietcc.csic.es

#### Specialty section:

This article was submitted to Structural Materials, a section of the journal Frontiers in Materials

Received: 08 January 2019 Accepted: 11 March 2019 Published: 04 April 2019

#### Citation:

Fernández-Jiménez A, Garcia-Lodeiro I, Maltseva O and Palomo A (2019) Mechanical-Chemical Activation of Coal Fly Ashes: An Effective Way for Recycling and Make Cementitious Materials. Front. Mater. 6:51. doi: 10.3389/fmats.2019.00051 Keywords: recycling, fly ash, alkaline activation, geopolymer, mechanical-chemical activation, NMR

## INTRODUCTION

Portland cement manufacture consumes around 4 GJ of energy per ton and accounts for approximately 8% of worldwide CO<sup>2</sup> emissions, or 2.981 GtCO2eq in 2014 (CSI/ECRA—Cement Sustainability Initiative/ European Cement Research., 2017).

The Roadmap for a resource-efficient Europe (COM, 2011) 0571 describes industrial strategies to mitigate the effects of industrial activity and ensure that by 2050 Europe will have a sustainable economy, in which "resources are not simply extracted, used and thrown away, but are put back in the loop so they can stay in use for longer."

Cross-cutting concerns such as the need for more long-term, innovative thinking are also being addressed. The research supporting this paper is wholly committed to resource efficiency through the valorization of waste from other industries, and more specifically, fly ash as a raw material to produce concrete-like "geopolymers," which contain a drastically reduced proportion of portland cement.

Worldwide production of combustion products (CCPs), including fly ash, amounted to 780 Mt in 2011 and nearly 1 Gt in 2015. Whilst the rate of re-use varies widely from country to country (10.6–96.4%), overall 47% of the fly ash generated is presently landfilled (Thomas, 2007). Options for reusing CCPs, which have been under study around the world for the last 70 years, can be classified as non-beneficial, simply transformed manufactures (STM) or elaborately transformed manufactures (ETM). Like other industrial waste such as blast furnace slag or red mud, fly ash can be activated in alkaline media, setting and hardening to acquire properties similar to those characteristics of portland cement. Such materials, known as "alkali-activated cements" or "geopolymers," constitute optimal vehicles for valorizing and recycling coal fly ash (Palomo et al., 2014; Zhuang et al., 2016; Bai and Colombo, 2018).

The 500 Mt/year of coal fly ash presently landfilled could be reused to manufacture cement or concrete (Chindaprasirt et al., 2005; Duvallet et al., 2015; Hemalatha and Ramaswamy, 2017) if ash particle reactivity could be significantly enhanced. Such valorization would also help mitigate global warming by reducing the CO<sup>2</sup> emitted by the building materials industry. The primary objective of the present study was consequently to devise processing methods that would heighten fly ash reactivity.

The main reaction product formed in the alkali activation of fly ash is an amorphous, three-dimensional alkaline inorganic polymer generically known as N-A-S-H (Na2O-Al2O3-SiO2- H2O) gel (Fernández-Jiménez et al., 2005; Provis and van Deventer, 2014). Secondary reaction products may include zeolites such as Na-chabazite, zeolites A and P, and faujasite (Fernández-Jiménez et al., 2005; Panias et al., 2007).

Earlier studies have shown an intrinsic relationship between the composition and structure of N-A-S-H gel and its physical and mechanical properties (Fernández-Jiménez et al., 2006; De Silva et al., 2007). Given similar degrees of reaction, the formation of N-A-S-H gels with an Si/Al ratio of ≈1 (type 1 gels) yields materials with lower mechanical strength than when the gel formed has an Si/Al ratio of ≈2 (type 2 gels) (Fernández-Jiménez et al., 2006; Hajimohammadi et al., 2010; Nikolic et al., 2015). Further studies by a number of authors (Sagoe-Crentsil et al., 2005; Kovalchuk et al., 2007; Fernández-Jiménez et al., 2017) optimal Si/Al ranges from 2 to 4. Such findings attest to the importance of the reactive silica and alumina content in the starting materials, which largely conditions the ultimate composition of the gel formed. In high calcium based systems (such as type C fly ashes), the initial setting process is governed by the formation of C-S-H or C-A-S-H-like phases rather than pure N-A-S-H. Fast dissolution of highly active Al2O<sup>3</sup> and SiO<sup>2</sup> sources in high pH medium provides high initial concentrations of silicate species and aluminate to react with Ca forming C-A-S-H phase in the early reaction stages (Chindaprasirt et al., 2012). Other works developed in the field of hybrid alkaline cements (systems with small OPC contents + high content of an aluminosilicate source) point out the co-precipitation of both C-A-S-H+ (N,C)-A-S-H gels during the early reaction stages, which would evolve with time to a C-A-S-H gel (Garcia-Lodeiro et al., 2011, 2013).

In addition to the chemical composition of the ash, the vitreous phase content and the (SiO2/Al2O3)reactive ratio, particle size is a factor to be borne in mind when determining a material's reactive potential. A number of studies mention the effect of fly ash particle size when used either as a supplementary cementitious material to lower the clinker factor (Paya et al., 1997; Chindaprasirt et al., 2005; Hemalatha and Ramaswamy, 2017) or as a precursor in the preparation of alkaline (OPC-less) cements (Fernández-Jiménez et al., 2005; Marjanovic et al., 2014; Duvallet et al., 2015).

Many authors (Rakesh et al., 2007; Temuujin et al., 2009; Kumar and Kumar, 2011; Kumar et al., 2015; Shekhovtsova et al., 2018) have broached the effect of fineness on ash reactivity, separating the material either by sieving (particles under 45µm (Duvallet et al., 2015) or electrostatic precipitation (Kumar et al., 2015). Their findings have consistently shown that mechanical strength rises substantially with ash fractions of under 45µm. Particle size separation normally prompts an increase in the proportion of vitreous phase, for the particles eliminated (the largest) consist essentially in scantly reactive crystalline phases such as quartz, unburnt carbon, and hematite. The discards generated by such processes would, naturally, need to be managed.

Kumar et al. have published a host of papers on the mechanical activation of fly ash (Kumar and Kumar, 2011; Kumar et al., 2015). They observed that milling fly ash raised the mechanical strength of the resulting cements, which they attributed both to the stimulation of ash reactivity by the smaller particle size (particularly at early ages) and to physical changes in porosity and the pore size distribution of the end product. Temuujin et al. (2009), however, observed no clear correlation between the rise in strength and changes in material porosity and density. The general consensus in the literature to date (Kumar and Kumar, 2011; Kumar et al., 2015) is that the mechanical activation of fly ash enhances the mechanical strength of the end product because the changes in particle morphology along with size hasten dissolution-precipitation reactions, raising reactivity. That in turn yields more compact microstructures, with a larger amount of gel and smaller fraction of unreacted particle (Zhang et al., 2016).

Most papers published on the effect of the mechanical activation of ash focus on a single material, with very few comparing the effect in ash from different sources (Marjanovic et al., 2014; Kumar et al., 2015). In a recent article proposing an equation to assess the suitability of fly ash for the production of a high-strength geopolymer, Zhang et al. (2016) reported that physical properties (density, specific surface, and fineness) may have a greater effect on ash reactivity than its vitreous phase content.

In light of the foregoing, the present study aimed to determine the effect of the mechanical activation of ash (by milling) on strength development, reaction kinetics and micro- and nanostructure of the reaction products.

The materials used had different vitreous phase contents and (SiO2/Al2O3)reactive fractions. Their initial particle size also differed, although they were ground for a sufficient time to substantially modify the specific surface of the original ash. The three ashes were ground and alkali-activated, after which their mechanical strength was determined and the reaction products were characterized by using XRD, SEM/EDX, and NMR.

## EXPERIMENTAL

## Materials

The three fly ashes used in this study, sourced from different Chinese steam plants, had different chemical and mineralogical compositions. Their chemical composition (% of oxides by mass), found with a PHILIPS PW 2400 X-ray fluorescence spectrometer fitted with a PW 2540 VTC sample changer, is given in **Table 1**.


The table also shows the XRD-detected minority crystalline phases in each ash. The significant differences in aluminum content prompted a fairly wide-ranging initial SiO2/Al2O<sup>3</sup> ratio, from 3.10 to 0.86.

## Experimental

Particle size distribution was determined on a COULTER LS 130 laser sizer with a measuring range of 0.1–900µm. The ash was ground in a laboratory ball grinder for 4 h (regarded as sufficient time to ensure similar particle size distributions in all three varieties of ash), using the same number of balls and loading the grinder with the same weight of sample throughout.

The materials were exposed to an acid attack (continuously stirring 1 g of ash in 100.00 mL of a 1% HF solution for 5 h (Ruiz-Santaquiteria et al., 2013) to quantify the potential reactivity of both the unground original (F1, F2, and F3) and ground (GF1, GF2 and GF3) ash. The resulting solution was then filtered, the residue was rinsed with distilled water to a neutral pH and the amount of dissolved (=potentially reactive) silica and alumina in the leachate was quantified by ICP on a Varian 725-ES optical ICP atomic emission spectrometer with the following settings: power, 1.40 kW; plasma gas flow, 15.00 L/min; nebulizer gas flow, 0.85 L/min; read time, 5 s. The findings were used to calculate the (SiO2/Al2O3)reactive ratio for each ash.

The ground and unground ashes were mixed with an 8 M NaOH solution and the resulting pastes molded into prismatic specimens measuring 1 × 1 × 6 cm<sup>3</sup> to determine whether these materials were susceptible to alkali activation and able to yield a product with good mechanical properties. The alkaline solution/ash ratio varied depending on the ash (see **Table 2**), yet was the same for the ground and unground versions of each. The samples were cured for 20 h in an oven at 85◦C and a relative humidity of ≥90%. These conditions usually guarantee a proper TABLE 2 | Mechanical strength of alkali-activated paste prepared with ground or unground coal fly ash and cured for 20 h at 85◦C.


FA reactivity (Fernández-Jiménez and Palomo, 2007; Zhou et al., 2016). The compressive strength of the specimens was found on an Ibertest Autotest 200/10-SW frame 24 h after mixing. Ten specimens of each composition were tested.

The starting ashes and reaction products were characterized by XRD, scanning electron microscopy (SEM/EDX) and NMR. X-ray diffraction (XRD) was conducted on a Bruker AXS D8 ADVANCE diffractometer. The scanning spectrum ranged from 22 angles 5 to 60◦ with a nominal step size of 0.0119736◦ at a rate of 0.5 s/step. Powder samples were exposed to Cu-Kα1, α2 X-ray radiation.

Fly ash morphology and cured specimen microstructure were examined under a JEOL 5400 scanning electron microscope fitted with an OXFORD LINK-ISIS EDX (energy dispersive X-ray) system. EDS analyses (on spots) were done with accelerating voltage of 20 kV, working distance of 15 mm and beam current of 20 µA and acquiring for 60 s per spot analysis. An average of 50 analys were taken per sample.

The NMR scans were recorded on a BRUKER AVANCE-400 spectrometer. All the samples were spun at 10 kHz. For <sup>27</sup>Al MAS-NMR, the settings were: frequency, 104.3 MHz; 2 ms single pulse MAS; 5 s relaxation delay; external standard, Al(H2O)6. For <sup>23</sup>Na MAS-NMR they were: frequency, 105.8 MHz; 2 ms single pulse MAS; 5 s relaxation delay; external standard, NaCl. For <sup>29</sup>Si MAS-NMR the data were: frequency 79.5 MHz; 5 ms single pulse MAS; 10 s relaxation delay; external standard, TMS.

Ash reaction rate was determined in terms of heat flow calculated using isothermal conduction calorimetry. Pastes were prepared by adding 5 g of fly ash to 2.5 g of alkaline solution to ensure satisfactory mixing. The paste was mixed manually for 3 min ex-situ in a sample cell, which was subsequently sealed and then immediately placed in a TAM Air Thermometric calorimeter where heat flow was recorded at a constant external temperature of 85 ± 1 ◦C.

## RESULTS AND DISCUSSION

## Effect of Grinding on Ash Physical and Chemical Characteristics

**Figure 1** shows the particle size distribution of the fly ash as received from the steam-fired power plants and after grinding. Ash F1 initially had a median particle size, **d50**, of 9µm; ash F2 of around 20µm; and F3, the coarsest, a value of approximately 34µm. After grinding, these differences narrowed substantially, with all three types of ash exhibiting **d<sup>50</sup>** values of 5µm to 7µm.

The effect of grinding is graphically illustrated in **Figure 1C**, which shows the percentage of particles under 45 and 10µm before and after milling. Grinding clearly yielded a more uniform particle size in all three varieties of ash, with 98% of the particles under 45 and 67% under 10µm. The most significant change was observed in F3, the ash with the initially largest particle size. The data in **Figure 1** show that ashes with initially different particle sizes exhibited fairly similar distributions after grinding. The inference appears to be that below a given threshold particle size, the grinding system used was ineffective.

Based on their chemical composition (**Table 1**: SiO<sup>2</sup> + Al2O<sup>3</sup> + Fe2O<sup>3</sup> ≥ 70%; CaO < 10%), F1, F2, and F3 constituted what is regarded by European standard EN 450-1 and ASTM standard C 618 as class F fly ash, i.e., pozzolanic ash resulting from the combustion of anthracite or bituminous coal. Further, to Canadian standard CSA A3001, however, which classifies fly ash in three groups depending on CaO content, only ash F3 would qualify as class F: silicoaluminous, low calcium fly ash with <8% CaO, consisting predominantly in aluminosilicate glass with varying amounts of crystalline quartz, mullite, hematite and magnetite. In the Canadian system, both F1 and F2 would be class C1 (8–20% CaO) ash, with an intermediate proportion of calcium oxide.

The results of the acid attack conducted to determine ash reactivity are compared in **Figure 2** to the overall silica and alumina contents found with chemical analysis. As the sum of the reactive silica and alumina was not over 70% in any of the materials studied, all three could be defined as Canadian standard class C1 ash.

**Figure 2** shows that grinding-based mechanical activation barely modified the potentially reactive silica and alumina identified in the starting ash. In other words, particle size would not initially appear to have any significant effect on the reactive potential of the ash. The figure also reveals the clearly smaller reactive than bulk aluminum content in ground and unground ashes F2 and F3. In these two materials, bulk alumina was on the order of 30–50% whilst the potentially reactive compound accounted for no more than 10%.

The SEM micrographs (see **Figure S1**) showed that the ashes with coarser particle sizes (F2 and F3) prior to grinding exhibited a more heterogeneous morphology, whereas the finer ash (F1) was more uniform and had a higher proportion of spherical particles.

The diffractograms for the original ash contained a hump at 22 values of 20–35◦ associated with the vitreous content of the ash. The intensity of this hump is usually related to the amorphous content. The minority crystalline phases identified were quartz, mullite, hematite, corundum and traces of CaO. The corundum and mullite contents were high in ashes F2 and F3, which justify the lower intensity of the hump in these two with respect to F1. No significant mineral differences were observed between the pre-and post-ground ash diffractograms. The sole visible changes included a slight decline in the intensity of the quartz diffraction line, which might explain the slight increase in the soluble SiO<sup>2</sup> content in the ground ash (see **Figure 2**), and a rise in the diffraction line for hematite, possibly attributable to the use of iron grinding balls.

The <sup>27</sup>Al NMR spectra (see **Figure 3**) for the original ash contained an aluminum signal at around 58 to 60 ppm associated with tetrahedral aluminum (AlT), present primarily in the vitreous phases of fly ash. Ash F1 clearly contained no octahedral aluminum (AlO), a finding consistent with the HF chemical attack results (see **Figure 2**), according to which nearly all the alumina present in this ash was reactive. In varieties, F2 and F3 mullite were detected as a shoulder at ∼2.5 ppm, associated with Al(VI), while resonances observed at ∼60 and 45 ppm were attributed to Al(IV). The intense signal at 45 ppm on the spectra for fly ash F3 was attributed to the tetrahedral aluminum in mullite, where three shares one oxygen atom surrounding Al tetrahedral (Sanz et al., 1988; Merwin et al., 1991). Ashes

F2 and F3 also exhibited a very intense signal at around 12 ppm, generated by the octahedral aluminum in corundum. This octahedral aluminum would explain the findings shown in **Figure 2**, according to which although the bulk aluminum content (∼33–50%) was high in these materials, less than half (14–15%) was dissolved by the acid. Hence the intensity of the Al<sup>O</sup> signals on the spectra for these two ashes.

The <sup>27</sup>Al spectra for F2 and F3 after grinding, also shown in **Figure 3** by way of example, were practically identical to those for the unground ash. The sole difference, the slightly lower postgrinding intensity of the shoulder at 2.5 ppm associated with mullite, may have been the result of the reduction in mullite crystal size, together with surface amorphisation of the mineral. In a study on the mechanical amorphisation of mullite drew similar conclusions (Schmucker et al., 1998). After grinding mullite for 240 h, they observed that the signal at 45 ppm was gradually replaced by one at 30 ppm, which appeared to grow at the expense of the resonances for both the tetrahedral and octahedral aluminum. In this study, whilst the intensity of the Al(VI) shoulder clearly declined, no signal was detected at 30 ppm. The apparent thrust of these findings was that mechanical activation did not significantly alter the coordination of the aluminum present in the ash: i.e., the tetrahedral and octahedral environments were the same in the ground and initial materials.

The <sup>29</sup>Si MAS-NMR spectra for the original ash (see **Figure 4**) exhibited a wide, poorly defined signal that actually enveloped a series of overlapping resonances. The area of the spectrum with absolute values greater than−108 ppm was attributed to the presence of crystal phases with Q<sup>4</sup> (0Al) environments, such as quartz (-108 to−109 ppm) and cristobalite (-113 to−114 ppm) (Engelhardt and Michel, 1987). The area between−80 and−104 ppm was primarily identified with Q<sup>4</sup> (mAl) units in the vitreous component of the ash (Criado et al., 2008). The signal at around−87 to−88 ppm, clearly visible on the F3, and as a shoulder on the F2 spectrum, was attributed to the presence of a less polymerized (more reactive) glassy component. This area may also contain a signal attributed to the Q<sup>3</sup> (3Al) units in crystalline mullite. The ash F1 spectrum exhibited a small peak at−71 ppm. Also observed but less clearly on the spectra for the other materials, that resonance may be due to the presence of scantly crystalline calcium silicate, which would partly explain the CaO content present in this ash (∼10%).

**Figure 4** also reproduces <sup>29</sup>Si spectra for F2 and F3 after grinding (GF2 and GF3), by way of example. Analogously to the Al spectra, the Si spectra showed no significant variations in the Si atoms after mechanical activation. In short, milling fly ash had a substantial effect on fineness and its distribution, but a much lesser impact on its potentially reactive silica and alumina content. A slight variation in the aluminum environments was nonetheless observed, possibly due to surface amorphisation of minority crystalline particles such as mullite. This effect was more visible in the ashes with a higher percentage of such minority crystalline phases, namely F2 and F3, despite the rather shorter ("medium" intensity) grinding time than applied in other studies (Marjanovic et al., 2014).

## Alkali Activation of Fly Ash

**Table 2** gives the compressive strength of the pastes obtained by activating the unground and ground ashes with 8 M NaOH. All the ground ash pastes exhibited strength values at least 40– 50% higher than their unground counterparts. The most striking difference was observed for ash F3. When mixed with 8 M NaOH, the unground ash hardened but was scantly compact and broke readily. Such an indication of poor quality might lead to the rejection of this ash. When the ash was ground, however, the resulting paste reached a strength of 18 MPa, a value comparable to those observed for the other materials before grinding. These findings attest to the significant role played by fineness in the

(unground) and GF2Na (ground) (C) F3Na (unground) and GF3Na (ground).

development of mechanical strength in cements made with alkali-activated fly ash.

**Figure 5** shows the heat curves for the ground and unground fly ash pastes activated with an 8 M solution of NaOH at 85◦C. Peak intensity was nearly double and appeared slightly earlier in all the pastes made with ground ash. The latter effect was more visible in ash F3, which initially had the coarsest particle size. These findings corroborated earlier reports (Zhang et al., 2016) to the effect that reducing the mean particle size raised the early age reaction rate in the material.

The XRD patterns for the initial fly ash and the activated pastes made with both ground and unground ash are reproduced in **Figure 6**. The most prominent features of the diffractograms were:


The <sup>27</sup>Al and <sup>29</sup>Si MAS NMR spectra for the initial ash and ground and unground activated materials are reproduced in **Figures 3**, **4**. A comparison of these spectra revealed clear differences.

The signal at +58 to +60 ppm associated with Al(IV) (AlQ<sup>4</sup> (4Si)-type units) present both in the N-A-S-H gel formed and in the zeolites was narrower and more intense on the <sup>27</sup>Al spectra (**Figure 3**) for the activated pastes made with ground and unground ash than on the spectra for the respective initial ashes. Pastes F1Na and GF1Na generated no signal that could be associated with Al(VI), for which no resonance was observed in the original ash either. The pastes prepared with ashes 2 and 3 (F2Na, GF2Na, F3Na, and GF3Na), in contrast, contained an intense signal at around 12 ppm, together with a shoulder at 2 ppm associated with the presence of Al(VI) in the crystalline phases of the initial ash (corundum and mullite). The signal for Al(IV) was narrower and more intense than the Al(VI) signal on the spectra for the pastes, whereas the opposite was observed on the spectra for the initial materials. The Al(VI) signal on the spectrum for paste GF2Na was very weak. Normally the Al<sup>O</sup> present in corundum and mullite is scantly responsive to alkaline activation. Here, however, ash grinding might be thought to have prompted a significant decline in corundum crystal size, along with surface amorphisation. As a result of the enhanced reactivity of this phase, part of the Al present would have subsequently been taken up in the matrix. That would explain the high strength values observed for the ground fly ash, despite its apparently low reactive Al content.

activated original ash (blue patterns) and activated ground ash (red patterns): (A) F1; (B) F2; (C) F3 (Q, quartz-SiO2; M, mullite- Al6 Si2O13; Fe, hematite-Fe2O3; Ca, calcium oxide-CaO; A, corundum; C, Calcite; H, herschelite (19-1178) NaAlSi2O6 3H2O; S, Hydroxysodalite (11-0401) Na4Al3Si3O12(OH); Y, zeolite Y (76-0108)-Na3.325H3.675Al7Si17O48).

The <sup>29</sup>Si spectra for the pastes also differed substantially from the spectra for the initial ashes (see **Figure 4**). Here the spectra for the unground ash pastes had more clearly defined signals, attributed to the higher zeolite content in these pastes, as observed in the XRD findings. As zeolites are more orderly structures than gels, their spectra are characterized by higher resolution. The <sup>29</sup>Si spectra for all the pastes exhibited a broad signal that masked overlapping resonances at positions at around−88,−93,−98,−104, and−108 ppm, attributed to tetrahedral silicons, respectively, surrounded by Al4, SiAl3, Si2Al2, Si3Al, and Si4. The intensity of the signals varied with the ash (unground and ground). A signal detected at around−87 ppm, particularly intense in paste F3Na, was partly generated by the mullite present in the starting ashes (see spectra F3 and F3Na in **Figure 4C**).

The <sup>23</sup>Na spectra for the same pastes are reproduced in **Figure 7**. All contained a signal centered over−6.5 to−10 ppm, attributed to the presence of partially hydrated sodium, and needed to balance the charge in the geopolymer when silicon was replaced by aluminum in the aluminosilicate hydrate gel structure. As a rule, in aluminosilicate glass this signal shifts to more negative values when the Si/Al ratio rises and more positive values when it declines, very likely as the result of the increase in the Na-O interatomic distance or in the coordination numbers (Duxson et al., 2005; Provis et al., 2009). Paste F3Na, which exhibited the lowest strength, had two shoulders on its spectrum, one at around 3 ppm to 5 ppm associated with the Na in unbound NaOH.nH2O (Koller and Engelhardt, 1994) which is consistent with the lower degree of reaction in this paste, and the other at around−15 ppm, possibly attributable to sodium bonded with oxygen atoms only (Koller and Engelhardt, 1994; Duxson et al., 2005; Provis et al., 2009).

**Figure 8** compares the microstructure of the chemically activated, ground and unground pastes made with ashes F1, F2, and F3. The pastes made with unground ash (F1Na, F2Na, and F3Na) were less compact on the whole than ground ash (GF1Na, GF2Na, and GF3Na) and contained a substantial number of unreacted glass microspheres, hollow (cenospheres) or filled with very small particles. The ground ash pastes (GF1Na, GF2Na, and GF3Na), especially GF3Na, had more compact matrices and a higher proportion of reacted ash. Zeolites were observed in the SEM micrographs of both ground and unground matrices, whereas XRD could detect them only in the pastes with unground ash. These data suggest that activation takes place so quickly in pastes with a high proportion of very fine particles that the new phases formed barely had time to develop a crystalline structure: i.e., reaction kinetics prevailed over thermodynamics.

Although qualitative, EDX analysis can approximate the compositional variability and homogeneity of the gels formed. The values in **Figure 9** are for CaO-SiO2-Al2O<sup>3</sup> and Na2O-SiO2.Al2O<sup>3</sup> diagrams. Further to these data, very low calcium content (∼2%) ash F3 formed which what was essentially a N-A-S-H gel. Ashes F1 and F2, with calcium contents of ∼10%, exhibited greater compositional variability. The matrices containing unground ash appeared to contain a heterogeneous mix of N,C-A-S-H-like gels with variable calcium contents (Garcia-Lodeiro et al., 2011; Walkley et al., 2016). A comparison

of the diagrams for F1 and F2 revealed that in the gels formed, the silicon/aluminum ratio was greater in the pastes prepared with ash F2, possibly due to the lower proportion of reactive aluminum present. When the ash was ground, the point distribution (red points) was slightly more homogeneous, indicating that mechanical activation, in addition to reducing the particle size, afforded the original ash a more uniform composition. As a result, the mean composition of the gels forming in pastes with ground ash was more homogeneous.

The findings for a given ash before and after grinding, compared here to determine the effect of fineness in materials with the same chemical composition, showed that the proportion of potentially reactive phase was only slightly modified by mechanical activation. Reaction kinetics and mechanical strength, however, rose significantly. These data are consistent with results reported by Zhang et al. (2016) suggesting that physical properties may play a more prominent role in ash reactivity than vitreous phase content or composition.

The smaller particle size in ground ash modified reaction kinetics, hastened dissolution and raised the initial degree of reaction, thereby generating higher early age strength. Smaller sized ash particles favored Al and Si dissolution in the medium as well as the precipitation of an N,C-A-S-H-like cementitious gel, which also affected material mineralogy and microstructure. The speedier reaction allowed practically no time for zeolites to crystallize or for the crystals to grow large enough to be detected by XRD. Another factor that may affect zeolite formation is the

need for space. As **Figure 8** shows, zeolites form preferably inside cenospheres where they have sufficient space to crystallize and grow. Grinding the ash breaks these hollow particles, leaving a smaller number of voids. Fewer zeolites and more N-A-S-H gel raise mechanical strength (Panias et al., 2007; Criado et al., 2008).

Further to these data, for ash with similar vitreous phase contents, reducing the particle size substantially raises the mechanical strength that can be developed during alkaline activation. The question posed is: given a similar particle size, how does vitreous phase content affect ash reactivity? Although after grinding the three ashes had very similar particle size (d50 ranges between 5 and 7µm), their mechanical strength values differed significantly. According to the proportion of potentially reactive phase identified by attacking the ash with acid (see **Figure 2**), the highest strength should have been recorded for ash GF1 followed by GF2 and GF3. Nonetheless, mechanical strength was the greatest in ash GF2.

Many studies have attempted to establish a criterion for assessing ash quality for the manufacture of alkaline cements (Fernández-Jiménez and Palomo, 2003; Fernández-Jiménez et al., 2005; Zhuang et al., 2016). The present authors deem that vitreous phase content, (SiO2/Al2O3)reactive ratio and fineness are

determinants that provide an approximate idea of the reactive potential of an ash. Nonetheless, the findings reported here suggest the existence of yet another parameter that may affect ash reactivity. The effect of minority elements in the ash (network modifiers such as alkalis, calcium, titanium, or iron) must be taken into consideration, inasmuch as their inclusion in the glass structure may modify its solubility. Presumably, the higher the concentration of such elements in the vitreous phase, the greater its glass reactivity. That hypothesis concurs partially with the one put forward by Zhang et al. (2016).

Determining the presence of network modifiers in the vitreous phase with the techniques deployed here is hardly straightforward. Comparing the <sup>29</sup>Si NMR spectra for ashes F1 and F2, however, revealed that the zone between−80 and−100 ppm was more intense in the latter than in the former. That may be indicative of a less polymerized and consequently more reactive glass, which would explain the higher strength in F2, fruit of a higher initial degree of reaction. The present authors feel, however, that the structure and composition of the gel forming as the main reaction product is another factor to be borne in mind when explaining mechanical strength. The EDX findings (see **Figure 9**) showed that as a rule the gel forming in pastes F2Na and GF2Na had a higher SiO2/Al2O<sup>3</sup> ratio, a finding consistent with the higher (SiO2/Al2O3)reactive ratio determined by acid attack. Reports can be found in the literature showing that more silica-rich N-A-S-H gels yield higher strength alkaline cements (Fernández-Jiménez et al., 2006; Criado et al., 2008; Provis et al., 2009).

The suitability of coal fly ash for manufacturing alkaline cements can in any event be predicted fairly precisely on the grounds of physical, mineralogical and chemical parameters (fineness, chemical composition, vitreous phase content...), which can approximate reactivity. Nonetheless, given a vitreous phase content that can be regarded as acceptable (>30%, see ash F3), fineness prevails over glass content.

This fact is promising with a view to raising fly ash reactivity. The mineralogical and chemical compositions of ash depend essentially on the coal and burning conditions used in steam-fired power plants and are consequently very difficult to control or modify. Fineness, in contrast, can be readily established and manipulated. This study shows that in ash with <80–90% of particles under 45µm, grinding would suffice to raise reactivity substantially. Ash such as F3 that initially failed to harden into a sufficiently compact specimen when alkali-activated, after grinding yielded mechanical strength of up to 18 MPa, a value similar to that was observed for unground ash F1.

## CONCLUSIONS

The conclusions drawn from the findings of this study are listed below:


## REFERENCES


## AUTHOR CONTRIBUTIONS

IG-L and OM conducted the experiments, data analyses and characterizations. AF-J and IG-L drafted the manuscript. AP helped with manuscript preparation and supervised the overall research. AF-J and IG-L finalized the manuscript.

## ACKNOWLEDGMENTS

This research was funded by the Spanish Ministry of the Economy and Competitiveness and FEDER under research projects BIA2013-43293-R, BIA2016-76466-R, and RTC2016- 4872-S.

## SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fmats. 2019.00051/full#supplementary-material

Figure S1 | SEM Micrographs of the coal fly ashes before (F1, F2 and F3) and after grinding (F1G, F2G and F3G).


of ground fly ashes (GFA) – cement mortars. Cem. Concr. Res. 27, 1365–77. doi: 10.1016/S0008-8846(97)00129-4


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Fernández-Jiménez, Garcia-Lodeiro, Maltseva and Palomo. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Studies About the Hydration of Hybrid "Alkaline-Belite" Cement

#### M. J. Sánchez-Herrero, Ana Fernández-Jiménez\* and A. Palomo

Department of Materials, Eduardo Torroja Institute (CSIC), Madrid, Spain

The mechanical performance of a hybrid cement containing 48% fly ash, 48% mineralized belite clinker, 1. 5% gypsum and 2.5% Na2SO<sup>4</sup> was analyzed and the hydration products were identified. The findings showed that the newly designed cement met European standard EN 197-1 compositional, mechanical and setting time requirements (cement type IVB). The hydration products forming in this blended "alkaline-belite" cement (as it was determined through XRD, <sup>29</sup>Si and <sup>27</sup>Al MAS NMR and electron microscopy) consisted in a mix of cementitious gels: C-(A)-S-H and N-(C)-A-S-H, which interacted and over time evolved toward the latter with no detriment to the mechanical strength developed by the cement.

Keywords: hybrid cement, belitic cement, alkaline cement, NMR-MAS, XRD

## INTRODUCTION

#### Edited by:

Miroslav Komljenovic, University of Belgrade, Serbia

#### Reviewed by:

Yiannis Pontikes, KU Leuven, Belgium Ionut Ovidiu Toma, Gheorghe Asachi Technical University of Ia ¸si, Romania

> \*Correspondence: Ana Fernández-Jiménez anafj@ietcc.csic.es

#### Specialty section:

This article was submitted to Structural Materials, a section of the journal Frontiers in Materials

Received: 22 January 2019 Accepted: 01 April 2019 Published: 18 April 2019

#### Citation:

Sánchez-Herrero MJ, Fernández-Jiménez A and Palomo A (2019) Studies About the Hydration of Hybrid "Alkaline-Belite" Cement. Front. Mater. 6:66. doi: 10.3389/fmats.2019.00066 The CO<sup>2</sup> emissions due to Portland cement production are environmentally problematic mainly due to the very high volume of demand for cement in view of concrete needs in the world. Actually, Portland cement production is continuously increasing at global level and there exists a general agreement to consider Portland cement responsible of 7–8% of the total emissions of carbon dioxide in the world (Scrivener et al., 2018a). The use of alternative fuel, the addition of minerals to Portland cement, the inclusion of mineralizers and fluxes in the raw mix, along with the development of alternative binders, are some of the sustainable solutions envisaged today by the cement industry and the scientific community (Lothenbach et al., 2011; Zhang et al., 2012). Belite (Chatterjee, 1996; García-Díaz et al., 2011; Stanek and Sulov ˇ ský, 2015), alkaline and hybrid alkaline cements (Deschner et al., 2012; García Lodeiro et al., 2013, 2016) are among the binders studied as alternatives to the traditional Portland cement.

The use of supplementary cementitious materials (SCMs, pozzolans, fly ash, blast furnace slag, limestone, etc.) as a partial substitute of Portland cement clinker in the so-called "blended cements" (Lothenbach et al., 2011; Zhang et al., 2012; Wa et al., 2018) represents not only a feasible solution but, for sure, the most investigated one. However, due to the low hydraulic activity of the SCMs, blended cements usually show delayed setting times and low early mechanical strength development, which is actually limiting the replacement of clinker by SCMs in the cement (Lothenbach et al., 2011). Obviously, many investigations are being focused on how to improve the reactivity of these SCMs. The main choices are: (i) Mechanical activation (Knop and Peled, 2016; Liu et al., 2018) (controlling the particle size distribution); (ii) Thermal activation (Vizcaíno-Andrés et al., 2015; Scrivener et al., 2018b) (especially for calcined clays—LC3 technology—); (iii) Chemical activation (Ruiz-Santaquiteria et al., 2013; Allahverdia and Malekia Mahinroosta, 2018) ("alkaline activated cement"). It is quite probable that future will lead to us to a combination of those options above mentioned like the most effective choice.

In this context the main goal of this work is to describe and discuss part of a research project (leaded by authors of this paper) linked with the development of some different types of "Hybrid Alkaline Cements" (HYC). HYC are binders containing a small fraction of Portland clinker (or even blast furnace slag) as a source of calcium, a large fraction of aluminosilicate materials (from natural origin or byproducts from some industries, such as fly ashes (Deschner et al., 2012; García Lodeiro et al., 2013, 2016) as a source of Si and Al, and obviously some alkaline additive. Due to a wide compositional range and to a high technological, economical and environmental potential, HYC are the object of considerable research (Qu et al., 2016; Al-Kutti et al., 2018).

Inasmuch as the presence of alkalis in the medium affects, aluminosilicate and Portland clinker reactivity, this study undertook an in-depth analysis of calcium silicate and calcium aluminate behavior after hydration by a number of alkaline solutions of different concentrations. The findings revealed that the compressive strength of C3A, C3S and β-C2S pastes rose in the presence of alkaline additives (Sánchez-Herrero et al., 2012, 2016). Most significantly, given the low initial reactivity of β-C2S under normal hydration conditions, the presence of such alkaline



additives hastened its hydration, affecting setting and hardening times substantially.

A mineralized belite clinker (with a belite content of over 50%) was first synthesized in the laboratory. In light of its good performance and given that the clinkerization temperature could be lowered considerably in the presence of the fluxing pair formed by "CaSO<sup>4</sup> and CaF2," the clinker was subsequently manufactured on an industrial scale with satisfactory results. The industrial belite clinker was then blended with fly ash, gypsum (the setting regulator) and Na2SO<sup>4</sup> as an alkaline additive to produce the "alkaline-belite" cement (some type of HYC).

This paper consequently describes and discusses the main singular characteristics (mechanical behavior and microstructural development) of this "alkaline-belite" cement initially formulated in the laboratory and successfully produced on an industrial scale. This advance in the design of an alkalinebelitic binder (reasonably sustainable one) was achieved due to a close collaboration between the scientific and industrial communities who simultaneously deployed some different technologies (different fields of knowledge within the "cement science") in order to produce and study the new "alkalinebelite" hybrid cement which was the object of the project. The different fields of knowledge simultaneously applied in this project were: the use of mineralizers and fluxes to produce low energy clinkers (Blanco-Varela et al., 1997; García-Díaz et al., 2011), the use of industrial by-products to generate blended cements with a low clinker factor (Deschner et al., 2012; García Lodeiro et al., 2013); and the use of affordable and innocuous natural mineral salts (alkaline additives with a minimal environmental impact) to stimulate the alkaline activation processes (García Lodeiro et al., 2016).

## EXPERIMENTAL

## Materials

A mineralized belite clinker produced in an industrial rotary kiln, a fly ash, gypsum (all supplied by the manufacturer) and 99.9% pure sodium sulfate (MERCK laboratory reagent) were

FIGURE 1 | (A) <sup>29</sup>Si MAS NMR spectrum for fly ash and second derivative of the spectrum; (B) second derivatives of the original <sup>29</sup>Si MAS NMR spectrum for fly ash and of the envelope generated by deconvolution.

used to prepare the hybrid cement. The respective chemical compositions are given in **Table 1**. The mineralogical phases present in the hybrid cement were identified with X-ray diffraction and quantified via Rietveld refinement of the XRD patterns (Rietveld, 1969). The software used: GSAS (General Structure Analysis System).

The fly ash used in this research was supplied by a thermal power plant located at Compostilla (Spain) and owned by Endesa (Spanish Electricity Company). This fly ash is a type F one (ASTM C 618) and is considered like a "good" fly ash. Some Spanish cement producers and concrete producers located at northern Spain are usual consumers of such a fly ash to produce their materials.

## Methodology

The alkaline-belite cement or hybrid cement contained 48% fly ash, 48% mineralized belite clinker, 1.5% gypsum and 2.5% Na2SO4. The belite-alkaline cement was observed to have 95% of the particles are smaller than 32µm, with a d50 of 6µm. The smaller particle size of hybrid "belitic-alkaline" was associated by operators of the mill (at least partially) with the presence of high contents of fly ash in its composition.

The hydration reactions were explored and the hydration products analyzed on cubic (3 × 3 × 3 cm<sup>3</sup> ) paste specimens prepared with a water/cement ratio of 0.29, based on consistency (EN 196-3) testing. The setting time of the pastes was determined according to the standard EN 196-3:2017 Methods of testing cement—Part 3: Determination of setting times and soundness for a water/cement ratio of 0.29.

Mechanical strength was determined on 4 × 4 × 16 cm<sup>3</sup> prismatic mortar specimens with a sand/cement ratio of 3:1 (EN 196-1) and a water/cement ratio of 0.45. The considerable rise in cement paste fluidity in the presence of fly ash generally translates into smaller mixing water content. Although European standard (EN 196-1 and EN 197-1) stipulates a water/cement ratio of 0.5 for mortars, here the value used was the result of slump testing, in line with the practice recommended in the ASTM standard (ASTM C230: Flow Table for Use in Tests of Hydraulic Cement).

The specimens were removed from the molds after 24 h of initial chamber curing at 21◦C and 95% relative humidity and then stored in a climatic chamber until the test age (2 d, 7 d, 28 d, 90 d, or 180 d). Mortar specimens were tested for compressive and flexural strength (three for bending and six for compressive strength at each test age) on an Ibertest Autotest-200/10-SW frame (Madrid, Spain).

After the paste specimens were tested for strength they were crushed to a powder and mixed with acetone and ethanol to detain hydration and prepare the materials for mineralogical and microstructural characterization via XRD and NMR.

The XRD patterns were recorded on a Bruker D8 Avance diffractometer (Karlsruhe, Germany). The instrumental conditions were: Cu K´α1,2 radiation (1.540 Å, 1.544 Å); no monochromator; goniometer radius, 217.5 mm. The settings for the hydrated pastes were: variable 6-mm divergence slit (hydrated pastes); 2θ angle, 5◦ to 60◦ ; step time, 0.5 s; step size, 0.02◦ . The anhydrous samples were scanned under different conditions for Rietveld quantification: divergence slit, 0.5◦ ; 2θ scanning range, 5◦ to 70◦ ; step time, 2 s; step size, 0.02◦ .

FIGURE 2 | Anhydrous and 2 d, 7 d, 28 d, and 180 d hydrated paste diffractograms Legend: b, C2S(86-0398); a, C3S(70-8632); q, quartz [SiO2](46-1045); g, gypsum [Ca2SO4.2H2O] (33-0311); l, mullite[Al6Si2O13] (73-1389); n, Na2SO4 (74-2036); e, ettringite [Ca6Al2(SO4)3(OH)12.26H2O] (41-1451); c, calcite[CaCO3] (86-0174); \*, C12A7 (70-2144); v, vaterite [CaCO3](74-1867); (in brackets, PDF codes).

TABLE 2 | Mortar strength (4 × 4 × 16 cm<sup>3</sup> specimens) and setting times for "alkaline-belite" hybrid cement.


The <sup>29</sup>Si and <sup>27</sup>Al nuclei were scanned on a Bruker MSL 400 NMR spectrometer. The conditions for <sup>29</sup>Si were: resonance frequency, 79.5 MHz; spinning rate 10 kHz; pulse sequence, single pulse (5 µs); recycle delay, 10 s; number of transients, 4912; internal standard for measuring chemical shift (in ppm), tetramethylsilane (TMS). For <sup>27</sup>Al they were: resonance frequency, 104.3 MHz; spinning rate, 10 kHz; recycle delay, 5 s; pulse sequence, single pulse (2 µs); number of transients, 400; internal standard for chemical shift measurements (in ppm), Al(H2O)3<sup>+</sup> 6 .

The <sup>29</sup>Si NMR spectra were deconvoluted in order to determine the number and position of the signals analyzed. The second derivative criterion was used (Kucherov and Kochubei, 1983; Pérez et al., 2014). This procedure is highly useful when analyzing spectra with wide, poorly defined signals, inasmuch as it resolves closely juxtaposed resonant frequencies and determines the exact position of each. In this particular case two software applications (OriginPro 9 and MestReNova) were used with the objective of avoiding practical mistakes. The <sup>29</sup>Si MAS NMR signals exhibited a Lorentzian/Gaussian (L/G) ratio of 0.5 (independent on the software). First derivative analysis reveals the presence of peaks that appear as shoulders or weak signals on the original spectrum, whilst the second derivative furnishes more precise information in the form of minima, the value of which concurs with the peak on the main signal. The number of points taken into account during the analysis (in the zone of the spectrum between −60 and −130 ppm) was 450 approximately. **Figure 1** is an example of <sup>29</sup>Si MAS NMR spectra of our fly ash. This spectrum at **Figure 1** is an example of analysis by following the already explained procedure.

Finally some samples were vacuum dried and carbon-coated for examination under a HITACHI S-4800 (Tokyo, Japan) scanning electron microscope fitted with a Bruker energy dispersive spectroscope to confirm the working hypotheses. For BSEM analysis, the samples were embedded in epoxy resin, cut and polished with ethanol prior to vacuum drying and carbon coating.

## RESULTS AND DISCUSSION

## Material Properties

Material compressive strength and setting time data are given in **Table 2**. The mechanical strength development of mortars (4 × 4 × 16 cm<sup>3</sup> prismatic) was 10.8 MPa at 2 days 40.0 MPa at 28 days, and 51.9 MPa at 180 days. The initial setting time was t<sup>i</sup> = 83 m. and the final t<sup>F</sup> = 193 m; the belitic hybrid cement met European standard requirements for 32.5 type cement (EN 197- 1 Cement—Part 1: Composition, specifications and conformity criteria for common cement).

As noted, this cement contained 48% mineralized belite clinker, which is characterized by low early age reactivity (Taylor, 1997). According to these findings, the constituents of the alkaline-belite cement generated synergies during hydration that contributed favorably to strength development. The heat of hydration released by the belite clinker fraction, in conjunction with the presence of the Na2SO4, sufficed to alkali-activate the fly ash, affording the system good initial mechanical behavior. Furthermore, as noted in prior studies (Sánchez-Herrero et al., 2017), the data appeared to indicate that the presence of the alkaline additive stimulated greater strength development in the

calcium silicates and aluminates present in the clinker than observed during normal water-based hydration.

In previous papers (Donatello et al., 2013) authors have already demonstrated that hybrid cements behave quite well TABLE 3 | Chemical shifts (ppm) recorded after deconvolution of <sup>29</sup>Si MAS NMR spectra for hybrid cement and 28 d and 180 d cement pastes.


when located in aggressive environmental conditions (sulfates, marine, acidic, etc.). Additionally, they do not undergo the ASR since alkalies are consumed in the reaction of activation of fly ashes (Malvar and Lenke, 2006).

## Reaction Products

The XRD patterns for the anhydrous hybrid cement and the 2 d, 7 d, 28 d, and 180 d hydrated pastes are reproduced in **Figure 2**.

All the diffractograms exhibited a hump associated with the vitreous fraction of the fly ash and the existence of a cementitious gel precipitating during hydration. The intensity of the diffraction lines for belite was observed to decline steadily with hydration time.

All the pastes generated diffraction lines attributed to quartz and calcite. Ettringite was identified in the 2 d and 7 d patterns, whereas the 28 d and 180 d materials contained vaterite. No signals associated with portlandite were identified at any of the ages studied.

**Figure 3** contains the <sup>29</sup>Si MAS NMR and <sup>27</sup>Al MAS NMR spectra for the anhydrous cement and the 28 d and 180 d pastes. Some of the signals on the hydrated paste spectra were absent on the scan for the anhydrous cement. These wide, poorly defined signals may be indicative of the presence of a mix of several components.

The <sup>29</sup>Si MAS NMR deconvoluted spectra for the hybrid cement and its hydrated pastes are reproduced in **Figure 4**, whilst **Table 3** give the chemical shifts and area for each signal identified on the spectra analyzed.

The <sup>29</sup>Si MAS NMR spectrum for the anhydrous cement exhibited a narrow signal in the clinker fraction at around −71 ppm attributed to belite, and a series of wider signals associated with the C2S and C3S Q<sup>0</sup> units in alite (Sánchez-Herrero et al., 2016, 2017). The area of the spectrum between −81 ppm and −120/−130 ppm was associated with the fly ash fraction in the cement. Deconvolution of this area of the spectrum revealed the presence of several components. The signals between −85 ppm and −108 ppm in the anhydrous cement, concurrent with the presence of the ash, were attributed to Q<sup>4</sup> (mAl) units (**Figure 4** and **Table 3**) for, as these signals tended to disappear in the 28 d and 180 d spectra, they were associated with the vitreous phase in the fly ash (Engelhardth and Michael, 1987; Palomo et al., 2004). The signals at (absolute) values higher than −109 ppm, however, remained practically unchanged during cement hydration and were therefore associated with Q<sup>4</sup> (0Al) units in the various crystalline forms of silica (Engelhardth and Michael, 1987). Mullite, in turn, a crystalline phase detected in this material through XRD, generated resonance frequencies at −86 ppm, −91 ppm, and −95 ppm (Merwin et al., 1991). The Q<sup>0</sup> and Q4 (mAl) units, respectively associated with the unreacted clinker and fly ash fractions, declined in intensity in the 28 d and 180 d pastes (García Lodeiro et al., 2013, 2016).

Attribution of the new signals appearing in the 28 d and 180 d pastes was not straightforward, for different units in the cementitious gel may overlap. In such complex systems, a single signal may be assigned to different units of the gel or gels precipitating during the reaction. The signal at −85 ppm for instance, may be attributed to a Q<sup>2</sup> unit in a C-S-H gel (Skibsted and Andersen, 2013), a Q<sup>3</sup> (3Al) unit in a C-A-S-H gel (Puertas et al., 2011; Myers et al., 2015) or a Q<sup>4</sup> (4Al) unit in an N-A-S-H gel (García-Lodeiro et al., 2010, 2011). In this study, attribution was based on the literature as well as the findings of other techniques discussed below (27Al MAS NMR and BSEM/EDX).

Further to those criteria, the signals at −78.8/−79.0 ppm were associated with Q<sup>1</sup> units in C-S-H gel. Similarly, the signals at −81.6/−81.3 ppm and −84.6/−83.8 ppm were attributed to Q2 (1Al) and Q<sup>2</sup> (0Al) units in C-(A)-S-H-like gels. C-S-H gel is the main reaction product that precipitates during ordinary Portland cement hydration. The uptake of aluminum in bridging tetrahedral positions in the C-S-H gel structure gives rise to Q2 (1Al) units, favoring the formation of C-(A)-S-H gels. Both gels precipitated during hydration of the mineralized belite clinker fraction of the cement.

The gels that form in cementitious systems with a very low calcium content (such as in the alkaline activation of fly ash) are alkaline aluminosilicate hydrates (N-A-S-H gels) (Fernández-Jiménez et al., 2006). Such gels, which have a three-dimensional structure consisting in Q<sup>4</sup> (mAl) units, constitute the main hydration product of fly ash alkaline activation. Further to that criterion, the signals located at around −88.5 ppm, −92.2 ppm, −95.7 ppm, −101.0 ppm, and −106.4 ppm on the 28 d and 180 d hybrid spectra would be respectively associated with the Q<sup>4</sup> (4Al), Q4 (3Al), Q<sup>4</sup> (2Al), Q<sup>4</sup> (1Al), and Q<sup>4</sup> (04Al) units in a N-A-S-H gel (Engelhardth and Michael, 1987; Fernández-Jiménez et al., 2006).

In these complex systems, however, the presence of calcium in the medium around the N-A-S-H gels may induce partial substitution of calcium for sodium, prompting the formation of N-(C)-A-S-H-like gels (García-Lodeiro et al., 2010, 2011) and with it slight shifts in resonance frequencies. That would explain the fluctuations observed in the positions of the signals on the 28 d and 180 d spectra, in addition to the evolution toward the precipitation of an N-(C)-A-S-H-like gel.

Two new signals were detected in the 28 d spectra at around −93.8 and −95.8 ppm, associated with Q<sup>3</sup> (mAl) units. The presence of those signals can be explained on the grounds of the existence of phases precipitating in accordance with the following two mechanisms.


AFt. The <sup>27</sup>Al MAS NMR spectra reproduced below in **Figure 3B**, support that theory, as they contain signals for scantly crystalline AFt and AFm, phases that would not be detected with XRD.

An analysis of the tetrahedral aluminum based on the <sup>27</sup>Al MAS NMR spectra for the alkaline-belite hybrid cement and its pastes (**Figure 3B**) revealed that the intensity of the signal at +80 ppm, associated with calcium aluminates in the clinker (Baptiste d'Espinose de Lacaillerie et al., 2008; Sánchez-Herrero et al., 2012; Skibsted and Andersen, 2013), was barely perceptible in the 28 d materials and disappeared almost entirely after 180 d.

In the cement pastes, the intensity of the signal associated with the tetrahedral component of the fly ash (+52/54 ppm) declined and shifted, with the appearance of a number of signals at around +60 ppm. The inference would be that part of the aluminum was taken up into a gel or mix of gels and surrounded by two, three or four silicon atoms, providing further evidence of complex structures comprising a mix of cementitious gels of very different characteristics. According to the literature (García-Lodeiro et al., 2010, 2011; Puertas et al., 2011; Myers et al., 2015) the Al<sup>T</sup> for N-A-S-H gel appears close to +57/+60 ppm while Al<sup>T</sup> in C-A-S-H gel appears at +65/+68 ppm. By comparing <sup>27</sup>Al NMR spectra at 28 and 180 days, the area of the signal at around +65/+68 ppm attributed to a C-A-S-H gel rose substantially after 180 days of hydration.

The <sup>27</sup>Al NMR spectra for the 28 and 180 days pastes (see **Figure 3B**) exhibited signals at around +12 ppm, +9.4 ppm, and +4 ppm, indicative of octahedral aluminum and respectively associated with ettringite (C6AˆS3H32), calcium monosulfoaluminate (C4AˆSH12) or calcium monocarboaluminate (C4ACHˆ <sup>11</sup>) and the third aluminate hydrate (TAH) (Shimada and Francis Young, 2004; Baptiste d'Espinose de Lacaillerie et al., 2008; Sánchez-Herrero et al., 2012; Skibsted and Andersen, 2013) none of these hydration products were detected with XRD. In contrast, with its high sensitivity to <sup>27</sup>Al nuclei, MAS NMR identifies compounds difficult to detect with other techniques due to the small amounts present or their low crystallinity.

**Figure 5** shows the areas attributed to the Q<sup>1</sup> , Q<sup>2</sup> , Q<sup>3</sup> , and Q<sup>4</sup> units obtained after deconvoluting the <sup>29</sup>Si MAS NMR spectra. The sum of the areas for the Q<sup>1</sup> + Q<sup>2</sup> units (associated with a C-(A)-S-H gel) rose in the 180 d materials, whilst the sum of the Q<sup>4</sup> units, corresponding to a N-A-S-H gel, declined at that age. The appearance of Q<sup>3</sup> units in the 180 d pastes denoted a rise in the total area of the Q<sup>1</sup> + Q<sup>2</sup> + Q<sup>3</sup> units associated with C-A-S-H gel precipitation, further supporting the evolution over time of the cementitious gels that precipitate during alkaline-belite hybrid cement hydration. These chemical and nanostructural changes in the cementitious gels did not compromise mechanical strength development in either the pastes (**Figure 5**) or the mortars (**Table 2**).

These findings corroborated the analysis of the <sup>29</sup>Si MAS NMR spectra, namely that the cementitious gel precipitating during hybrid "alkaline'-belite" cement hydration was actually a mix of interacting gels that evolved over time toward a gel with characteristics reminiscent of N-(C)-A-S-H gel. The evolution of cementitious gels toward a single product implies that part of the tetrahedral aluminum released into the medium was taken up into the secondary hydration products in the form of octahedral aluminum.

Further to the above analysis of the <sup>29</sup>Si and <sup>27</sup>Al MAS NMR findings, three types of cementitious gels precipitated during hybrid "alkaline-belite" cement hydration: C-(A)-S-H, C- (N)-A-S-H and N-(C)-A-S-H. That was possible due to system characteristics, determined by its 48% clinker, 48% fly ash, 1.5% gypsum, and 2.5% Na2SO<sup>4</sup> contents. These cementitious gels interacted and their characteristics evolved over time toward those of a gel similar to C-(N)-A-S-H.

The EDX microanalyses of BSEM micrographs showed that the majority elements in paste chemical composition were calcium, silicon and aluminum, although different amounts of sulfur were also identified in most (**Figure 6**). The microanalysis conducted at point 1 on the micrograph detected the presence of a C-S-H-like gel with small amounts of aluminum (C-(A)-S-Hlike gel), whereas the microanalysis at point 2 was indicative of the presence of a C-A-S-H-like gel (**Figure 6**).

The CaO/SiO<sup>2</sup> and Al2O3/SiO<sup>2</sup> molar ratios for the two pastes shown in **Figure 7** were calculated from the microanalyses of the BSEM micrographs. The 28 d cementitious gels were observed to precipitate primarily in two areas with very different compositional ranges. In one, low calcium N-A-S-H-like gels precipitated where some cenosphere or fly ash particle was previously located, whilst C-(A)-S-H-like gels were found in the other zones (where some particle of C3S or C2S was existing). The 180 d findings showed that the chemical composition of both types of gels evolved toward a single C-A-S-H-like gel.

In short, the BSEM/EDX findings for the 28 d and 180 d pastes confirmed the precipitation of different types of cementitious gels and hence the existence of a mix of gels. Due to the interaction between these gels as hydration progressed, this cementitious mix evolved toward the precipitation of a hydration product with characteristics reminiscent of C- (N)-A-S-H gel. The excellent technological behavior of the belitic-alkaline hybrid cement can only be explained by this nanostructural circumstance: two different cementitious gels being formed simultaneously. The extent to which formation of caco3 might contribute to the mechanical strength development has not been analyzed in this paper but authors consider it will be quite interesting point to analyze in the future since xrd seems to indicate that caco3 is increasing with time of curing.

## CONCLUSIONS

The most prominent conclusions to be drawn from the present study are as follows.


## AUTHOR CONTRIBUTIONS

MJS-H conducted the experiments, data analyses and characterizations. MJS-H and AF-J drafted the manuscript. AP helped with manuscript preparation and supervised the overall research. AF-J finalized the manuscript.

## ACKNOWLEDGMENTS

This research was funded by the Spanish Ministry of the Economy and Competitiveness and FEDER under research projects BIA2016-76466-R and RTC2016-4872-S.

## REFERENCES


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Sánchez-Herrero, Fernández-Jiménez and Palomo. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Microgravity Effect on Microstructural Development of Tri-calcium Silicate (C3S) Paste

Juliana Moraes Neves <sup>1</sup> \*, Peter J. Collins <sup>1</sup> , Ryan P. Wilkerson<sup>2</sup> , Richard N. Grugel <sup>2</sup> and Aleksandra Radlinska ´ 1

<sup>1</sup> Department of Civil and Environmental Engineering, The Pennsylvania State University, University Park, PA, United States, <sup>2</sup> Marshall Space Flight Center-NASA, Huntsville, AL, United States

For the first time, tricalcium silicate (C3S) and an aqueous solution were mixed and allowed to hydrate in the microgravity environment aboard the International Space Station (ISS). The research hypothesis states that minimizing gravity-driven transport phenomena, such as buoyancy, sedimentation, and thermosolutal convection ensures diffusion-controlled crystal growth and, consequently, lead to unique microstructures. Results from SEM micrographs, image analysis, mercury intrusion porosimetry, thermogravimetry, and x-ray diffraction revealed that the primary differences in µg hydrated C3S paste are increased porosity and a lower aspect ratio of portlandite crystals, likely due to a more uniform phase distribution. Relevant observations led by the presence or absence of gravity, including bleeding effect, density, and crystallography are also presented and discussed.

#### Edited by:

Maria Juenger, University of Texas at Austin, United States

#### Reviewed by:

Jinrui Zhang, Tianjin University, China Hongyan Ma, Missouri University of Science and Technology, United States

#### \*Correspondence:

Juliana Moraes Neves jmn273@psu.edu

#### Specialty section:

This article was submitted to Structural Materials, a section of the journal Frontiers in Materials

Received: 19 January 2019 Accepted: 08 April 2019 Published: 24 April 2019

#### Citation:

Moraes Neves J, Collins PJ, Wilkerson RP, Grugel RN and Radlinska A (2019) Microgravity Effect ´ on Microstructural Development of Tri-calcium Silicate (C3S) Paste. Front. Mater. 6:83. doi: 10.3389/fmats.2019.00083 Keywords: microgravity, C3S hydration, microstructure, porosity, portlandite CH, sedimentation, bleeding, phase distribution

## INTRODUCTION

Exploration of extraterrestrial bodies, such as the moon, can be facilitated by extended stays, which will require fabrication of habitats and other supporting infrastructure. Given the cost of transporting materials into space, it is envisioned that in situ materials, e.g., lunar regolith, would be required in such construction and a cement-like binder would be needed for creating resilient habitats. Solidification of cement in a microgravity (10−<sup>6</sup> g or µg) environment, however, is not well-understood. As a first step, the primary objective of this research is to experimentally document microgravity influence on the microstructure development of a tricalcium silicate (C3S) paste.

The Microgravity Investigation of Cement Solidification (MICS) research project is in strong alignment with NASA's human and space exploration needs. The experiment directly addresses two objectives identified by NASA Space Technology Roadmaps and Priorities (National Research Council, 2012), listed as (1) exploring the evolution of the solar system and the potential for life elsewhere and (2) expanding the understanding of the Earth and the universe. Furthermore, it is possible that indigenous materials can be used to manufacture a binder to make concrete on the surface of the Moon, which would enable a cost-effective solution for long-duration human missions (Khoshnevis, 2004). The role the extreme environment would have on life has been researched (lsson-Francis et al., 2018), and the extraterrestrial building material would have to help combat these challenges.

Cement has been studied for over a century with pioneering research conducted by Bogue, T. C. Powers, and T. L. Brownyard. Together they remarkably contributed to the science of cement and concrete materials (Powers and Brownyard, 1946; Bogue, 1947; Powers et al., 1954; Powers, 1958), including characterization of basic fresh and hardened properties. Since then, understanding the complex hydration of portland cement has improved substantially, but is not yet fully understood. One example is the insufficient comprehension of the mechanism controlling the main heat evolution peak of cement hydration (Bullard et al., 2011; Scrivener et al., 2015), which hampers development of accurate prediction models (Bullard et al., 2010). The need for improved methods of studying cement hydration has been noted (Provis, 2015) and MICS contributes by providing a novel approach to understanding how gravity affects materials processing. Moreover, this novel insight will promote our understanding of Earth-based cementitious science, and advance development of binders for use on extraterrestrial bodies.

It should be noted that the work described here is part of the on-going MICS research project and focuses on the microstructural development of high water-to-cement ratio tricalcium silicate (C3S) paste mixed and cured in a microgravity environment. C3S is the main mineral component of typical commercially available portland cement and dictates most of the kinetics and early properties (Taylor, 1990; Mindess et al., 2003; Scrivener, 2004; Thomas et al., 2009). In principle, investigating pure compounds, such as C3S, constitutes elementary research, which eases and enhances the analysis of more complex systems such as portland cement mixtures. In portland cement systems, the presence of aluminates, sulfates, alkalis, and other impurities affects pH and supersaturation of the solvent (i.e., pore solution), which makes interpretation of results and the conclusion-drawing process more complicated. As such, a thorough understanding of pure phases is necessary before more complex systems can be analyzed. To date, pure C3S and its hydration products (C-S-H and CH<sup>1</sup> ) have been extensively studied theoretically (Young and Hansen, 1987; Bentz and Garboczi, 1991; Joseph et al., 2017) and experimentally (Thomas et al., 2009; Bazzoni, 2014; Hu et al., 2016a,b).

As of now, no experiments have fully documented the effect of microgravity on the microstructural development of portland cement or any of its compounds. There have been cement hydration studies performed aboard a parabolic flight path, which allows for ∼20 s of a 10−<sup>2</sup> -10−<sup>3</sup> gravity environment (Meier et al., 2015; Lei et al., 2016; Meier and Plank, 2016). The short time duration limited the experiments to investigating instantaneous ettringite precipitation. The authors reported that the microgravity environment generally leads to increase in the amount of ettringite precipitated. However, the change in gravity has only a minor effect on the aspect ratios of ettringite formed from a neat cement by a slight reduction. The most notable changes occur in the presence of PCE polymers as the aspect ratio of the ettringite was significantly changed depending on the anionicity of the polymer. Outside of the cementitious materials TABLE 1 | Percentages (10, 50, and 90) of C3S particles with a diameter equal to or smaller than the indicated values.


For example, 10 percent of the C3S particles are smaller than 1.60 µm. Results obtained from laser diffraction.

realm, one experiment documenting the effects of microgravity on the growth of an inorganic crystal (sodium chloride) was conducted aboard the ISS (Fontana et al., 2011). The authors reported morphological differences between crystals precipitated in µg and typical NaCl precipitated on Earth (Fontana et al., 2011). Qi et al. (2017) modeled how settling motion could affect the morphology of a growing dendritic crystal. These studies (Fontana et al., 2011; Meier et al., 2015; Lei et al., 2016; Meier and Plank, 2016; Qi et al., 2017) advocate that gravitational forces impact crystal growth and can provide ground for interpreting the results of the present research.

## MATERIALS AND EXPERIMENTAL SETUP

## Characterization of Anhydrous Materials: Tri-calcium Silicate (C3S) and Hydrated Lime

Pure compound C3S, hydrated lime, and distilled water were used in this work. High purity triclinic C3S with a surface area of 3,500 cm<sup>2</sup> /gram was obtained from Mineral Research Processing (Meyzieu, France). Calcium hydroxide (Alfa-Aesar) was added to the mixing water, which is explained in detail in Mixture Proportions. The present study refers to the calcium hydroxide used in the mixing water as hydrated lime, so it can be easily differentiated from the calcium hydroxide precipitated during the C3S hydration process, which is referred to as portlandite.

The evaluation of the anhydrous powder of C3S and hydrated lime was first conducted to characterize their size and morphology. Dry-dispersion laser diffraction (Malvern Mastersizer 3000) was employed to determine the particle size distribution of C3S. Approximately 100 mg of powder was dispersed by an Aero S dry dispersion unit using a standard venturi at 4 bar air pressure and a 25% feed rate. The refractive index and adsorption were assumed to be, respectively, 1.68 and 0.1. **Table 1** shows diameters D10, D50, and D90. Determining the particle size distribution of hydrated lime was not performed as most of the powder was expected to dissolve once added to water.

Scanning electron microscope (SEM) images of the anhydrous C3S and hydrated lime powders were also taken (**Figures 1**, **2**). In the specific case of hydrated lime, the images prove that it differs from the hydrated materials, which are shown and discussed in the results section. The samples were prepared by sprinkling the powders over carbon tape and placing them in an SEM at low vacuum mode. It can be seen that the C3S particles are consistent with the results obtained through laser diffraction, as most of the particles are smaller than the 10µm scale bar.

<sup>1</sup>The usual cement chemistry notation is used throughout this manuscript, where C=CaO, S=SiO2, and H=H2O for both the anhydrous and hydrated phases.

FIGURE 1 | SEM image of anhydrous C3S particles showing irregular morphology, smooth surface, and sizes often smaller than the 10µm (scale bar), in agreement with laser diffraction results.

## Mixture Proportions

The mixture evaluated in this research consisted of C3S mixed with lime-water at a water-to-cement ratio (w/c) of 2.0 by mass (5 grams of C3S and 10 grams of lime-water). The lime-water consisted of a 15 mmol/liter solution of calcium hydroxide, where 1.12 grams of hydrated lime was added to 1,000 g (1 liter) of distilled water, sealed, and continuously stirred for 24 h at room temperature to ensure stabilization of the solution.

In addition to enhancing the degree of hydration, a high w/c (2.0) enhances crystal growth through coarsening the porosity of the paste (Scrivener, 1989; Taylor, 1990; Mindess et al., 2003; Thomas and Jennings, 2019). Magnifying the overall microstructural development also nourishes the µg effects and aids in their identification. Moreover, the high w/c allowed for the easiness of the manual mixing process (discussed in Experimental Setup) without need for chemical admixtures. If pure water was used rather than lime-water in this high w/c system, the initial reaction rate would be too rapid and as such, unrealistic. Using lime-water mimics the supersaturation degree with respect to portlandite in the pore solution of cement [pH of 12.5 ± 0.1(Rajabipour et al., 2015)], controls the initial reaction rate (Hu et al., 2016a,b), and still enhances nucleation and crystal growth (Bazzoni, 2014).

## Experimental Setup

Commercially available plastic bags (Burst Pouches <sup>R</sup> ) were used as mixing containers (**Figure 3**). A total of 12 identical pouches were prepared; 6 were sent to the ISS on the OA-9 resupply mission and 6 remained on the ground as control samples. The pouches entailed two separate compartments divided by an internal, burstable seal which allowed mixing materials under sealed conditions. Individual compartments of the pouches were filled with C3S and aqueous solutions under laboratory conditions on Earth. This setup allowed the first contact between the C3S and lime-water to take place while aboard the ISS. Space and ground samples were mixed simultaneously and at the same conditions (temperature of 20 ± 2 ◦C). The pressure and relative humidity conditions (1 ATM and 35% RH) were also constant, since they are characteristics of the air confined in the sealed pouches during the preparation. As such, gravity is considered the only variable between the space and ground experiments.

The mixing procedure consisted of exerting pressure on the lime-water by rolling the pouch and effectively pushing the solution forward until the internal seal was broken, after which the solution came in contact with the anhydrous C3S. The hydrating C3S paste was then manually mixed with a rubberized spatula for ∼3 min, or until homogeneity was achieved. Finally, a clamp was used to assure consolidation of the paste, as shown in **Figure 3**. The processed ground and space samples remained within the sealed pouch for the entire hydration period at undisturbed, controlled temperature (20 ± 2 ◦C) conditions. The space samples were allowed to hydrate for 42 days aboard the ISS until returning to Earth for analysis. Samples were retrieved after splash down and immediately transported in insulated containers to Marshall Space Flight Center in Huntsville, AL, where the initial analysis took place.

Both the experimental setup and mixing procedure were kept simple to meet two major requirements: (1) to reduce the volume of materials and shipping cost, which is estimated to be at least \$10,000 per pound (0.454 kg) of material (Futron Corporation., 2002), and (2) to meet NASA safety requirements that consider cement and its compounds to be hazardous. To address this issue, the pouches were tested for pressure, temperature, and longterm storage of high pH materials. In addition, the samples had to be triple-contained while the experiments were executed by the astronauts on the ISS. The first containment level was the pouch itself, the second was a plastic bag seen in **Figure 3**, and the third was a portable glovebag which was fixed on top of the Maintenance Workbench Area (MWA). Note that the hardened samples were given the lowest toxicity level (0 rating).

## METHODS

## Scanning Electron Microscopy (SEM)

Shortly after returning to Earth, the space samples underwent a visual comparison with the ground samples. This assisted in identifying the most pronounced differences and subsequently selecting the series of most relevant tests. The fractured surfaces of samples were examined using a Hitachi S-3700N SEM immediately after their removal from the pouch, at day 56 from the initial hydration. The microscope operated under variable pressure conditions (10 to 20 Pa) with an accelerating voltage of 10 or 15 keV. In the variable pressure mode, the capability of the instrument is limited to backscattered electron (BSE) micrographs. A probe current of 25 µA and an aperture of 4 were used to maximize image resolution.

Additionally, polished cross sections were prepared at day 152 after initial hydration. The sample preparation consisted of immersing the samples in 200 ml of isopropyl alcohol for 48 h,

followed by 30 h of drying under vacuum and room temperature (22◦C). Subsequently, the samples were mounted in medium grade acrylic resin, ground through a series of SiC papers, and finally polish using 0.25µm diamond paste. The polished crosssections were examined in a FEI Q250 and BSE images were taken. The operating conditions are specified at the bottom of each micrograph presented in the results section.

## Image Analysis

Analysis of the porosity through the BSE images taken of the polished samples was done by creating a binary image separating the pores from all other phases. The greyscale value to determine the upper boundary of the pores was done using the overflow method (Wong et al., 2006). There were 20 images captured for the space samples and 30 images for the ground samples at a magnification of 500x. A statistical analysis was done ensuring the number of images acquired for each sample was enough to be 95% confident that the true mean is within 10% of the sample mean. The images captured were 1,536 × 1,024 pixels with each pixel representing 0.27 × 0.27µm. The brightness and contrast settings were adjusted to ensure a greyscale histogram with a broad range of values for the present phases. The images showed a large amount of porosity, C-S-H, and portlandite with minimal anhydrous C3S being visible.

After creating the binary image using the greyscale value obtained from the overflow method for thresholding, the area fraction of porosity was calculated. Regions with <10 pixels were regarded as noise and not included in the porosity calculation, which is consistent with the Wong et al. study (Wong et al., 2006). As such, the smallest pore size that was detected and counted in this method had diameter of 0.96 microns. As will later be discussed in more detail, the ground sample exhibits a layered structure due to sedimentation. Each layer of the ground sample was individually analyzed. Multiple size measurements were taken and averaged to calculate the corresponding fraction of each layer. The values obtained were used to calculate a weighted average of porosity.

## Mercury Intrusion Porosimetry (MIP)

Prior to the MIP test, ground and space samples were immersed individually in 200 ml of isopropyl alcohol for 48 h for solvent exchange, which was followed by 34 h of drying under vacuum at room temperature. Subsequently, a Micromeritics AutoPore V 9620 MIP mercury penetrometer was used to assess the porosity, pore size distribution, bulk density, and skeletal density of the pastes. The mercury temperature during testing of ground and space samples was 19◦C.

## Thermogravimetric Analysis (TGA)

Prior to TGA analysis, the space samples underwent 42 days of hydration under µg condition, followed by at 109 of hydration under 1 g, totaling 151 days of hydration. Unopened ground and space samples of same age were removed from the pouches and stored for 10 days in a nitrogen-purged chamber, at 32◦C and 65% relative humidity. Subsequently, the partially dried samples were finely ground with mortar and pestle immediately before the TGA measurement. SDT Q600 (TA Instruments) was then used to estimate the degree of hydration and to quantify the amount of portlandite in the pastes. Approximately 20 grams of

FIGURE 3 | A typical processed pouch containing hydrating cement paste and a clamp.

ground material was placed in the nitrogen purged chamber of the instrument and heated to 110◦C. This temperature was held for 20 min to eliminate any remaining free water and then raised in increments of 10◦C per minute until 800◦C or higher.

## X-Ray Diffraction (XRD)

Samples of 151 days of age were taken from the pouch immediately before the analysis and, without additional preparation, were crushed using a mortar and pestle for approximately 3 min. Due to the high water-to-cement ratio, the crushed material was somewhat wet when back-filled into a zerobackground sample holder for x-ray diffraction. XRD data was collected using a PANalytical X'Pert PRO MPD diffractometer in a conventional Bragg–Brentano θ-θ configuration. CuKα X-ray (λ = 1.5418 Å) was generated using 40 mA and 45 kV operating conditions. The optics setup included incident divergence slit fixed at 0.25◦ , 0.04 soller slits, incident/diffracted anti-scatter slits and a nickel filter. The samples were scanned continuously between 5◦ and 65◦ 2θ at a step size of 0.013◦ , and spinning at 60 rpm to maximize the intensity (counts).

## RESULTS AND DISCUSSION

## SEM

The micrographs obtained at different magnifications provided a wide range of detail, revealing information ranging from the overall distribution of phases to the micron-scale crystal morphology. As it can be seen in **Figures 4**, **5**, the pastes cured in µg differ from the control samples cured in 1 g in several aspects. First (**Figure 4**), is the greater number of trapped air bubbles found in the space specimens, which is attributed to the lack of buoyancy in a µg environment. Buoyancy is a gravity-driven phenomenon that results from differences in specific gravities among components comprising a given system (Turner, 1979). Under the influence of gravity, air bubbles, for example, are expected to move upwards through the initial watery cement paste due to their lower density [≈1.2 kg/m<sup>3</sup> (Cavcar, 2000)], while also promoting fluid flow (Barge et al., 2015).

The 1g micrograph in **Figure 5** clearly shows that layering and "bleeding" effects can manifest even in a paste that is merely 3 mm high, as is the case for the ground samples. Bleeding occurs at very early stages of cement hydration and generates a gradient water-to-cement ratio across the sample height, which results in a porosity gradient (Han and Wang, 2016). The physical explanation behind the commonly seen phenomenon is the gravity-driven sedimentation. Sedimentation promotes sinking of anhydrous C3S, as well as of C-S-H, and portlandite [specific gravities of 3.15, 2.3–2.6, and 2.24, respectively (Mindess et al., 2003)], while raising the calcium-saturated pore solution (specific gravity of ≈1) to the surface. The film of supersaturated solution promotes unimpeded growth of portlandite at the specimen's top surface, as seen in **Figure 6**. As expected, the sedimentation layering or bleeding effects do not take place in space; rather, the pastes hydrated in µg develop a more uniform density, as shown in the µg micrograph (**Figure 5**). Gravity-driven sedimentation also seems to affect the consolidation process of cementitious materials. **Figure 7** suggests that C3S hydrated in microgravity generates a more porous paste in comparison to the ground control. It also appears that, in 1g, the self-weight of the hydrating paste generates sufficient compaction to reduce the overall spacing between the hydrating phases. In µg, the lack of a directional gravitational force allows the anhydrous C3S to float in the solution, enhancing the porosity.

The size, shape and distribution of portlandite crystals also differs substantially between 1g and µg samples, in agreement with the ISS NaCl experiment (Fontana et al., 2011). The different morphologies of portlandite seen in **Figure 7** can also be justified by differences in phase distribution. When hydrated in 1g, the severe bleeding caused by the high w/c reduces the availability of calcium in the pore solution, as well as reduces room for crystal growth. As a result, portlandite is more frequently distributed throughout the matrix and grows filling up the tortuous pore space (Gallucci and Scrivener, 2007). In µg, the crystals have a rather elongated, plate-like morphology, which is in line with the preferential crystal growth found by Harutyunyan and coworkers (Harutyunyan et al., 2009). The higher effective w/c and the larger porosity provide the dissolved species with more freedom to develop a purely chemical diffusion-driven microstructure, without complications from the physical effects of gravity. Also shown by Harutyunyan et al. (2009), portlandite has very specific crystallographic growth preferences that are accordant with the crystals shown in the µg image, **Figure 7**.

FIGURE 6 | SEM image taken from the top surface of a C3S paste hydrated under 1 g environment for 56 days. The large, plate-like portlandite crystals sitting on the C-S-H matrix can reach over 1 mm of diameter and result from bleeding effect. The scale bar at the bottom of the micrograph shows 1 mm.

Smaller crystallites of well-defined hexagonal shape, and covered by C-S-H, were also found in pastes hydrated in ground and space (**Figures 8**, **9**). Characterized by the hexagonal basal plane, the euhedral crystallites are smaller in size than typical portlandite. Their smaller size indicates that they likely precipitate at later stages of the hydration period (Boistelle and Astier, 1988; Glasser, 2001) after the skeleton of the paste is formed. Prismatic crystallites were oftentimes found in samples hydrated in µg, whereas plate-like crystallites of similar basal plane size were rarely found in 1g samples (**Figure 8**).

## Image Analysis

The typical greyscale histogram seen throughout the images taken of the samples is seen on the left side of **Figure 10**. There is minimal anhydrous material observed both visually and in the greyscale histogram of polished samples images. This is indicative of advanced degree of hydration as the samples age was 152 days. The slight bump around the greyscale value of 200 would be the minimal anhydrous material. The right side of **Figure 10** shows the cumulative greyscale histogram used with the overflow method to determine the greyscale value for thresholding.

As previously discussed, there is a visual difference between the porosity of the ground and space samples (**Figure 7**) that was confirmed by the image analysis. The calculated porosity by the area fraction shows porosities for the ground and space samples of 48.1 and 71.7 percent, respectively (**Table 2**). These

FIGURE 9 | The micrographs show CH crystallite hydrated in 1 g (left), resulting in a completely misoriented C-S-H covering its surface, while the CH crystallite formed in µg (right) is covered by an oriented C-S-H. Note the difference in the scale bar, which also highlights that the crystallites are bigger when hydrated under µg condition.

values show that the porosity in the space sample is over 20 percent higher than in the ground sample. The lower porosity in the ground sample could be partially attributed to the bleeding effect, which would lower the effective w/c ratio of the sample and therefore the porosity (Han and Wang, 2016). It is important to keep in mind that high values of porosity were

TABLE 2 | Summary of measured porosity, pore size (D50), and densities of ground and space samples.


expected due to the samples having a w/c of 2.0. This analysis further confirms the effect that gravity has on the densities of the samples.

## MIP

The results obtained from MIP, summarized in **Table 2**, support the SEM observations and are in good agreement with the image analysis. As expected, the porosity [i.e., the network of interconnected pores (Aligizan, 2006)] is over 20 percent higher in the space sample than in the ground control. It should be noted that the air bubbles (voids) (**Figure 4**) are not included in the measurements of interconnected porosity.

By comparison, the samples hydrated on Earth and space not only differ in terms of the total porosity and density (**Table 2**), but also in terms of pore size distribution. The pore size distribution in cement pastes strongly affects permeability, ion diffusivity (Halamickova et al., 1995; Kim et al., 2012), conductivity (Rajabipour et al., 2007), and mechanical properties (Zhang and Wang, 2016) of cement pastes. As shown in **Figure 11**, the ground sample develops well-distributed pore sizes, ranging between 15 and 10,000 nm. The sample hydrated in space, contrarily, develops over 50 percent of its pores between 6,000 nm and 10,000 nm, which is characterized by the steep slope (**Figure 9**). The calculated D50s (i.e., the pore diameter (nm) at which 50% of the pores are smaller than) are summarized in **Table 2**.

FIGURE 11 | Pore size (diameter) distribution of C3S paste hydrated in 1 g (black) and µg (red). The dashed line indicates the size limit between gel pores (<10 nm) and capillary pores (10 nm < capillary pore < 10,000 nm). The paste hydrated in 1 g condition develops well-distributed pore sizes, whereas the paste hydrated in µg condition develops 50 percent of the pores larger than 6,015 nm (D50). This can be seen graphically by the steep slope at the range of 6,000–10,000 nm.

Also shown in **Figure 11**, the curve representing the pore size distribution of the ground control ends abruptly on a steep slope, whereas the curve representing the space sample seems to be slowly plateauing. This suggests that percentage of pores smaller than 3 nm is likely to be greater in the ground sample than the in the space one. The inability of mercury to penetrate into pores smaller than 3 nm (Aligizan, 2006) limits the measurements of pore size distribution but still support the differences found in the skeleton densities (**Table 2**). Additionally, it needs to be mentioned here that MIP measures the entry pore diameter only.

Another important feature pointed out by MIP is the bulk density, which also differs significantly between the space and ground control samples. ASTM D3766-08 (2012) defines bulk density as one including the skeleton of the paste, connected pores, non-connected pores, and air bubbles. The ground control under the influence of a gravity-driven compaction becomes denser, i.e., each gram of hydrated paste should occupy a

predictable volume, which can be theoretically calculated (Young and Hansen, 1987) or modeled (Lothenbach et al., 2011). Both the empirical equations and model for volume changes would not apply for cements hydrated in µg. In the absence of gravity, cement-based materials should occupy a larger volume, given the lower bulk density (**Table 2**), as verified by visual comparison of the 1g and µg samples. The skeleton density also differs between the space sample and ground control. In this case, the volume of the sample excludes the connected porosity from the volume calculations (ASTM D3766-08, 2012), and most likely the air bubbles, suggesting a possible variability of the C-S-H hydrated in µg.

## TGA

While it is possible that gravity impacts reaction rates, the results obtained here suggest similar reaction rates between ground and space samples. **Figure 12** shows that the mass of portlandite produced is essentially unaffected by the gravity level, despite the discussed morphological and spatial differences. The analysis of the collected data to quantify the amount of portlandite was performed as suggested by Kim and Olek (Kim and Olek, 2012). The portlandite content in space and ground samples are 35.3 and 37.2 percent by mass, respectively. This difference of <2 percent is attributed to sampling and/or instrumentation errors.

The degree of hydration (α) can also be inferred from the TGA curves through Equation 1 (Zhang and Scherer, 2011):

$$\alpha = \frac{W\_n}{n} - LOI \tag{1}$$

where W<sup>n</sup> is the normalized weight loss of the sample, n is the mass of non-evaporable water of the fully hydrated paste, and LOI is the loss on ignition or the normalized weight loss of the anhydrous cement. The non-evaporable water for C3S paste is a fixed value of 0.21 determined by Young and Hansen (Young and Hansen, 1987), whereas the LOI, measured according to ASTM C114 (ASTM C114 – 18., 2018) is 0.9 percent. That gives an 88 percent degree of hydration for both ground and space samples.

## XRD

Qualitative XRD data, shown in **Figure 13**, points out the presence of two crystalline phases, identified as portlandite and calcite. Firstly, the diffraction pattern confirms that µg tends to change the growth kinetics of portlandite, resulting in different crystal morphology. More specifically, the lengthto-width or aspect ratio of portlandite is reduced in space samples in comparison to the ground control. The diffraction pattern presents several discrepancies among the intensities most pronounced in the indexed portlandite peaks. The main difference observed is at the peak corresponding to the (001) basal plane (2θ =18.066◦ ) of calcium hydroxide. According to the Powder Diffraction File no. 01-076-0571 (Kabekkodu S. (ed)., 2018), the most intense peak (100% intensity) should be that of (011) (2θ =34.101◦ ). The presented spectrum (**Figure 13**) is subject to preferred orientation effects, or the tendency of dishlike crystals to lie on the basal plane, parallel to the surface, rather than being randomly oriented in the prepared sample (Grattan-Bellew, 1975). For portlandite, this effect erroneously enhances the relative intensity of the diffracted (001) hexagonal bases, while decreasing the intensity of peaks corresponding to other planes, such as the lateral facets. Despite the error, the diffraction pattern coupled with SEM observations highlight the influence of gravity level on crystal morphology.

Secondly, the peaks resulting from the presence of calcite indicates that the samples experienced carbonation. It is believed that carbonation occurred during sample preparation, since the pouches were sealed and contained limited air. Furthermore, during the extensive SEM analysis on fractured samples, crystals resembling the possible forms of calcite crystals (Mason and Berry, 1968) were not found.

Another important observation inferred from the diffraction pattern is that the peak positions (2θ) corresponding to portlandite in the space samples have not changed relative to the ground control. As such, it implies that the atom positions of the unit cell are the same as those measured by Busing and Levi (Busing and Levy, 1957). The results (**Figure 13**) are also in line with the NaCl crystal grown aboard the ISS (Fontana et al., 2011), which presented morphological differences from those typically grown on Earth, while maintaining the lattice parameters.

## CONCLUSIONS

This research presents an experimental comparison between high water-to-cement ratio C3S pastes hydrated under terrestrial gravity and microgravity conditions. The results revealed that microgravity alters several features of the microstructure of high water-to-cement ratio C3S paste, including, but not limited to the following:


image analysis and MIP. Moreover, the pore diameters for µg are one order of magnitude larger than the ground control.


The results presented here provide a first comparison of cement samples processed on the ground and in a microgravity environment. Lack of physical factors, including buoyancy, sedimentation, and fluid convection resulted in samples with uniform phase distribution and distinct morphologies. Despite using a very high w/c, the observations made here call attention to aspects worthy of being explored in future studies.

## AUTHOR CONTRIBUTIONS

JM collected and interpreted most of the data, as well as wrote the manuscript. PC was responsible for image analysis and helped substantially with writing the manuscript. RW contributed by collecting several XRD diffraction patterns and helped analyzing and interpreting the data. RG was present during SEM investigation, provided valuable mentoring and insights about his experience with microgravity research. AR is the principal investigator, advisor, and was awarded with the NASA grant.

## FUNDING

All the funding was provided by NASA.

## ACKNOWLEDGMENTS

The authors gratefully acknowledge the financial support from the National Aeronautics and Space Administration (NASA)-Grant No. NNX17AC48G, as well as use of MSFC's EM31 Materials Diagnostics laboratory and precious insights from David Donovan, Karen Stephens, and Pam Denkins. Appreciation is expressed to Reviewers of the manuscript, Dr. Farshad Rajabipour, Dr. Barry Scheetz for their feedback during manuscript writing, and the Science Concept Review committee: Jan Olek, Maria Juenger, Kim Kurtis, Jason Ideker, and Sulapha Peethamparan for their valuable suggestions in research planning stages. The authors also express gratitude for collaboration with NIST, in particular Dale Bentz and Jeff Bullard. Thanks to the Leidos team, particularly Micah Johnson, Cole Nelson, and Kile Mouch, for the pre-launch preparation of the samples. Lastly, the authors express immense gratitude for the astronauts of Expedition 56: Drew Fuestel, Serena M. Auñón-Chancellor, Ricky Arnold, Alexander Gerst, Sergey Prokopyev, and Oleg Artemyev, who executed the experiments aboard the ISS.

## REFERENCES


Zhang, Z., and Wang, H. (2016). The Pore characteristics of geopolymer Foam concrete and Their impact on the compressive strength and Modulus. Front. Mater. Struct. Mater. 3, 1–10. doi: 10.3389/fmats.2016.00038

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Moraes Neves, Collins, Wilkerson, Grugel and Radlinska. This ´ is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Effect of Nano Alumina on Compressive Strength and Microstructure of High Volume Slag and Slag-Fly Ash Blended Pastes

### Faiz Uddin Ahmed Shaikh\* and Anwar Hosan

*School of Civil and Mechanical Engineering, Curtin University, Perth, WA, Australia*

This paper presents the effect of nano alumina (NA) on compressive strength and microstructure of cement paste containing high volume blast furnace slag (HVBFS) contents of 70, 80, and 90% as partial replacement of cement and combined blast furnace slag (BFS) and class F fly ash (FA) contents of 70 and 80% as partial replacement of cement. FA is used at 30% as partial replacement of BFS. NA contents are varied from 1 to 4% as partial replacement of BFS and BFS-FA. Results show that the addition of NA improves the compressive strength of high volume BFS and BFS-FA pastes by 2 to 16%. The compressive strength of paste containing 69% BFS, 30% cement, and 1% NA exceeded the compressive strength of control cement paste while the compressive strength of paste containing 77% BFS, 20% cement, and 3% NA is 1% lower than control cement paste. NA significantly reduced the large capillary pores of >0.1 microns of high volume BFS and BFS-FA pastes. No evidence of reduction of Ca(OH)<sup>2</sup> in high volume BFS pastes is observed due to addition of NA, however, in high volume BFS-FA paste the Ca(OH)<sup>2</sup> is reduced due to addition of NA. Increase in intensity peaks of CAH, Ettringite and CSH in X-ray diffraction analysis is observed in high volume BFS and BFS-FA pastes due to addition of NA, which coincides with the observed more dense microstructure of high volume BFS and BFS-FA pastes containing NA than those without NA.

Keywords: nano alumina, blast furnace slag, fly ash, compressive strength, microstructure

## INTRODUCTION

Concrete is the most widely used construction materials in the world. Its demand continues due to rapid urbanization and population growth in the world. As a result the demand of cement, the key ingredient of concrete, also increases. However, the manufacturing of cement is energy intensive and contributes about 5–7% of total CO<sup>2</sup> release in to the atmosphere globally (Benhelal et al., 2013). Therefore, there is worldwide initiative to reduce the use of cement in concrete by partial replacement of cement using various industrial by-products such as fly ash, blast furnace slag (BFS), silica fume, rice husk ash, sugarcane bagasse ash, etc. Significant amount of research have been devoted worldwide to investigate various properties of concrete containing above byproducts. Very good level of understanding of their use in concrete at low volume as well as at high volume already exists, which result their use in concrete as partial replacement of cement in many countries. However, their use in high volume in concrete often results in reduction in early age mechanical and durability properties due to slow pozzolanic reaction of these materials with

#### Edited by:

*Maria Juenger, University of Texas at Austin, United States*

#### Reviewed by:

*Denvid Lau, City University of Hong Kong, Hong Kong Ali Behnood, Purdue University, United States*

> \*Correspondence: *Faiz Uddin Ahmed Shaikh s.ahmed@curtin.edu.au*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *20 December 2018* Accepted: *10 April 2019* Published: *30 April 2019*

#### Citation:

*Shaikh FUA and Hosan A (2019) Effect of Nano Alumina on Compressive Strength and Microstructure of High Volume Slag and Slag-Fly Ash Blended Pastes. Front. Mater. 6:90. doi: 10.3389/fmats.2019.00090* cement hydration product (Oner and Akyuz, 2007; Aghaeipour and Madhkhan, 2017). Various ultrafine and nano materials are used to address the above limitation of slow pozzolanic reaction at early stage of concrete. In a number of studies the effect of various nano materials e.g., nano silica, nano CaCO3, nano TiO2, nano alumina, carbon nano tube, etc. on the properties of concrete containing fly ash at various volume fractions have been evaluated and promising results in terms of improvement in mechanical and durability properties of concrete containing fly ash are reported (Sato and Diallo, 2010; Hou et al., 2013; Shaikh and Supit, 2014, 2015; Shaikh et al., 2014).

Blast furnace slag is the by-product of iron/steel industry and is widely used in concrete as partial replacement of cement.


Its use up to 60–70% as partial replacement of cement in concrete is researched (Elchalakani et al., 2014). It has been reported that above this limit the mechanical and durability properties are significantly affected (Oner and Akyuz, 2007). Despite the use of various nano materials in concrete containing high volume fly ash, not enough research is reported on the use of nano materials in concrete containing high volume slag above 60%, especially nano alumina which exhibited significant improvement in modulus of elasticity of concrete. Nano alumina, similar to other nano materials, provides synergistic effects in cement hydration process through seeding and filler effects and pozzolanic reaction. In a number of studies the effect of nano alumina on mechanical properties of concrete and cement composite is evaluated. Li et al. (2006) investigated the effect of various nano alumina contents of 3–7% on compressive strength and elastic modulus of cement mortar and reported about 143% improvement in 28 days elastic modulus of cement mortar due to addition of 5% nano alumina due to densification of cement matrix around interfacial transition zone of aggregates in mortar. However, at the same nano alumina content the improvement in compressive strength was not that great. In a separate study, Nazari and Riahi (2011) however, reported significant improvement (about 44%) in compressive strength of cement paste due to addition of 4% nano alumina. Reduction in harmful pores in cement paste due to addition of nano alumina is also observed in their study. On the other hand, Barbhuiya et al. (2014) reported no improvement in 7 days compressive strength of cement paste due to addition of nano alumina despite formation of dense microstructure of cement paste containing nano alumina. Improvement in compressive strength of cement pastes due to addition of nano alumina is also reported by others (Heikal et al., 2015; Gowda et al., 2017). The variation of results on properties of concrete due to addition of NA could be due to their different particle sizes e.g., the size of NA was ∼150nm in the study by Li et al. (2006) whereas Nazari and Riahi (2011) use much smaller size NA particles of about 15 nm. On the other hand, the size of NA was 27–43 nm in Barbhuiya et al. (2014) study.

Mohseni and Tsavdaridis (2016) reported improvement in compressive strength and durability of fly ash mortar containing 25% fly ash due to addition of nano alumina. Reduction in gel and medium capillary pores of fly ash mortar is reported in this study due to addition of nano alumina. So far only one study reported by Nazari and Riahi (2011) evaluated the effect of nano alumina on mechanical and microstructure of self-compacting concrete containing various BFS contents of up to 60% as partial replacement of cement. While nano alumina exhibited improvement in mechanical and durability properties of cement paste and mortars and those containing fly ash and BFS, no study so far reported its effect on the concrete containing high volume slag content above 70% as partial replacement of cement. This paper presents the effect of various nano alumina contents on compressive strength and microstructure of cement pastes containing high volume blast furnace slag (HVBFS) and high volume blast furnace slag and fly ash (HVBFS-FA) as partial replacement of cement.

## EXPERIMENTAL PROGRAM

This study is consisted of five parts. In the first three parts, three high volume slag pastes containing 70, 80, and 90% BFS by weight as partial replacement of ordinary Portland cement (OPC) are

considered. In the fourth and fifth parts, the effect of two hybrid combinations of BFS and class F fly ash (FA) at total contents of 70 and 80% by weight are considered, where FA is used at 30% by weight of BFS. In all parts the effect of nano alumina (NA) content of 1–4% by weight on the high volume BFS pastes and high volume BFS-FA pastes are evaluated. Therefore, total 26 mixes are considered in this study.

## MATERIALS, MIX PROPORTIONS, AND CASTING

The cement used in this study was OPC. The fly ash (FA) was class F fly ash supplied by Eraring power station of New South Wales in Australia. The ground granulated BFS was supplied by BGC cement of Australia. The nano alumina (NA) powder was purchased from Nanostructured and Amorphous Materials, Inc. of USA. **Table 1** shows the physical and chemical properties of OPC, FA, BFS, and NA, while **Figures 1**, **2** show the X-ray diffraction (XRD) analysis of FA, BFS and NA. It can be seen by comparing the results that NA was less amorphous compared to FA and BFS. Among FA and BFS, BFS is more amorphous with amorphous content of 95.7% compared to 67.8% amorphous content of FA based on quantitative XRD analysis. Detail mix

TABLE 2 | Mix proportions of high volume slag and high volume slag-fly ash blend cement pastes containing nano alumina (NA).


*Superplasticizer was used as wt.% of binder (OPC* + *BFS* + *FA* + *NA).*

FIGURE 3 | Dispersion of nano alumina in water contining superplasticizer solution in ultrasonication method.

proportions of high volume BFS pastes, high volume BFS-FA pastes and those containing NA are shown in **Table 2**. It can be seen that water/binder ratio of all mixes were kept constant at 0.4. In the case of mixes containing NA, a naphthalene sulfonate based superplasticizer was used to maintain the workability of pastes. All pastes were mixed in Hobart mixer, where cement, FA and BFS were mixed first for ∼3 min. Water was added and mixed for further 3 min. In the case of high volume BFS and high volume BFS-FA pastes containing NA, the NA powder was first ultrasonically dispersed in water containing superplasticizer using ultrasonic mixer for 30 min shown in **Figure 3** and then added with remaining water during mixing. The paste samples of 50 mm cubes were cast and demoulded after 24 h. The specimens were cured in water at room temperature for 28 days.

## TESTING METHODS

Compressive strength of the specimens were tested according to ASTM C109 (2012) using a loading rate of 0.7 MPa/s. In order to observe the changes in reaction phases in cement pastes containing high volume BFS and BFS-FA blends due to inclusion of NS thermogravimetric analysis (TGA), X-ray diffraction (XRD) analysis, Mercury intrusion porosimetry (MIP), scanning electron microscopy (SEM), and energy dispersive X-ray spectroscopy (EDS) were done on selected samples.

For the XRD analysis, the samples were measured with a D8 Advance Diffractometer (Bruker-AXS) using copper radiation and a Lynx Eye position sensitive detector. The diffractometer scanned the samples from 7◦ to 70◦ (2θ) in steps of 0.015◦ at a scanning rate of 0.5◦ / min. XRD patterns were obtained using Cu Ka lines (k = 1.5406 Å). A knife edge collimator was fitted to reduce air scatter. SEM analyses were performed using a Zeiss EVO 40XVP microscope equipped with an energy dispersive Xray analyser. The thermal stability of the samples was studied by thermogravimetric analysis (TGA). A Mettler Toledo TGA one star system analyser was used for all these measurements. Samples weighing 25 mg were placed in an alumina crucible and the tests were carried out in an Argon atmosphere at a heating rate of 10◦C/min from 25 to 1,000◦C.

Mercury intrusion porosimetry (MIP) was used to measure the pore volume and pore size distribution of pastes samples. The pore diameter and intruded mercury volume were recorded at each pressure point over a pressure range of 0.0083–207 MPa. The pressures values were converted into equivalent pore diameters using the Washburn expression 1921, as expressed in Equation (1):

$$d = -\frac{4\nu\cos\theta}{\rho} \tag{1}$$

where, d is the pore diameter (µm), γ is the surface tension (mN/m), θ is the contact angle between mercury and the pore wall (◦ ), and P is the net pressure across the mercury meniscus at the time of the cumulative intrusion measurement (MPa).

## RESULTS AND DISCUSSION

## Compressive Strength

The effects of various NA contents of 1–4% on compressive strength of high volume BFS pastes are shown in **Figure 4**. **Figure 4A** shows that due to addition of NA the compressive strength of high volume BFS paste containing 70% BFS is higher than that without NA. The compressive strength is increased by 2–16% due to addition of NA. In the case of high volume BFS pastes containing 80 and 90% BFS the improvement in compressive strength due to addition of NA is not better than 70% BFS content. It can be seen in **Figures 4B,C** about 2–8 and 1% improvement in compressive strength due to addition of various amount of NA in pastes containing 80 and 90% BFS, respectively. While the compressive strength of high volume BFS pastes containing 70% and NA is higher than that of control OPC paste, however, the compressive strength of paste containing 80 and 90% BFS is lower than OPC paste despite the addition of NA. This can be due to significantly lower level of Portlandite (Ca(OH)2) generation in these mixes due to very high volume of OPC replacement by BFS. It can also be seen no trend in compressive strength improvement with increase in NA contents in all high volume BFS pastes, however, beyond 3% no improvement in compressive strength is observed presumably due to agglomeration of NA as it has been reported by several researchers. Due to higher van der Waal's forces, the nano alumina particles in water causes agglomeration and hence,

adversely affects their uniform dispersion in the matrix (Senff et al., 2009; Quercia et al., 2012).

The effect of NA on compressive strength of pastes containing combined BFS and FA contents of 70 and 80% is shown in **Figures 4D,E**. It can be seen that the addition of NA improves the compressive strength by up to 9% of paste containing total BFS and FA content of 70%, which is slightly lower than paste containing 70% BFS. In the case of combined BFS and FA

content of 80% containing NA up to about 16% improvement in compressive strength is observed which is higher than those containing 80% BFS. However, the compressive strength of these two HVBFS-FA pastes for all NA contents is still lower than that of control OPC paste. The improved compressive strength due to addition of NA in pastes containing combined BFS and FA than the pastes containing BFS can be interpreted as the better dispersion of NA in those mixes due to spherical shape of the FA particles.

## Microstructure

It can be seen in above compressive strength results that the high volume BFS paste containing 70% BFS and NA exhibited higher compressive strength than that of control OPC paste. High volume BFS paste containing 80% BFS and combined BFS and FA content of 70% showed about 1 and 7%, respectively, lower compressive strength then that of OPC paste. These mixes contain significantly less OPC hence much lower carbon footprint of concrete made using these blends. Therefore, in the microstructural analysis the changes of microstructures in terms of their pore sizes and volumes, conversion of Portlandite to secondary calcium silicate hydrate (CSH)/calcium aluminate hydrate (CAH) through pozzolanic reaction and formation of new mineral phases are studied due to addition of NA.

## Mercury Intrusion Porosimetry (MIP)

**Figure 5** shows the cumulative pore volume and distribution of various pore sizes of high volume BFS and combined BFS and FA pastes with and without NA. It can be seen in **Figure 5A** that the volume of pores corresponding to pore diameter >0.1 microns increases in cement paste due to addition of high volume BFS contents of 70 and 80% as partial replacement of OPC. However, the volume of medium capillary pores and gel pores corresponding to pore diameters 0.05–0.01 micron and <0.01 microns, respectively, are decreased due to addition of 70% BFS compared to control OPC paste. In the case of paste containing 80% BFS the volume of medium capillary pores and gel pores are higher than OPC paste. By comparing the pastes containing two BFS contents, it can be seen that at 70% BFS content the total pore volume of all pores is less than at 80% BFS content, which might be the reason for observed higher compressive strength in the former paste than the latter as large pores might have caused stress concentration. On the other hand, in the paste containing combined BFS and FA with a total content of 70% the pore volume of larger capillary pores is reduced significantly compared to both high volume BFS pastes and control OPC paste. The distribution of pores having different diameters in

all pastes is shown in **Figure 5B**. The region under the curves represent the concentration of pores and the peak point of the curve represents the critical pore diameter. By comparing different curves its can be seen that the maximum concentration of pores in OPC paste is in 0.03–0.2 microns range which is in lower side of large capillary pores. However, in the case of pastes containing 70 and 80% BFS and combined BFS and FA content of 70% the peak of their curves is shifted toward smaller pore diameter and they fall within medium capillary pores. This indicates that due to addition of BFS and combined BFS and FA the concentration of pores is shifted from large capillary pores to medium capillary pores.

The cumulative pore volume and pore size distribution of high volume BFS and combined BFS and FA pastes containing NA are also shown in **Figure 5**. It can be seen in **Figure 5A** that due to addition of NA the volume of pores corresponding to pore size >0.1 microns is significantly decreased in all high volume BFS and combined BFS and FA pastes. Surprisingly it can also be seen that volume of gel pores and medium capillary pores is increased due to addition of NA in high volume paste containing 70% BFS. The cumulative pore volume of paste containing 80% BFS is reduced from 0.164 to 0.155 cc/g due to addition of 3% NA. In the case of paste containing combined BFS and FA content of 70% the cumulative pore volume is reduced significantly from 0.155 to 0.14 cc/g due to addition of 1% NA. The formation of additional CSH and CAH gels due to enhanced pozzolanic reaction of NA with Portlandite in those high volume BFS and combined BFS and FA pastes could be the reasons for such reduction in volume of large capillary pores in those pastes. However, by looking into the pore size distribution shown in **Figure 5B** it can be seen that the concentration of pores is not significantly changed in high volume BFS and combined BFS and FA pastes due to addition of NA, but the volume of pores in the concentration of pore sizes in each paste containing NA is reduced.

## Thermogravimetric Analysis (TGA)

TGA cures of high volume BFS and BFS-FA pastes and those containing NA are shown in **Figure 6A**. TGA curve shows the mass of a substance under control environment as a function of temperature. It has been reported that mass loss between 420 and 540◦C corresponds to the dehydration of calcium hydroxide (Keattch and Dollimore, 1975). The amount of calcium hydroxide (CH) can be quantified according to Taylor (1990) formula shown as follows in Equation (2):

$$CH\left(\%\right) = WL\_{CH}\left(\%\right) \frac{MW\_{CH}}{MW\_{H2O}} \tag{2}$$

Where, WLCH is the weight loss during the dehydration of CH between 420 and 540◦C as percentage of the ignited weight (%); MWCH is the molecular weight of CH; MWH2O is the molecular weight of H2O.

The calculated CH contents are shown in **Figure 6B** and it can be seen that CH contents of all high volume BFS and BFS-FA pastes and those containing NA are lower than that of OPC paste. This clearly indicate that due to pozzolanic reaction of SiO2 in BFS and FA with CH in those pastes CSH is formed.

paste. (B) 69BFS+1NA paste. (C) 80BFS paste. (D) 77BFS+3NA paste. (E) 49BFS21FA paste. (F) 48BFS21FA+1NA paste.

In **Figure 6B** it can also be seen slightly higher amount of CH content in high volume BFS pastes containing 77%BFS + 3%NA and 69%BFS + 1%NA pastes than that containing 80% BFS and 70%BFS pastes, respectively. This indicates that Al2O3 in NA did not produce any CSH in the matrix rather formed CAH which is confirmed in XRD and EDS analysis discussed in latter sections. Surprisingly, the CH content is much lower in paste containing 49%BFS + 20%FA + 1%NA then those of pastes containing 70% and 80% BFS. This can be interpreted due to reaction of SiO<sup>2</sup> of FA with CH in that paste.

## X-Ray Diffraction (XRD) Analysis

XRD patterns of high volume BFS and combined BFS and FA pastes with and without NA are shown in **Figure 7**. The horizontal scale is the diffraction angels measured in degrees and the vertical scale is the peaks height of the intensity of the diffractions measured in pulses/s. The diffraction spectra analysis of high volume BFS and combined BFA and FA pastes with and without NA indicate the predominance of calcium hydroxide (CH), calcium silicate hydrate (CSH), calcium aluminate hydrate (CAH), Ettringite (E), and Quartz (Q). By comparing the CH peaks of pastes with and without NA it can be seen that in most occasions intensity of CH peaks is not changed due to addition of NA except at 2θ angles 34.5◦ and 47◦ , where intensity of CH is reduced in pastes containing NA. Comparison also shows the formation of Ettringite in high volume BFS pastes containing NA at 2θ angles 43.5◦ and increase in intensity of Ettringite peaks at 2θ angles 19◦ , 25.6◦ , and 35◦ . The intensity of CAH peak at 2θ angle of 10.8◦ is also increased in high volume BFS and combined BFA and FA pastes due to addition of NA. In the case of high volume BFS-FA paste new mineral peaks of Mullite (M) and Gehlenite (G) can be seen at 2θ angles 16.5◦ and 20.8◦ , respectively, with increase of Gehlenite peak due to addition of NA.

## Scanning Electron Microscopy (SEM) and Energy Dispersive X-Ray Spectroscopy (EDS) Analysis

SEM images and EDS spectra of high volume BFS and high volume BFS-FA pastes with and without NA are shown in **Figure 8**. It can be seen by comparing **Figure 8B** with **Figure 8A** that the microstructure of high volume BFS paste containing 69% BFS+1% NA is much denser than paste containing 70% BFS. By comparing both figures it can also be revealed less black spots indicative of pores/voids in high volume BFS paste containing NA than high volume BFS paste indicating formation of additional hydration products such as CSH/CAH as shown in the figure. Similar, improvement can also be seen in the high volume BFS paste containing 80% BFS and that containing NA in **Figures 8C,D** where paste containing 77% BFS and 3% NA in **Figure 8D** exhibited much denser microstructure than high volume BFS paste containing 80% BFS in **Figure 8C**. The EDS spectra in the same figure shows the formation of more CSH/CAH in the high volume BFS pastes containing NA than that of high volume BFS pastes without NA. The similar densification of microstructure of high volume BFS-FA paste due to addition of NA can also be seen by comparing **Figure 8F** with **Figure 8E**, where the peak of silica in EDS is much higher in paste containing NA than without NA indicating formation of more CSH/CAH in the former than the latter. The higher peak of silica in EDS is also evident in high volume BFS pastes containing NA than those without NA.

## CONCLUSIONS

The effects of various nano alumina contents on compressive strength and microstructure of cement pastes containing HVBFS and high volume blast furnace slag-fly ash combination are

## REFERENCES

reported. Based on limited study the following conclusion can be drawn:


## AUTHOR CONTRIBUTIONS

FS planned the experimental program, analyzed results, and wrote the manuscript. AH cast and tested all specimens in mechanical and microstructural studies and compiled and analyzed the results.

ASTM C109 (2012). Standard Test Method for Compressive Strength of Hydraulic Cement Mortars Using 50mm cube Specimens. West Conshohocken, PA: ASTM.

Barbhuiya, S., Mukherjee, S., and Nikraz, H. (2014). Effects of nano-Al2O3 on early age microstructural properties of cement paste. Constr. Build. Mater. 52, 189–193. doi: 10.1016/j.conbuildmat.2013.11.010

Aghaeipour, A., and Madhkhan, M. (2017). Effect of ground granulated blast furnace slag (GGBFS) on RCCP durability. Constr. Build. Mater. 141, 533–541. doi: 10.1016/j.conbuildmat.2017.03.019


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Shaikh and Hosan. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

## Characteristics of Ferrite-Rich Portland Cement: Comparison With Ordinary Portland Cement

Yogarajah Elakneswaran<sup>1</sup> \*, Natsumi Noguchi <sup>1</sup> , Kazuki Matumoto<sup>1</sup> , Yuka Morinaga<sup>1</sup> , Takashi Chabayashi <sup>2</sup> , Hiroyoshi Kato<sup>2</sup> and Toyoharu Nawa<sup>3</sup>

*<sup>1</sup> Division of Sustainable Resources Engineering, Faculty of Engineering, Hokkaido University, Sapporo, Japan, <sup>2</sup> Cement Business Division, Development Department, Tokuyama Corporation, Yamaguchi, Japan, <sup>3</sup> Hokkaido University, Sapporo, Japan*

#### Edited by:

*Miroslav Komljenovic, University of Belgrade, Serbia*

#### Reviewed by:

*Frantisek Soukal, Brno University of Technology, Czechia Tongbo Sui, Sinoma International Engineering Co., Ltd., China*

\*Correspondence:

*Yogarajah Elakneswaran elakneswaran@eng.hokudai.ac.jp*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *11 January 2019* Accepted: *15 April 2019* Published: *01 May 2019*

#### Citation:

*Elakneswaran Y, Noguchi N, Matumoto K, Morinaga Y, Chabayashi T, Kato H and Nawa T (2019) Characteristics of Ferrite-Rich Portland Cement: Comparison With Ordinary Portland Cement. Front. Mater. 6:97. doi: 10.3389/fmats.2019.00097*

The cement industry is an energy-intensive industry, and improving the energy efficiency of cement has become necessary to reduce its carbon footprint and to compete in the global market. Clinker production consumes more than 90% of the total energy used in the cement industry. Therefore, a reduction in the burning temperature of the cement clinker can reduce the energy consumption; however, it alters the mineralogy of the clinker composition. Ferrite-rich Portland cement can be produced by lowering the burning temperature by ∼100◦C (i.e., at 1,350◦C), which can reduce the energy consumption by ∼5% in comparison with ordinary Portland cement (OPC) clinker. In this study, the hydration reaction and properties of the ferrite-rich Portland cement were examined by experimental techniques and thermodynamic modeling approach, and the results were compared with that of OPC. The produced ferrite-rich cement has almost twice the amount of ferrite phase and half the amount of belite phase contents present in the OPC. The hydration reaction and the composition of hydrates were studied by the X-ray diffraction (XRD)/Rietveld analysis and thermogravimetry (TG) and differential thermal analysis (DTA). The different proportions of the ferrite and belite phases in ferrite-rich cement change their hydration reaction from that of the OPC, but not the total hydration of cement. The XRD results reveal similar phases in both the cements, and the analysis could not identify the new phases formed in the ferrite-rich cement. An equal degree of hydration and quantified hydrates at the early age results in almost identical initial and final setting times in both the cements. The ferrite-rich cement demonstrates a high early strength and relatively slower strength development; however, it can develop adequate strength at 28 days. The thermodynamic model predicts the hydration of ferrite-rich cement and shows comparatively high amount of Fe-containing phases, mainly Fe-ettringite and Fe-siliceous hydrogarnet. Model predictions of the hydrates compositions agreed with the experimental results, and a relationship between the predicted total porosity and the measured compressive strength was derived.

Keywords: Fe-containing phases, hydration, thermodynamic modeling, X-ray diffraction, compressive strength, setting time

## INTRODUCTION

The cement industry is a high-energy consuming sector, and the production of ordinary Portland cement (OPC) is responsible for ∼5% of the total anthropogenic greenhouse gas emission by limestone calcination (Imbabi et al., 2012; Amato, 2013). Reduction of CO<sup>2</sup> emission associated with the cement manufacturing process without affecting the cement demand is a very challenging task for the cement and concrete industry. Investigations on novel cement systems and low-carbon concretes are thus important to reduce the environmental impact of cement. A number of novel cement systems and low-carbon concretes are currently being developed as alternatives for the OPC (Gartner, 2004; Gartner and Macphee, 2011). Replacement of clinker or OPC by supplementary cementitious materials can help to mitigate the carbon footprint of the OPC (Lothenbach et al., 2011), and various alternative clinkers have been proposed to reduce the limestone amount in the raw mix of cement production (Morin et al., 2017).

Over the years, cement industry has been reducing the CO<sup>2</sup> emission associated with the heating process by improving the energy efficiency, increasing the use of alternative fuels, and deploying renewable energy sources (Morin et al., 2011; Folliet et al., 2017). However, it needs further energy conservation technologies to reduce its carbon footprint. The conventional OPC clinker is burnt at a very high temperature of about 1,450◦C, and the heating process is the most energy-consuming part in cement manufacture. Therefore, decreasing the clinker burning temperature reduces both the energy consumption and CO<sup>2</sup> emission. Various technologies have been developed to reduce the energy consumption during the burning process, such as changing the mineral composition of the clinker or using alternative clinkers for OPC and using mineralizers to reduce the burning temperature (Yamashita and Tanaka, 2011). However, studies on the systematic adjustment of the mineral composition of clinker to reduce the burning temperature are limited.

Authors' previous study has reported that compared to the other components of the clinker, the ferrite phase contributes the most to the sintering of the clinker, and the contributions of the various components to clinker burnability is in the order, C4AF>C3A>C2S>C3S (Chabayashi et al., 2012, 2015). A low-burning-temperature clinker, which sinters at 1,350◦C,


which is ∼100◦C lesser than temperature required for the OPC clinker, was produced by increasing the C4AF and decreasing the C2S proportions while targeting a strength equal to that of the conventional OPC. It was successfully burnt in an actual kiln as well (Chabayashi et al., 2015). The cement manufactured using this low-burning-temperature clinker is called ferriterich Portland cement (FC), and it can reduce ∼5% of total CO<sup>2</sup> emission compared to the OPC during the clinkering process (Chabayashi et al., 2015). Furthermore, the amount of limestone required for cement production, calculated from mineral composition of clinker, proved that compared to the OPC, the FC would reduce ∼1.29% of CO<sup>2</sup> emission during the heating of limestone. Therefore, the overall CO<sup>2</sup> emission in cement manufacturing process can be decreased by 2.77% by replacing OPC with FC (Shinmi et al., 2017). The characteristics of FC are investigated in this study.

To the best of authors' knowledge, only a few studies have been reported on the production, hydration, property development, and performance of FC. It has been reported that use of FC delays the strength development compared to those with less ferrite and more aluminates (Gartner and Myers, 1993). Furthermore, chemical admixtures are necessary to enhance the hydration reaction of the ferrite phase (Gartner and Myers,

FIGURE 3 | Hydration degree of cements as a function of hydration time determined by XRD.

1993; Schwarz, 1995); this in turn contributes to later-age strength development. Poor crystalline Fe-containing phases could form during the hydration of the ferrite phase (Dilnesa et al., 2011, 2014), influencing the properties of the cement matrix. Difficulties in understanding the hydration reaction of the ferrite phase and the identification of Fe-containing phases in the hydration products limit the development and application of FC. Therefore, the objectives of this study are to understand and evaluate the hydration reaction and mechanical performance of FC and to compare the results with that of the conventional OPC. The X-ray diffraction (XRD)/Rietveld analysis and thermogravimetry (TG) and differential thermal analysis (DTA) were used to identify and quantify the clinker minerals and their hydration products as well as to estimate the amount of amorphous in the hydration products. The thermodynamic hydration model developed in authors' previous study (Elakneswaran et al., 2016) was applied to the FC to predict the hydration reaction and to estimate the hydration products. The modeling results were verified with the experimental results. Finally, a relationship between the measured compressive strength of the cement paste and predicted total porosity was derived.

## MATERIALS AND METHODS

## Clinker and Cement Production

The limestone, silica stone, coal ash, and iron source were used for the clinker production of both cements. The mixture of raw materials for OPC was first heated to 1,000◦C using electric furnace in 30 min and burnt at that temperature for another 30 min. Further, the temperature was increased to 1,470◦C in 47 min and kept constant for 90 min to limit the amount of free lime to <1%. The calcination of FC up to 1,000◦C follows the same procedure as that used for OPC; thereafter, the temperature is raised to 1,370◦C in 37 min and sustained for 90 min.1 The clinkers were rapidly cooled to room temperature in the air. Finally, gypsum (gypsum: bassanite = 1:1) was added to the produced clinkers in an amount equal to 2% of SO3, and the mixture was then ground using a ball mill to the target Blaine value of 3,200 ± 50 cm<sup>2</sup> g −1 to produce the cements. The Blaine specific surface area and the density based on JIS R 5201, the oxide composition by X-ray fluorescence (XRF), and the mineral composition by Bogue analysis were determined for the produced cements; the results are listed in **Table 1**.

### Sample Preparation

The cement was mixed with distilled water at a water to cement ratio of 0.5. The mixture was stirred manually until bleeding stopped. Further, it was cast into cylindrical molds and sealedcured at 20◦C. The samples that reached the predetermined curing time (1, 6, and 12 h and 1, 2, 3, 7, 14, and 28 days) were ground and immersed in acetone for 1 h to stop the hydration. Thereafter, the samples were removed from the acetone solution by suction filtration using an aspirator. Finally, the samples were kept in an oven at 40◦C until they reached a constant mass. The prepared samples were ground and powdered for XRD and TG measurements. The 2 × 2 × 8 cm paste specimens were prepared for compressive strength measurement using a mixer given in JIS R 5201.

## Experimental Methods

The fine powders of the samples were blended with 10 wt.% of corundum (α-Al2O3) and analyzed by XRD. The XRD was performed using a Rigaku MultiFlex X-ray generator with CuKα radiation. The measurement was performed under the following conditions: tube voltage of 40 kV, tube current of 40 mA, scan speed of 1◦ 2θ min−<sup>1</sup> , scan range of 2θ = 5–70◦ , and step width of 0.02◦ . Siroquant Version 4.0, manufactured by Sietronics, was used for quantitative Rietveld analysis. In the Rietveld analysis, monotrinic and triclinic C3S, α- and β-C2S, cubic C3A, C4AF, Gypsum, Bassanite, Portlandite, Ettringite, Monosulfoaluminate (simply called monosulfate), and Corundum (α-Al2O3) were assigned as targets. The quantitative value of C3S was taken as the sum of the measured values of polymorphic minerals with different crystal structures. The amount of amorphous phase content was calculated according to Equation (1) from the quantitative value obtained by adding 10% mass of corundum (α-Al2O3).

$$\mathbf{P} = \frac{100}{(100 - \mathbf{A})} \times \left(1 - \frac{\mathbf{A}}{\mathbf{R}}\right) \times 100 \tag{1}$$

where P is the amount of amorphous phase content (% mass); A is the mixing rate of corundum (% mass); and R is the quantitative value of corundum (% mass). The amount of amorphous phase content in the sample was calculated from the quantitative value obtained by the Rietveld analysis using corundum as an internal standard substance, according to Equation (1). This result was used to correct the determined values of cement minerals. The loss on ignition was obtained as a mass loss between 105 and 950◦C, using the mass measured after the ignition loss. To determine the hydration degree against the unhydrates, the value obtained by the Rietveld analysis was corrected with the ignition loss value using Equation (2).

$$\mathbf{Q} = \mathbf{Q}\_0 \times \frac{100}{100 - \text{LOI}} \tag{2}$$

where: Q is the corrected quantitative value (%); Q<sup>0</sup> is the quantitative value before the correction (%); and LOI is the loss on ignition (%).

TG/DTA measurements were performed in nitrogen using a HITACHI TG/DTA 7220 analyser at a heating rate of 5◦C/min up to 1,000◦C. Cement setting time experiment was conducted according to JIS R 5201 and the compressive strength of the paste specimens was measured on 3, 7, 14, and 28 days of curing.

## Modeling Approach

The coupled thermodynamic model developed in our previous work (Elakneswaran et al., 2016) was used to model the hydrate assemblage of OPC and FC as a function of hydration time. The thermodynamic model combines the IPhreeqc module (Charlton and Parkhurst, 2011) and Excel for computations. The thermodynamic properties of various phases including Fecontaining hydrates and minerals in the cement system were collected from Cemdata18 (Lothenbach et al., 2008a, 2019) and others (Myers et al., 2014), and the data were converted into a format suitable for PHREEQC. The converted data by Elakneswaran et al. (2016) along with the PHREEQC default thermodynamic database (Parkhust and Appelo, 1999) were used for the calculations.

## RESULTS AND DISCUSSION

The XRD patterns of un-hydrated cements and their pastes hydrated for 1 h and 28 days are shown in **Figure 1**. In addition to the cement clinker minerals, peaks of portlandite, ettringite, and monosulfate were identified in both the cements, and calcium silicate hydrates (C-S-H) exists in amorphous phase. The quantitative amounts of the phases were determined by Rietveld analysis and TG/DTA, and the results are discussed later. No qualitative difference is observed in the product formation in the decreasing of the C2S proportion and increasing of the C4AF proportion in FC. Ettringite production is noted from 1 h in the hydration products, and it transformed to monosulfate after 3 days of hydration in both the cements. Further, the results suggest that the Fe-hydrates may exist in their amorphous states as they cannot be identified by XRD. **Figure 2** illustrates the hydration degree of each clinker mineral in OPC and FC, while the total degree of hydration as a function of time is plotted in **Figure 3**. As given in **Table 1**, both the cements have almost the same proportions of C3S and C3A; therefore, variation in their hydration degree with time is the same. However, the change in the proportions of C2S and C4AF affects their hydration; ∼20% difference was observed between the C4AF hydrations of OPC and FC at 28 days. Nevertheless, the differences did not influence on the total degree of cement hydration (**Figure 3**).

The XRD/Rietveld analysis reveals that the portlandite, ettringite, monosulfate, and amorphous phase contents are the

main hydration products of both OPC and FC, and the variation in their quantities with respect to the hydration time are plotted in **Figure 4**. The phenomenon of hydration in FC is similar to that observed in the conventional Portland cement system: portlandite and amorphous phase contents are formed from the beginning of the hydration process and increase with hydration time, and the decrease in the amount of ettringite correlates with the formation of monosulfate. It can be observed that the increase in ferrite and the decrease in belite in FC do not significantly influence the quantitative value of hydration products at an early age. However, at a later age, a difference is observed particularly in amorphous phase contents. The hydration degree of belite and ferrite in FC further affects the difference (**Figure 2**). Although it is difficult to distinguish the amorphous phase contents using XRD/Rietveld analysis, it is believed that the hydration products which containing Fe are quantified as amorphous.

**Figure 5** illustrates the setting time results of OPC and FC. With an identical value of Blaine specific surface area, both the cements demonstrate similar initial and final setting times. As shown in **Figures 3**, **4**, the hydration reaction and the amounts of solid products formed in FC are equivalent to that of OPC, which attribute to the similar setting time in both the cements. The measured compressive strength of the hydrated cements, mean of 4 samples, is shown in **Figure 6** (standard deviation is shown by bars). It demonstrates that the strength continues to develop

with the hydration time in both the cements, and the ferriterich cement displays a higher compressive strength at early age and becomes equal to OPC at 14 days. Early age compressive strength of OPC paste is slightly lower than the values reported in the hydrated paste, however it shows the similar strength at 28 days (Hoshino et al., 2006). Although the 28 days compressive strength of FC is lower than that of OPC, but it is sufficiently developed.

A coupled thermodynamic model, proposed in a previous study for OPC and slag-blended cementitious materials (Elakneswaran et al., 2016), is employed without any modifications in this study to understand the hydration reaction of the ferrite-rich cement system. The coupled model solves the kinetic equation of the clinker minerals in Excel, and the dissolved quantities are input to PHREEQC for the thermodynamic equilibrium and solid solution calculations. An ideal solid-solution model consisting of six end-members was considered for C-S-H. Prior to predicting the hydration products, the kinetic model is used to predict the hydration degree of the clinker minerals and the total degree of hydration in both the cements. The comparison of results estimated from the model and that obtained from XRD/Rietveld analysis are shown in **Figure 7**. Despite certain discrepancies between the calculation results and the experimental data on individual clinker hydration, the kinetic equations and constants (Elakneswaran et al., 2016) can predict the hydration reaction of ferrite-rich cement and its clinker minerals.

The hydration degree or amount of dissolved clinker phases was used in the thermodynamic calculations with the thermodynamic database for cement minerals and phases. It has been reported that the Fe-siliceous hydrogarnet is expected to form in the hydrated cement system, and it is more stable than Al-siliceous hydrogarnet (Lothenbach et al., 2019). Feettringite, Fe-monosulfate, and Fe-siliceous hydrogarnet were considered in addition to other hydrates commonly found in the hydrated cement system for the calculations of both the cements. For the input of the clinker composition of cement determined by XRD/Rietveld analysis and the mixing conditions, the predicted phase-assemblage composition of the hydrated OPC and FC in mass percentage, relative to the total amount of solid products, is shown in **Figure 8**. Both the cement systems form the same kind of phases with the hydration time, but the initial clinker proportion influences their quantitative values. The presence of Fe-containing phases from the model prediction is consistent with the experimental observation (Dilnesa et al., 2011). Fe-ettringite formed at early age, destabilized and transformed to form Al-ettringite; Fe-monosulfate was not produced. The stable Fe-siliceous hydrogarnet starts to form from early age and continues with the hydration time. The dissolved Fe from the ferrite phase transforms into Feettringite and Fe-siliceous hydrogarnet. The high quantity of ferrite phase in the FC enhances the formation of the Fe-containing phases.

The predicted mass percentage of portlandite, ettringite, and amorphous phase content are compared with the experimental data in **Figure 9**. As reported, the poor crystalline phases of AFm and Fe-containing phases are difficult to quantify by XRD/Rietveld analysis, since they might exist in the amorphous phase as well (Matschei et al., 2007; Lothenbach et al., 2008b). Therefore, the summation of the total amorphous phase content and monosulfate was compared with the addition of the predicted C-S-H, monosulfate, and Fe-siliceous hydrogarnet. Apart from certain variations, the calculated phase assemblage generally agrees well with the experimentally determined values, both qualitatively and quantitatively. A high proportion of the Fe-containing phases formed from the hydration of ferrite phase in the FC contributes significantly to the increase in the calculated amorphous phase content. The composition of the phase assemblage in terms of volume fraction is shown in **Figure 10**. The volume of each hydrate was calculated from its molar volume, and the total porosity is estimated based on the volumes of hydrates and un-hydrated cement at any time and the initial volume of the paste, assuming that chemical shrinkage does not occur. The calculated total porosity is found to be closely associated with the measured compressive strength, as shown in **Figure 11**, and the relationship is found to be more prominent in FC. As shown in **Figure 10**, the porosity does not change significantly with the increase in ferrite and decrease in belite phase proportions in FC, though certain differences are observed between the quantitative values of the hydration products of the two cements.

## CONCLUSIONS

Burning of clinker at ∼1,350◦C, which is 100◦C lower than the burning temperature of OPC, produces a high-ferrite and lowbelite content cement. The manufacturing process of this cement may release 5% lesser CO<sup>2</sup> than that of OPC, thus demonstrating a potential to substitute the OPC. The hydration behavior of the FC is similar to that of OPC, but the different proportions of the clinkers influence their hydration degree; low hydration degree of C2S and C4AF are found in FC, in contrast to the OPC. The hydration products were identified and quantified by the XRD/Rietveld analysis and TG/DTA, except the C-S-H and Fe-containing phases that were accounted as amorphous phase contents. The degree of hydration and hydration products of ferrite-rich cement are comparatively similar to that of the OPC; therefore, both the cements exhibit the same initial and final setting times. The ferrite-rich cement develops higher initial compressive strength than the OPC, and displays slow strength development at later age, but achieves the required final strength. Thermodynamic modeling results reveal the hydration reaction of FC. Fe-siliceous hydrogarnet is the main Fe-containing phase formed during the hydration of both FC and OPC, and Feettringite is produced at a very early age and then converted to Al-ettringite. The predicted quantities of the Fe-containing phases strongly depend on the proportion of ferrite in the cement. The estimated compositions of portlandite, ettringite, and amorphous phase contents agree satisfactorily with the

## REFERENCES


XRD/Rietveld analysis and TG/DTA results. Furthermore, the predicted total porosity correlates well with the measured compressive strength.

The present study focuses on the hydration reaction of FC and compares it with that of the OPC. The next step would be identifying and quantifying the Fe-containing phases formed in the cement and verifying it with a thermodynamic model, which requires additional thermodynamic data for the other Fe-containing phases. In addition, Mancini (Mancini et al., 2018) has recently pointed out that Fe can incorporate into C-S-H as well, which should be included in the C-S-H model by either surface complexation or phase equilibrium.

## AUTHOR CONTRIBUTIONS

YE, TC, HK, NN, YM, and TN designed the experimental plan. NN, KM, YM, TC conducted the experiments and data analysis. YE and NN carried out the modeling work. YE, NN, YM, and TC drafted the manuscript. TC, HK, and TN contributed to the discussions.


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Elakneswaran, Noguchi, Matumoto, Morinaga, Chabayashi, Kato and Nawa. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

## Effect of Pre-dispersing Metakaolin in Water on the Properties, Hydration, and Metakaolin Distribution in Mortar

Haining Geng<sup>1</sup> , Wei Chen<sup>2</sup> , Qiu Li <sup>2</sup> \*, Zhonghe Shui <sup>2</sup> and Bo Yuan<sup>3</sup>

*<sup>1</sup> Department of Civil Engineering, Hubei Urban Construction Vocational and Technological College, Wuhan, China, <sup>2</sup> State Key Laboratory of Silicate Materials for Architectures, Wuhan University of Technology, Wuhan, China, <sup>3</sup> School of Materials Science and Engineering, Wuhan University of Technology, Wuhan, China*

The effects of pre-dispersing metakaolin (MK) in water on the properties and hydration of mortar and distribution of MK particles in mortar were characterized by analytical techniques. Zeta-potential of MK-water dispersion significantly decreased with the increase of pH of solution, resulting in a stable dispersion at pH of 8 with zeta-potential of −40 mV. The bulk density of MK-water slurry rapidly increased with water/MK ratio due to the surface tension force of water and filling of spaces by water, resulting in a maximum bulk density of 665.2 kg/m<sup>3</sup> at water/MK ratio of 1.0. Comparing to that of powder MK, pre-dispersing MK in water did not alter hydration products, slightly improved compressive strength, and pore structure, but further reduced shrinkage and significantly improved the uniformed distribution of MK particles in mortar.

#### Edited by:

*Maria Juenger, University of Texas at Austin, United States*

#### Reviewed by:

*Maria Stefanidou, Aristotle University of Thessaloniki, Greece Zehra Canan Girgin, Yildiz Technical University, Turkey*

> \*Correspondence: *Qiu Li qiu-li@whut.edu.cn*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *09 February 2019* Accepted: *17 April 2019* Published: *03 May 2019*

#### Citation:

*Geng H, Chen W, Li Q, Shui Z and Yuan B (2019) Effect of Pre-dispersing Metakaolin in Water on the Properties, Hydration, and Metakaolin Distribution in Mortar. Front. Mater. 6:99. doi: 10.3389/fmats.2019.00099* Keywords: zeta-potential, viscosity, hydration products, compressive strength, shrinkage, coefficient of variation

## INTRODUCTION

Concrete is one of the most used man-made materials for construction, due to its availability, suitable workability, properties, and durability. However, concrete has been criticized in recent decades for its contribution to CO<sup>2</sup> footprint, which is mainly from one of its components, Portland cement (Scrivener and Kirkpatrick, 2008). To reduce the CO<sup>2</sup> footprint of construction industry, mineral additives have been introduced to concrete in the aim to reduce the amount of Portland cement required. Mineral additives include fly ash, silica fume, slag, and metakaolin (MK). Mineral additives are usually amorphous in nature with high silicate/aluminate content, which can react with calcium hydroxide formed during hydration of Portland cement through pozzolanic reaction (Taylor, 1997).

MK attracted much attention among mineral additives recently due to its high contribution to strength and durability when introduced to concrete (Frias and Cabrera, 2000; Aquino et al., 2001; Gueneyisi et al., 2008; Goncalves et al., 2009; Shui et al., 2010; Fernandez et al., 2011; Rashad, 2013), resulting from its effect of acceleration of hydration of Portland cement and refinement of pore structure of concrete or mortar (Buchwald et al., 2007; Lagier and Kurtis, 2007; Cassagnabere et al., 2010; Li et al., 2010; Kadri et al., 2011). MK is manufactured by calcining kaolin at 650–800◦C to remove interlayer water and OH<sup>−</sup> resulting in the deterioration of long-range order of silicon-aluminum layer structure (Murat and Comel, 1983; He et al., 1995; Badogiannis et al., 2005; Rashad, 2013). Amorphous MK with pozzolanic activity was obtained with formula of Al2O3·2SiO<sup>2</sup> (Coleman and McWhinnie, 2000; Sabir et al., 2001; Love et al., 2007; Bich et al., 2009; Fernandez et al., 2011). When MK was added into concrete, the compressive strength could be improved, permeability, and porosity could be decreased and pore structure could be refined due to the filler effect, acceleration of hydration, and pozzolanic reaction with CH to form C-S-H (Wild et al., 1996; Aquino et al., 2001; Lagier and Kurtis, 2007; Gueneyisi et al., 2008; Shui et al., 2010; Fernandez et al., 2011; Kadri et al., 2011). The compressive strength increased with the content of MK, and the maximum effect was reached at 28 days as observed by many researchers previously (Justice and Kurtis, 2007; Khatib, 2008, 2009; Arikan et al., 2009; Johari et al., 2011; Kadri et al., 2011; Antoni et al., 2012; Li et al., 2015). The early age (before 24 h) autogenous

TABLE 1 | Chemical composition and physical properties of cement and metakaolin used in this study.


shrinkage of concrete decreased with the increase of content of MK due to the dilution effect of MK, although the longterm autogenous shrinkage of concrete containing MK could be higher than that of control specimen, due to the acceleration of Portland cement hydration, and pozzolanic reaction by MK with calcium hydroxide (Kinuthia et al., 2000; Brooks and Johari, 2001). With the increase of MK addition, the drying shrinkage of concrete could be significantly reduced by 50%, comparing to the control specimen, due to the reduced amount of evaporable water as a result of hydration, and pozzolanic reaction (Brooks and Johari, 2001).

Due to the high specific surface area and small particle size, there is a high tendency for MK particles to agglomerate, resulting in the non-uniform properties of concrete when MK was introduced as SCMs. The low stacking density of MK, which is ∼1,500 kg/m<sup>3</sup> , reduced the capacity and efficiency of transportation. Approximately over 40% of MK particles is under the size of 2.5µm, which could pollute the environment and be harmful to the respiratory system of human. The authors' preliminary study suggested non-uniform hydration products,



mechanical properties, and durability of concrete specimens with MK addition, resulted from the uneven distribution of MK particles after introducing into concrete due to the agglomeration of MK particles. Some previous research also reported that the agglomeration of MK reduced the pozzolanic activity of MK, together with a lower compressive strength, different microstructure and macrostructure of paste, and geopolymer (Medri et al., 2010; Ilic et al., 2018).

In this study, aiming to improve the capacity and efficiency of transportation, prevent the air pollution caused by MK powder, reduce the agglomeration and improve the uniform distribution of MK particles in mortar matrix and the properties of mortar, MK was pre-dispersed in water as slurry before being added in mortar to increase the density and homogeneity. A range of analytical techniques was employed to characterize the effect of pH of solution and water/MK ratio on the physical properties and stability of dispersion of MK-water slurry, and the effect of pre-dispersed MK slurry on the properties and hydration of cement mortar was evaluated, together with the improvement of distribution of MK particles in mortar after pre-dispersing.

## MATERIALS AND EXPERIMENTAL

## Materials

In this study, ordinary Portland cement (OPC) (Huaxin Cement Co. Ltd., China) was used MK was provided by Maoming Kaolin Science and Technology Co. Ltd., China, with a specific surface area of 2,800 m<sup>2</sup> /kg. **Table 1** shows the chemical composition and properties of OPC and MK. The particle size distribution of MK is shown in **Figure 1**, which was analyzed with Malvern Mastersizer 2000 particle size analyser (Malvern, UK) by dispersing in sodium hexametaphosphate solution. According to X-ray diffraction (XRD) pattern in **Figure 2**, MK is mainly amorphous. Fine aggregate is natural sand with an apparent density of 2,640 kg/m<sup>3</sup> , stacking density of 1,458 kg/m<sup>3</sup> and fineness modulus of 2.91. The grading of fine aggregate is shown in **Table 2**.

## Experimental Details Zeta Potential Test

MK-water colloidal dispersion was prepared with water/MK ratio of 1,500 for zeta potential test under various pH value of solution. The pH value of the MK-water dispersion was varied by addition of HNO<sup>3</sup> and NaOH. The pH value of MK-water dispersion was measured by Inesa PHSJ-3F pH Meter (China). Zeta potentials of MK particles in solution with various pH value were conducted by Zeta-Meter System 3.0+ (USA). All measurement was conducted under room temperature.

## Bulk Density and Viscosity of MK-Water Slurry Test

The bulk density of MK-water slurry was measured by adding water into MK powder with various water/MK ratio, mechanically mixing the slurry with a ball miller for 30 min,


TABLE 4 | Zeta-potential of MK dispersion under various pH in solution with MK/water ratio of 1,500.


then measuring the weight, and volume of slurry. The viscosity of MK-water slurry with various water/MK ratio was evaluated by Brookfield R/S-SST (USA) with shear rate of 60 r/min. The VANE rotor for the test was V40-20, the container used was Griffin beaker of 600 ml and the temperature was constant at room temperature of 25 ± 5 ◦C.

#### Stability of MK-Water Slurry With Time

The stability of MK-water slurry with water/MK ratio of 1.86 was evaluated by the differences of viscosity at 0 h and 24 h after mixing and the segregation of water from the slurry at 24 and 48 h after mixing. The segregation of water was measured as follows.

One hundred milliliter of slurry was put in the measuring cylinder after mixing. The volume of water segregated from slurry was recorded from the top of slurry after 24 and 48 h. The measurement was performed under room temperature and no vibration was applied.

#### Mix Design and Curing Conditions

The pre-dispersed MK slurry used in this study was with water/MK ratio of 1.86. The mix design of mortar and paste is shown in **Table 3**. OPC was substituted by 5 or 10 wt% powder MK or pre-dispersed MK slurry in mortar and paste. The water/binder ratio was 0.5 for all mortar and paste mixes. Specimens were demoulded at 1 day and cured in a fog room under standard moist curing condition with temperature of 20◦C and relative humidity above 90%.

#### Compressive Strength and Porosity Test

For compressive strength test, the mortar and paste specimens were cast into the size of 40 × 40 × 160 mm. Nine specimens from each mix were used for compressive strength test. The compressive strength test was performed at 3, 7, and 28 days for all mixes according to Chinese standard GB/T 50081-2002<sup>1</sup> . The results from all specimens were included for the calculation of compressive strength and standard deviation. Porosity test was conducted on paste specimens according to ASTM C642-13<sup>2</sup> .

#### Phase Assemblage

XRD analysis was conducted on the paste specimens at 3 and 28 days. Crystalline minerals in freshly ground paste specimens was characterized by Bruker D8 Advance with Cu Kα = 1.5406 Å, step size of 0.019◦ , measuring time of 141.804 s/step, start position 5◦ and end position 70◦ . The obtained XRD pattern was analyzed by PANalytical X'pert Highscore Plus software with PDF2004 database.

#### Shrinkage

The shrinkage of mortar specimens was evaluated according to Chinese standard JGJ/T70-2009<sup>3</sup> . Mortar specimens were cast into the size of 40 × 40 × 160 mm and cured under standard moist conditions for 7 days before demoulding, where cement mortar without MK addition, cement mortar with 10 wt% MK addition as dry powder and cement mortar with 10 wt% MK addition as dispersed slurry were named as M0, MMK10, and MDMK10, respectively. Then the initial length of specimens were recorded and the specimens were cured under temperature of 20 ± 2 ◦C and relative humidity of 60 ± 5%. The lengths of specimens were measured at 7, 14, 21, 28, 35, and 42 days after demoulding. Three measurement were performed on each age. The shrinkage was calculated by the following equation.

$$\varepsilon\_{t} = \frac{L\_{0} - L\_{t}}{L - L\_{d}} \tag{1}$$

Where ε<sup>t</sup> is the shrinkage of mortar at t days, L<sup>0</sup> is the length of specimen at demoulding (mm), L is the standard length of 160 mm, L<sup>d</sup> is the length of test head in the mortar (mm), and Lt is the length of specimen at t days (mm).

#### Coefficient of Variation of Compressive Strength

The coefficient of variation (Cv) is a standardized measure of dispersion of a probability distribution. It is defined as the ratio

<sup>1</sup> Standard for Test Method of Mechanical Properties on Ordinary Concrete. Ministry of Construction of the People's Republic of China (2002).

<sup>2</sup>ASTM C642-13: Standard Test Method for Density, Absorption, and Voids in Hardened Concrete. West Conshohocken, PA: ASTM International (2013).

<sup>3</sup> JGJ/T70-2009: Standard for Test Method of Performance on Building Mortar. Beijing: Ministry of Housing and Urban-Rural Development of People's Republic of China (2009).

of the standard deviation (SD) to the mean (µ), as follows.

$$\mathcal{C}\_{\text{v}} = \frac{\text{SD}}{\mu} \tag{2}$$

In this study, the coefficient of variation was used to express the distribution of compressive strength test results of mortar specimens of M0, MMK10, and MDMK10 at 28 days. Eighteen specimens were tested for each mix. The mean value, standard deviation and coefficient of variation were calculated to assess the distribution of MK particles in mortar matrix. The results from all specimens were included for the calculation of mean value, standard deviation, and coefficient of variation, to evaluate the homogeneity.

## RESULTS AND DISCUSSION

## Stability of MK-Water Dispersion Under Various pH

Zeta potential is a key indicator of the stability of colloidal dispersion. The magnitude of zeta potential indicates the electrostatic repulsion between adjacent, similarly charged particles in the dispersion, with a higher zeta potential indicating better stability, and less aggregation. According to the relationship between zeta potential and pH value of MK-water dispersion with water/MK ratio of 1,500 observed in this study (**Figure 3** and **Table 4**), zeta potential of MK-water dispersion significantly decreased with the increase of pH of solution, from 16.4 to −48.2 mV when pH increased from 1.5 to 10. The concentration of OH<sup>−</sup> increases with pH value, providing more OH<sup>−</sup> available to attach to the surface of metakaolin particles, resulting in the continuous decrease of zeta potential to ∼-40 mV at pH of 8, indicating that the MK-water dispersion was stably dispersed. In solution of pH over 10, pozzolanic reaction of MK will take place, which is not beneficial for the purpose of storage and dispersion of metakaolin for construction. Severe aggregation was observed when pH was more than 10 in this study, indicating the activation of MK due to the pozzolanic reaction (He et al., 1995). Considering the stability of dispersion of MK slurry, dispersing MK with water, which has a pH value of 7, should result in a stable dispersion without pozzolanic reaction. In this study, de-ionized water was used for dispersing at pH of 7.

## Bulk Density and Viscosity Under Various Water/MK Ratio

To improve the bulk density of MK slurry and distribution of MK in mortar, MK-water slurry with water/MK ratio of 0–2.5 were prepared and the bulk density and viscosity were evaluated. According to the results shown in **Figure 4**, the bulk density of MK slurry rapidly increased with the addition of water due to the decrease of distance between particles by surface tension force of water and filling of water between particles, reaching the maximum of 665.2 kg/m<sup>3</sup> at water/MK weight ratio of 1.0, where the packing of MK particles reached the closest packing. Then the bulk density of MK slurry gradually decreased with the increase of water/MK ratio, due to the increase of distance between MK particles by continuous addition of water.

The viscosity was measured at water/MK ratio over 1.3, where the bulk density of MK slurry started to decrease (**Figure 4** and **Table 5**). The viscosity decreased rapidly with the increase of water/MK ratio, from 693.5 mPa·s at water/MK ratio of 1.25 to 30.0 mPa·s at water/MK ratio of 1.85. To obtain a low viscosity and relative high MK content, water/MK ratio of 1.86 was chosen to evaluate the effect of pre-dispersed MK slurry on the properties and hydration of cement mortar in this study.

The stability of MK-water slurry with water/MK ratio of 1.86 was evaluated. The viscosity of MK slurry increased from 28.6 mPa·s at 0 h after mixing to 49.3 mPa·s at 24 h after mixing. The

FIGURE 5 | Compressive strength of mortars with powder MK, pre-dispersed MK, and per-dispersed MK settled for 1, 2, and 4 days before addition.

TABLE 5 | Density and viscosity of MK slurry under various water/MK ratio.


water segregation of MK slurry was 4.7 ml water per 100 ml slurry at 24 h after mixing and 8.2 ml water per 100 ml slurry at 48 h after mixing.

## Compressive Strength

The compressive strength of specimen M0, MMK10, MDMK10 at 3, 7 and 28 days were shown in **Figure 5** and **Table 6**. For MDMK10 specimens, the pre-dispersed MK slurry was added in mortar after preparing or settling for 1, 2, and 4 days, respectively. Comparing to the compressive strength of M0 specimen, by addition of 10 wt% powder MK, the compressive strength increased by 29.4, 51.6, and 24.7% at 3, 7, and 28 days, respectively. Comparing the compressive strength of mortar with pre-dispersed MK with those with powder MK, pre-dispersing of MK did not hinder the effect of MK on the compressive strength of mortar. At 28 days, the compressive strength of mortar with pre-dispersed MK was slightly higher than that with powder MK. With a pH value of 7, water could not activate MK, which was activated later by calcium hydroxide formed during the hydration of cement in mortar. The improvement of compressive strength of mortar by MK was due to the accelerated hydration of cement, filler effect, and pore structure refinement by MK, as reported in previous studies (Li et al., 2015).

Comparing with the compressive strength of mortar with non-settled pre-dispersed MK specimen, the settlement of predispersed MK up to 4 days did not alter the effect of MK on the compressive strength of mortar with pre-dispersed MK, indicating that there was no reaction between water and MK during settlement up to 4 days and MK particles were stably dispersed in water as being suggested by results from zetapotential test under various pH values.

## Porosity

The porosity of paste specimens is shown in **Figure 6** and **Table 7**, together with the relationship between porosity and compressive strength. For the pastes with powder or pre-dispersed MK addition, the compressive strength and porosity did not alter much at 3 days, comparing to that of the control specimen P0. However, at 7 and 28 days, while the porosity was slightly lower than that of the control specimen at the same age, the compressive strength greatly increased with ages, comparing to those of the control specimen P0. There is a linear relationship between the compressive strength and porosity of mortar for each group, where the compressive strength increased with the decrease of porosity, which is similar to previous studies (Frias and Cabrera, 2000). Although the relationship between porosity and compressive strength kept linear for the control, powder MK and pre-dispersed MK specimens, these results indicate that the pore structure was refined through addition of powder or pre-dispersed MK while the porosity did not change much. The relationships between porosity and compressive strength of powder and pre-dispersed MK specimens were similar, indicating that the pore structure refinement effect of MK was not affected by pre-dispersing MK in water. It is notable that the relationship is valid in the porosity range of 28–46% in this study. When the porosity is close to 0, the relationship could be invalid.

## Hydration Products

The crystalline phases of paste specimens were characterized by XRD (**Figure 7**). According to these results, the hydration products of ettringite, Portlandite, calcite, hemicarboaluminate (Hc), and monocarboaluminate (Mc) were identified for P0, PMK10, and PDMK10 specimens at 3 and 28 days, except for P0 at 3 days, where monocarboaluminate was not identified. The hydration products identified were similar to those from previous studies (Frias et al., 2012; Gameiro et al., 2012; Tironi et al., 2013). The intensity of peaks attributed to monocarboaluminate

cement paste, pastes with powder MK, and pre-dispersed MK. Specimens with a porosity of ∼45% were cured for 3 days, those with a porosity of 35–42% were cured for 7 days and those with a porosity of 25–35% were cured for 28 days.

TABLE 6 | Compressive strength of mortar specimens (MPa).



TABLE 7 | Compressive strength with standard deviation and porosity of paste specimens.

increased with age for P0, PMK10, and PDMK10 specimens, indicating the increase of content of monocarboaluminate with age. The intensity of peaks of these hydration products was similar in PMK10 and PDMK10 specimens at the same age, indicating that the effects of powder MK and pre-dispersed MK on the hydration of cement were similar and the effects of MK on cement hydration were not hindered by dispersing in water, as being indicated by the compressive strength and porosity tests. The reason is that MK was not activated by water, with a pH value of 7. There was no reaction between water and MK observed.

## Shrinkage

The shrinkage of mortar was evaluated and the results are shown in **Figure 8** and **Table 8**. The shrinkage of mortar was reduced at all ages through addition of powder MK, due to the refined pore structure of mortar by addition of MK, as indicated by previous studies (Li and Yao, 2001), which suggested that the reduced shrinkage by MK addition can be attributed to the lower amount of evaporable water available due to the hydration and pozzolanic reaction of cement and MK. At 28 days, the shrinkage was reduced by 21%, which was similar to previous studies (Brooks and Johari, 2001; Khatib, 2008; Guneyisi et al., 2010, 2012; Cheng et al., 2016). Pre-dispersing of MK reduced the shrinkage of mortar by 41% at 28 days, due to the internal curing effect by water stored in the interlayer of MK during the predispersing process.

## Distribution of MK in Mortar Matrix

Preliminary study showed that the agglomeration of MK in mortar or concrete matrix resulted in the non-uniform properties or hydration degree of mortar or concrete. In this study, the distribution of MK particles of powder or pre-dispersed MK were evaluated by coefficient of variation of compressive strength results of 18 specimens from each mortar mix at 28 days. The variation of coefficient is unit-less and represents the distribution of results. The variation of coefficient of results with different mean value can be compared, unlike the standard deviation. The compressive strengths with the range of results, standard deviation, and coefficient of variation were shown in **Figure 9** and **Table 9** for M0, MMK10, and MDMK10 at 28 days. According to these results, the compressive strength of mortar improved by MK addition, and pre-dispersing of MK further improved

the compressive strength, as discussed above. Comparing to that of the control specimen, the range of compressive strength of specimens from the same mix significantly increased by addition of powder MK and slightly increased by addition of pre-dispersed MK. Similar to the range of compressive strength, standard deviation (SD) of compressive strength was greatly increased by powder MK and slightly increased by pre-dispersed MK, comparing to that of control specimen. The coefficient of variation (Cv), which indicates the evenness of distribution of MK particles, shown as the size of circle in **Figure 9**, significantly increased by powder MK addition and approximately doubled by pre-dispersing MK addition, comparing to that of M0 specimens.

at water/MK ratio

Zeta-potential of MK-water dispersion significantly decreased with the increase of pH of solution, with zeta-potential of ∼-40 mV at pH of 8, indicating a stable dispersion, due to more OH<sup>−</sup> attachment on the surface of MK particles without

During the addition of water, the bulk density of MK-water slurry firstly rapidly increased with water/MK ratio, due to the decrease of distance between MK particles by surface tension force of water and filling of spaces between particles by water.

of 1.0, indicating the closest packing of MK particles. Then the bulk density gradually decreased with the increase of water/MK ratio, due to the increase of distance between MK particles by continuous addition of water. Viscosity of MK-water slurry

The maximum bulk density was 665.2 kg/m<sup>3</sup>

decreased rapidly with the increase of water/MK ratio.

activation of MK.

These results suggest that the particle of powder MK distributed unevenly in the mortar and pre-dispersing of MK by water improved the uniformness of distribution of MK particles in mortar. As indicated by the zeta-potential of dispersed MK slurry, pre-dispersing improved the distribution of MK particles in slurry, and formed a stable dispersing system. Viscosity was also significantly improved through dispersing in water, so as to improve the uniformness of distribution of MK particles in mortar during mixing, although the hydration of cement and MK was not affected by pre-dispersing, as indicated by XRD results.

## CONCLUSIONS

The effects of dispersing MK in water on the properties of MK, the properties and hydration of mortar and the distribution of MK particles in mortar were characterized by a range of analytical techniques. According the results, the following conclusions can be drawn.

or pre-dispersed MK. The results are shown as mean value of 18 specimens of each mix, range of results (fc-max and fc-min), standard deviation and coefficient of variation (size of circle).




TABLE 9 | Average value, standard deviation, range, and coefficient of deviation of compressive strength at 28 days by calculating 18 specimens.

Both powder MK and pre-dispersed MK increased the compressive strength of mortar. There was linear relationship between the compressive strength and porosity of mortar. Neither pre-dispersing MK in water nor settlement of pre-dispersed MK slurry hindered the effect of MK on the compressive strength. Pre-dispersing MK in water slightly decreased the porosity of mortar, comparing to that with powder MK. The hydration products were not affected by pre-dispersing MK in water. Addition of powder MK reduced the shrinkage of mortar, which was reduced further by pre-dispersing MK in water.

Comparing to powder MK, pre-dispersed MK improved the uniformed distribution of MK particles in mortar according to the range, standard deviation, and coefficient of variation of compressive strength results, due to the stable dispersing of MK particle in pre-dispersed MK slurry.

## REFERENCES


## AUTHOR CONTRIBUTIONS

HG, WC, QL, and ZS contributed conception and design of the study. BY performed the properties tests. HG, WC, and QL performed the microstructure and statistical analysis. HG wrote the first draft of the manuscript. WC, QL, ZS, and BY wrote sections of the manuscript. All authors contributed to manuscript revision, read, and approved the submitted version.

## ACKNOWLEDGMENTS

Thanks to Dr. Jian Huang (Wuhan University of Technology) for beneficial discussion on the viscosity test. This research was financially supported by the Natural Science Foundation of China (project 51611530545, 51672199), the YangFan Innovative and Entrepreneurial Research Team Project (No. 201312C12) and the Research Fund of State Key Laboratory of Silicate Materials for Architectures (Wuhan University of Technology).


pastes hydrated at 25 degrees C. Cement Concrete Res. 37, 109–117. doi: 10.1016/j.cemconres.2006.11.012


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Geng, Chen, Li, Shui and Yuan. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Silico-Aluminophosphate and Alkali-Aluminosilicate Geopolymers: A Comparative Review

Yan-Shuai Wang, Yazan Alrefaei and Jian-Guo Dai\*

*Department of Civil and Environmental Engineering, The Hong Kong Polytechnic University, Hong Kong, Hong Kong*

Chemically activated materials (often termed as geopolymer) have received attracting attentions in civil, material and environmental research fields as a toolkit alternative to traditional Portland cement in specific applications. This paper presents a comparative review on silico-aluminophosphate (SAP) geopolymers in terms of definition, chemistries involved during geopolymerization, mechanical performance, durability, environmental impacts, and their potentials in applications relative to conventional alkali-aluminosilicate (AAS) geopolymers. Recommendations for future applications are also highlighted. It is found that S-A-P gels with six-coordinated aluminum environment dominate in SAP geopolymers, while the aluminum in N-A-S-H gels formed in the AAS geopolymers is characterized by four-coordinated features. Besides, the slow performance development of SAP geopolymer matrix under ambient temperature curing can be compensated through incorporating additional countermeasures (e.g., metal sources) which allow the tailored design of such geopolymers for certain *in-situ* applications. Generally, the calcium-bearing C-(A)-S-H gels co-existing with N-A-S-H gels are dominant in AAS geopolymers, while the S-A-P gels enhanced by phosphate-containing crystalline/amorphous phases are the main products in SAP geopolymers. The SAP geopolymers show their environmental friendliness relative to the AAS geopolymers due to the utilization of phosphate activators that require lower production energy relative to silicate-containing activators. However, the higher cost of phosphate activators may confine the applications of SAP geopolymers in some exquisite or special fields.

#### Edited by:

*Miroslav Komljenovic, University of Belgrade, Serbia*

#### Reviewed by:

*Kenneth John MacKenzie, MacDiarmid Institute for Advanced Materials and Nanotechnology, New Zealand Yingwu Zhou, Shenzhen University, China Hubert Rahier, Vrije University Brussel, Belgium*

> \*Correspondence: *Jian-Guo Dai cejgdai@polyu.edu.hk*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *14 January 2019* Accepted: *23 April 2019* Published: *07 May 2019*

#### Citation:

*Wang Y-S, Alrefaei Y and Dai J-G (2019) Silico-Aluminophosphate and Alkali-Aluminosilicate Geopolymers: A Comparative Review. Front. Mater. 6:106. doi: 10.3389/fmats.2019.00106* Keywords: silico-aluminophosphate, alkali-aluminosilicate, geopolymerization mechanism, environmental impacts, sustainable development

## GENERAL ABOUT CEMENT AND GEOPOLYMER BINDERS

Cement, routinely denoted as ordinary Portland cement (OPC), is described as a cohesive material that can set and bind sands and gravels together; such binder is excessively used to prop up the modern urban agglomerations (Taylor, 1997; Neville, 2011; Shi et al., 2011). However, the annual huge consumption of contemporary cement has caused several environmental issues worldwide due to the high carbon footprint (e.g., the inescapable decarbonation of limestone) during cement production. Geopolymer binders, characterized by higher durability and greenness relative to the conventional OPC, have gained many attentions recently as a promising partial substitute to OPC (Provis and Bernal, 2014; Provis et al., 2015).

Wang et al. Review of Silico-Aluminophosphate Geopolymers

Initially, the geopolymer was strictly defined as the alkaliactivation of raw aluminosilicates (e.g., metakaolin as a precursor) without (or with little) other components (i.e., Ca, Mg, etc.) (Davidovits, 2011). The geopolymer formation can be described as a polymer-like condensation process (as shown in Equation 1) where the loss of water between two hydroxyl groups occurs in inorganic materials (Wilson and Nicholson, 2005; Davidovits, 2011). Such chemistry is also extended to the reaction between two different hydroxylcontaining groups, e.g., the isolated aluminate (IV) and silicate (IV) units to reconstruct connected gels with chain or ring network structure through condensation chemistry in alkaliactivated aluminosilicate binders, as shown in Equation (2) (De Jong et al., 1983). Thus, such high-connectivity molecular structures of the formed products can perform excellently in terms of strength and durability (Walkley et al., 2017). Later on, the phosphate-activated aluminosilicate was developed and incorporated into the geopolymer category due to the similar condensation process (Liu et al., 2010; Davidovits, 2011; Guo et al., 2016; Wang et al., 2017, 2018b).

$$q \equiv Q - OH + HO - Q \equiv \overset{\text{Catalyst}}{\Longrightarrow} Q - O - Q \equiv + H\_2O \tag{1}$$

where, Q denotes the hydroxyl-containing group (e.g., orthosilicic group).

$$I \equiv \text{Si} - OH + HO - Al \equiv \overset{Alkaline}{\Longrightarrow} \equiv \text{Si} - O - Al \equiv + H\_2O \text{ (2)}$$

The catalytic activation process of un-doped aluminosilicate precursors is usually enhanced by external energies (e.g., heat, ultraviolet, and microwave), otherwise it requires a long hardening duration, depending on the reactivity of the aluminosilicate sources (Agarwal, 2006; Zhang et al., 2016). Such special curing methods are acceptable in precast geopolymer industry, yet very challenging to be implemented in in-situ constructional operations (Provis and Bernal, 2014). Therefore, it is imperative to develop other chemical processes for roomtemperature hardening mechanism in geopolymers.

In alkali-activated systems, the calcium-containing solid aluminosilicate sources, such as ground granulated blast-furnace slag (GGBS) or high-calcium fly ash (FA), are habituated under high alkaline environment to form C-(A)-S-H gels rapidly (Wang et al., 1994; Shi and Qian, 2000; Li et al., 2010). It should be noted that the high volume of calcium species in FA (e.g., Class C) greatly influences the operating time of their corresponding binders due to their flash setting, which needs to be overcome by the use of retarding additives, e.g., sodium tetraborate (Nicholson et al., 2005; Tailby and MacKenzie, 2010). The C-(A)-S-H gel shows some similar features to that of C-S-H gel that is dominant in the conventional OPC cement. In the meantime, soluble silicates (sometimes mixed with aluminates) are generally introduced into the alkaline activating solution with the same aim of obtaining a rapid sol/gel transition (Benavent et al., 2016). Such actions provide the alkali-activated binders matrices with high efficiency in terms of mechanical performance and durability (Ding et al., 2016, 2018; Shang et al., 2018; Xue et al., 2018).

As a comparison, in phosphate-activated systems, some solid agents (e.g., alkali metals) can be incorporated into the aluminosilicate sources to induce an acid-base chemistry which generally facilitates the formation of crystalline (e.g., products in magnesium/calcium phosphate cement) or amorphous phosphate phases during the chemical activation of aluminosilicates (Wagh, 2016; Wang, 2018). Such acidbase chemistry can be described in Equations (3–5). Initially, the release of cations from the metal oxides occurs through hydrolysis reaction (shown in Equation 3), following, the metal cations react with water to form the aquo ions (hydration reaction in Equation 4). The polyacids generated by the acidic salt or acid (e.g., phosphate radicals) react with the aquo ions (e.g., Mg[H2O]2<sup>+</sup> 6 ) as described in Equation (5) (Ding et al., 2012). The whole conversion process from sols (i.e., colloidal solutions) to gels can be described as exothermic and rapid. Such fast acid-base chemistry can be used for possible improvements in the early performance of the phosphate activation of aluminosilicate.

$$^{1}MO\_{\frac{\theta}{2}} + \frac{p}{2}H\_{2}O \Rightarrow M^{p+} + pOH^{-} \tag{3}$$

$$M^{p+} + qH\_2O \Rightarrow [M(H\_2O)\_q]^{p+} \tag{4}$$

$$[H\_nPO\_4^{-(3-n)} + [M(H\_2O)\_q]^{p+} \Rightarrow [M(H\_nPO\_4)]^{n+p-3} + qH\_2O \text{ (5)}$$

Due to the environmental concerns, the currently used aluminosilicate sources in the geopolymer production are usually industrial byproducts (e.g., FA and GGBS) rather than the materials available naturally (e.g., metakaolin [MK]) (Toniolo and Boccaccini, 2017). The MK [2SiO2·Al2O3] is the product of the dehydroxylated clay mineral kaolinite [Si2O5Al2(OH)4] upon calcination under high temperature (Sperinck et al., 2011) to yield disordered aluminosilicates with Si-O-Al(1O) (i.e., Al in five-fold coordination) and Si-O-Al(2O) (i.e., Al in four-fold coordination), which is an ideal precursor for geopolymer preparation. Due to the limited availability of clay mineral and high temperature calcination that may cause undesirable high cost and energy consumption, the MK is still confined to laboratory use with the aim of investigating the mechanism of geopolymer formation (Khale and Chaudhary, 2007); however, the use of MK is recommended for some practical low-volume high-value applications (e.g., bioactive ceramics, electronic ceramics, drug delivery agents, etc.) (MacKenzie, 2015). On the other hand, the industrial raw materials (typically the byproducts of power stations like FA and iron & steel industries like GGBS) are chemically complex due to their diverse production technologies (Pacheco-Torgal et al., 2008; Yao et al., 2015; Xu and Shi, 2018). Thus, some "impurities" (the impurities mean different available components in the precursor other than aluminosilicates) in the byproducts will induce non-geopolymeric chemistries, thus the acidic or alkaline activators will generate different chemical compositions, crystallinities and phases. Accordingly, the formation process of the impurity-contained geopolymer (i.e., silico-aluminophosphate [SAP] geopolymer and alkalialuminosilicate [AAS] geopolymer) can be illustrated based on the acidic and alkaline activation conditions as shown in **Figure 1**. In addition to the different pH conditions, the resulting molecular structures of the SAP and AAM geopolymers vary to wide extent, especially the chemical environment of Al in the formed gels, which behaves as Al(VI) and Al(IV) in SAP geopolymer and AAS geopolymer, respectively. Such formation model and mechanism will be discussed in section Formation Model and Mechanism of Geopolymer.

Recently, the AAS geopolymers have been investigated from multiple angles, such as the synthesis mechanism and chemistry (Provis, 2013; Provis and Bernal, 2014; Provis et al., 2015), properties and durability (Bernal et al., 2014; Arbi et al., 2016; Ding et al., 2016), life cycle analysis (Habert et al., 2011; Ouellet-Plamondon and Habert, 2015), and multi-field applications (MacKenzie, 2015; Rao and Liu, 2015; Luukkonen et al., 2019). Yet, the available review related to the SAP geopolymer is limited even though such type of geopolymer was termed as aluminosilicate phosphate cement (Khabbouchi et al., 2017; Katsiki, 2019), phosphoric acid-based geopolymer (Liu et al., 2012; Guo et al., 2016), phosphate-based geopolymer (Wang et al., 2017), acid-based geopolymers (Mathivet et al., 2019), and phosphate-bonded materials (MacKenzie, 2015). This paper mainly reviews the development of the SAP geopolymers in reference with the conventional AAS geopolymers.

## TERMINOLOGY AND CLASSIFICATION

The research era of alkali-activated materials (AAMs) started in 1908 (Kühl, 1908; Provis and Van Deventer, 2013), and since then AAMs were explained as cement-like materials formed by the reaction between an aluminosilicate precursor and an alkaline activator. Then, the tranquility was broken by the term "geopolymer" coined by Davidovits who inadvertently produced an inorganic polymer-like material with good fireresistance (Davidovits, 2011). The traditional geopolymer was predominantly confined to the binder prepared by the aluminosilicate sources with marginal impurities. Such geopolymer binder mainly highlights the chemical behaviors of aluminate and silicate species, including bonds breaking (i.e., decomposition reaction) and remaking (i.e., polymerization reaction). Afterwards, the activation approach of aluminosilicate precursors was extended to phosphate species and termed as phosphate-based geopolymer (Wagh, 2004; Davidovits, 2011). The empirical formulas of the AAS and SAP geopolymers are shown in Equations (6, 7), respectively (Cui et al., 2008; Davidovits, 2011; Wagh, 2016).

$$M\_n[-(\text{SiO}\_2)\_z - AlO\_2]\_n \cdot mH\_2O \tag{6}$$

$$[-(\text{Si}-O)\_z - Al-O-P]\_n \cdot mH\_2O \tag{7}$$

where n is the degree of geopolymerization, z is 1, 2, or 3, and M is an alkali cation (e.g., K or Na).

Recently, the term "chemically-activated materials" (CAMs) was also proposed to describe the aluminosilicate sources that can be activated to create cementitious matrices using some chemical solutions (e.g., alkali-silicate, phosphate and other chemicals) (CAM2017 International Conference, Gold Coast, Australia). However, the term "geopolymer" seems to be more attractive to (even widely-used by) engineers because such term is more concise and easy-spreading. In this paper, all the binders primarily prepared by aluminosilicate materials and some chemical activators were named as modern geopolymer for the convenient circulation in academic, engineering and business communities. The modern geopolymers are further classified into two categories according to the activation approach: alkali-aluminosilicate (AAS) geopolymer and silicoaluminophosphate (SAP) geopolymer. The terminology and development of the geopolymers are summarized in **Figure 2**. In modern geopolymer, the AAS geopolymer denotes the alkali activation of aluminosilicate precursors such as the calciumcontaining sources (e.g., alkali-activated GGBS binder), the lowimpurity aluminosilicates (e.g., alkali-activated FA/MK binder), or their blends (e.g., alkali-activated GGBS/FA blends), while the SAP geopolymer implies that the aluminosilicate sources are pre-mixed with some metals and activated by the acid phosphate or other phosphates. It is good to mention that the activation methods assisted by external energies [e.g., microwave (Chindaprasirt et al., 2013a), autoclaved (Rashad et al., 2012), and mechanical activation (Kumar and Kumar, 2011)] are not involved in this review.

## FORMATION MODEL AND MECHANISM OF GEOPOLYMER

## AAS Geopolymer

As commonly explained, the three-dimensional network of alkali-aluminosilicate geopolymer is configured with the negative [AlO4] <sup>5</sup><sup>−</sup> and [SiO4] <sup>4</sup><sup>−</sup> tetrahedrons in addition to positive alkali metal ions (e.g., Na<sup>+</sup> or K+) (Duxson et al., 2006; Buchwald et al., 2011). Recent research has indicated that the hydroxyl free radicals (·OH) released in the alkaline solutions catalyze the disintegration of the aluminosilicate precursor by breaking the bonds of Si-O-Si, Al-O-Si, and Al-O-Al, besides catalyzing the promotion of isolated aluminate and silicate anions reconstruction via remaking the bonds of Si-O-Si, Al-O-Si, and Al-O-Al (Feng et al., 2016). Thus, the tetrahedrons of silicon and aluminum are combined by oxygen bridges to form the high-connectivity chain or ring network structures. The negative charge given by tetra-coordinated aluminum is compensated by the alkali metal ions (Xu et al., 2001). However, the lowersilica gels in such geopolymer matrix (e.g., high molar ratio of Na2O/SiO<sup>2</sup> in the activating solution or low molar ratio of Si/Al in the aluminosilicate source) tend to be transformed into some loose particulate units [i.e., pseudo-zeolitic structure (mainly Q<sup>4</sup> )] rather than continuous well-connected N-A-S-H gels (He et al., 2013; Provis et al., 2015). It was found that a reasonable content of silicate species in the activating solution could yield compact alkali-aluminosilicate geopolymer matrix with more continuous gels, especially when using the solid aluminosilicate precursor with lower Si/Al ratio (Duxson and Provis, 2008).

The results of solid-state nuclear magnetic resonance (NMR) confirmed the chemical environment of silicon and aluminum in the calcium-free geopolymer made by lowimpurity aluminosilicate precursor (Buchwald et al., 2007). The

higher intensity of the Q<sup>4</sup> (4Al) and AlO<sup>4</sup> signals shown by NMR curves indicated that the tetra-coordinated silicate and aluminate species dominate the N-A-S-H gels formed in alkalialuminosilicate geopolymer (Xu et al., 2018).

The microstructure and strength developments of the lowcalcium FA-based geopolymer activated by an alkaline solution were descriptively explained beforehand (Fernández-Jiménez et al., 2005, 2006). When reacted with alkaline solution, the FA grains are chemically opened which dissolves the active silicate and aluminate phases to form microscopic imperfections (e.g., points or holes). The dissolved species are condensed into gels through geopolymerization and nucleation which result in compact geopolymer matrix with relatively adequate strength (Pacheco-Torgal et al., 2008). A three-stage corresponding relationship between microstructure evolution and geopolymerization is established to interpret the mechanical strength development in the AAS geopolymer (Fernández-Jiménez et al., 2006). The first stage is the "dissolution stage" of the glassy aluminosilicate phases. Since the matrix is not yet formed during this period, no mechanical strength can be achieved. The second stage is the "induction period," where massive metastable gels (denoted Gel 1 in Fernández-Jiménez et al. (2006)) are formed and precipitated. As the dissolution of silicate is comparatively slower than that of aluminate, the Gel 1 includes relatively high amount of reactive aluminum which is associated with the initial setting of the geopolymer paste. It is good to mention that the mechanical strength development is usually initialized at the second stage. The final stage is the "silicon incorporation stage" that describes the transformation of Gel 1 into Gel 2 (as indicated in Fernández-Jiménez et al., 2006) wherein the reactive aluminum is embedded into the Si-rich gels to form Gel 2. According to the aluminosilicate precursor and activator used, the mechanical strength of the formed matrix will usually be lower when the content of Gel 1 is higher than that of Gel 2 in hardened geopolymer matrix. Conversely, the mechanical strength can be higher in case the content of Gel 2 outnumbers that of Gel 1, as reported by Criado et al. (2008).

## Effect of "Impurities" in AAS Geopolymer

Regarding the role of calcium in alkali-aluminosilicate geopolymer, it is usually claimed that calcium aluminate silicate hydrate (C-A-S-H) gels with tobermorite-like structure (mostly Q2 and some Q<sup>1</sup> ) are generated during the geopolymerization (Richardson et al., 1994). The calcium source can induce the aluminosilicate precursor to produce different chemical products in AAS geopolymer. The chemical environments of silicon in calcium-incorporated AAS geopolymer behave as Q<sup>1</sup> and Q<sup>2</sup> (1Al) forms, while the Q<sup>4</sup> (2Al), Q<sup>4</sup> (3Al), and Q4 (4Al) chemical forms of silicon appear in calcium-free AAS geopolymer (Buchwald et al., 2007). The AlO4-dominated C-(A)-S-H gels and the hydrotalcite (AlO6) exist in the calcium-incorporated AAS geopolymer as the main chemical configurations of aluminum, which shows significant differences compared to calcium-free AAS geopolymer (AlO4-dominated N-S-A-H gels).

The combined utilization of GGBS and low-calcium FA can produce N-A-S-H and C-A-S-H gels which are entangled together to form strength-giving gels. The molecular structure of C-A-S-H gels in high-calcium AAS geopolymer can be generalized as the Al-substituted C-S-H gels (Richardson et al., 1993; Brough et al., 2001; Lodeiro et al., 2010). Such gels have Ca/Si ratio lower than that formed in Portland cement (Wang and Scrivener, 1995). This interpretation is similar to the chemistry of supplementary cementitious materials (SCMs) in OPC. The Si-rich phase improves the lower Ca/Si ratio of hydrates by consuming portlandite (i.e., pozzolanic reaction), while Al-rich phase promotes the Al-uptake of hydration products (Lothenbach et al., 2011).

Although the glassy aluminosilicate source with high calcium content can easily generate C-A-S-H gels in alkali environment, the reactivity of available calcium in raw materials may be varied substantially depending on its existence forms. Besides, the GGBS and high-calcium FA have similar calcium and glassy aluminosilicate contents, yet they behave differently in the alkali activation process. The high reactivity of GGBS is mainly attributed to the fast cooling technology (e.g., water quenching granulation) which generates amorphous phases with the network-forming anions [SiO4] <sup>4</sup>−, [AlO4] <sup>5</sup>−, and [MgO4] 6− in addition to the network-modifying cations Ca2+, Al3+, and Mg2<sup>+</sup> (Garcia-Lodeiro et al., 2015). The existence of calcium species can lower the polymerization degree of silicon network which can equip the GGBS with high pozzolanic activity. On contrary, the high-calcium FA particles usually form crystalline phases (e.g., mullite and quartz) due to the slow cooling collection method which presents minimal or no cementitious properties. Therefore, the use of high-calcium FA as main raw material for the preparation of alkali-aluminosilicate geopolymer is significantly challenging at room temperature.

In many existing studies on the durability and mechanical properties of AAS geopolymers regardless of "two-part" (i.e., liquid activators) or "one-part" (i.e., solid activators) mixing approaches, the use of GGBS and FA combination with a certain ratio as an aluminosilicate precursor achieves a geopolymer binder that can roughly meet the engineering requirements (Ismail et al., 2014; Bernal et al., 2015, 2016; Alrefaei and Dai, 2018; Alrefaei et al., 2019). The incorporation of highcalcium materials (i.e., to seed calcium species) with low-calcium raw materials (i.e., main source of alumina and silica) greatly affects the formation of the main binding gels (Bernal et al., 2016). The N-A-S-H type gels generated in low-calcium systems exhibit a relatively slower reaction rate compared to the C-A-S-H type gels in high-calcium systems; however, N-A-S-H gels promote highly cross-linked and compact products through extending the curing time (Bernal et al., 2016). In other words, the synergistic effect of the high- and low-calcium aluminosilicate precursors is witnessed in the coexistence of these molecular structures, where the C-A-S-H gels ensure early performances (e.g., setting time and early strength), while the N-A-S-H gels contribute to long-term characteristics (e.g., durability and strength development).

## SAP Geopolymer

The Egyptian pyramids proved their reliability and durability for thousands of years although the construction materials used in the Pyramids are not fully clear (MacKenzie et al., 2011). However, limited research declared that the surface of pyramids' blocks spontaneously grew a white/red coating that was complicated geopolymeric phosphate compounds made of crystalline hydroxyapatite and brushite embedded in a poly(sialate)/SiO<sup>2</sup> matrix (Barsoum et al., 2006; Davidovits, 2011). Later on, the aluminosilicate precursors consolidated by phosphates or orthophosphoric acids were explored for preparing refractory castables (Kingery, 1950, 1952) which can be defined as phosphate activation of the aluminosilicates to yield an SAP geopolymer with excellent durability and high temperature resistance (Wagh, 2016). The use of phosphate or phosphoric acid as an activator allows the chemical interaction with the aluminosilicates to produce geopolymeric binder with condensed poly(silico-alumino-phosphate) phases (Cao et al., 2005; Liu et al., 2010).

Regarding the molecular structure in SAP geopolymer, the [AlO4] <sup>5</sup>−, [SiO4] <sup>4</sup>−, and [PO4] <sup>3</sup><sup>−</sup> units act as basic blocks (He et al., 2013). Compared with AAS geopolymer, the [PO4] 3− unit replaces the alkali metal ion, as reported by Cui et al. (2011). Cui's model revealed that partial replacement of [SiO4] 4− unit by [PO4] <sup>3</sup><sup>−</sup> in the broken Si-O-Al bonds may occur in low polymeric network. Thus, the charge balance within the molecular structure can be achieved without the involvement of monovalent cations. Previous research showed that the updated Si-O-Al linkage (i.e., Al-O-P) can balance the charge due to the aluminum chemical shift ranging from octahedral or/and pentahedral coordination to tetradentate form (Cao et al., 2005). However, many studies indicated that the hexa-coordinated aluminum dominates in SAP geopolymers (Davidovits, 2011; Louati et al., 2014, 2016). It should be noted that an amorphous Si-Al-P structure was reported in AAS geopolymer wherein the tetrahedral P was identified to occupy a proportion of site of silicate (Mackenzie et al., 2005).

Since the model of microstructure evolution of SAP geopolymer was barely reported in previous literature, some result-derived interpretations and hypotheses are introduced in this paper based on existing studies (Duxson et al., 2006; Cui et al., 2008, 2011; Wagh, 2016). Cui et al. (2011) synthesized SAP geopolymer using pure Al2O3-2SiO<sup>2</sup> powders (synthesized in the lab) and phosphoric acid with P/Si molar ratio of 1.2. The XRD patterns showed no clear trace of newly-formed crystalline peaks in the phosphate activated geopolymers. Instead, a broad diffused peak in the range of 15–40◦ (2θ) appeared in all the samples. The diffraction peak was shifted from 22◦ (2θ) in calcined Al2O3- 2SiO<sup>2</sup> powder to about 26◦ (2-theta) in geopolymers, with a marginal change in the diffraction intensity. It was claimed that the molecular structure and interlayer spacing were chemically shifted after geopolymerization in acid phosphorus environment relative to the raw Al2O3-2SiO<sup>2</sup> powders.

Besides, the natural aluminosilicate materials like Tunisian clay and metakaolin were also investigated when activated by phosphate acid with various Si/P molar ratios (Douiri et al., 2014; Louati et al., 2014). A shift in the wide broad band from 18–30◦ (2θ) to 22–35◦ (2θ) was observed in both calcined clay and synthesized geopolymers, as well as the diffraction intensity changed in some crystalline peaks e.g., aluminum phosphate (AlPO4), augelite (Al2(PO4)(OH)3), and monetite (CaHPO4). Such differences in the XRD patterns of the raw aluminosilicates and phosphate activated aluminosilicates (i.e., SAP geopolymer) could be attributed to the geopolymerization of the aluminosilicate precursors.

Perera et al. (2008) also performed a comparative study on the metakaolin-based geopolymer when activated by alkali and phosphate solutions. Similar to Cui et al. (2011) observations, the XRD results confirmed the amorphousness of SAP geopolymer. Further investigations by the transmission electron microscope (TEM) equipped with energy dispersive X-ray spectroscopy (EDXS) and selected area electron diffraction (SAED) indicated that such amorphous phases contained Si, Al, P, and O. The SAED image exhibited some blurred rings and spots instead of crystal lattices which further ensured the availability of amorphous phases in SAP geopolymer (Perera et al., 2008).

Furthermore, magic angle spinning-nuclear magnetic resonance (MAS-NMR) analysis was conducted on the SAP geopolymer by Perera et al. (2008). A single [PO4] <sup>3</sup><sup>−</sup> resonance was observed at around−16 ppm which could be deconvolved into two resonances (−4 and −16 ppm) due to its asymmetric line shape. A cross polarization (CP) NMR of P was identified to enhance the dominance at −16 ppm. The <sup>27</sup>Al results indicated that the main chemical environment of aluminum in SAP geopolymer was six-coordination (VI) as evidenced by the single resonance at an apparent shift of around −12 ppm. Thus, the phosphate activation promoted the metakaolin to produce mixed octahedral Al(VI) environment with both P and Si, which influenced the oxo-bridged second coordination sphere. This was mainly unlike the AAS geopolymer where Al(VI) was converted to tetrahedral sites, i.e., Al(IV). The three resonances (e.g., −91.8, −101.6, and −110.8 ppm) in <sup>29</sup>Si NMR curve described the Q<sup>3</sup> , Q<sup>4</sup> (Al or P), and Q<sup>4</sup> silicon chemical environments, respectively (Provis et al., 2015). The intensity of Q<sup>3</sup> was the lowest which meant that most of metakaolin was dissolved. The mixed environment of Si-Al-P phase in SAP geopolymer was assigned as the Q<sup>4</sup> (Al or P), while the Q<sup>4</sup> was described as siliceous or Al-free portions that might not fully connect to P phase (Cao et al., 2005).

## Effect of "Impurities" in SAP Geopolymer

In order to improve the modest early performance of SAP geopolymer, the incorporation of some alkali metals as setting agents (e.g., solid Mg/Ca source) with aluminosilicate precursors can produce a setting-controlled SAP geopolymer matrix (Luz et al., 2015). Such agents essentially initiate acid-base chemistries (i.e., between phosphate and setting agents) in addition to geopolymerization of the aluminosilicates, which enables the fast formation of amorphous or crystalline phosphates to modify the workability and early strength of such geopolymer (Wang et al., 2018b). Relative to the aluminosilicate precursors, the added setting agents can be regarded as "impurities" here. The chemistry of such setting agents is elaborated below.

"Phosphate cement" is a general term that describes the phosphate as a main ingredient of the cement (Walling and Provis, 2016). Since phosphate cement is usually a binder that hardens through acid-base and/or hydrolysis chemical reactions (usually generate crystalline phosphates as main products) at room temperature, it is also named as chemically bonded phosphate ceramic (CBPC) (Roy, 1987). CBPCs were discovered and developed as dental cements in the nineteenth century which were extended to construction field later on as patching materials for rapid rehabilitation of runway and bridge. The formation of CBPC experiences the dissolution of alkali metals and phosphates to form crystalline structures (i.e., phosphate compounds) with excellent adhesive and mechanical properties. It is expected that such chemical mechanism (i.e., the chemistry between alkali metals and phosphates) can be introduced into SAP geopolymer as a room temperature hardening mechanism to address its weak early properties (Wang, 2018). The reaction products of CBPCs are directly linked to the raw materials used; more specifically, the material recipe (e.g., ratio, reactivity etc.) is related to the reaction environments (e.g., pH condition, ions concentration etc.). Here, the calcium- and magnesium-based sources are comprehensively introduced because other alkali-metal sources (e.g., Fe and Zn) fails to form well-crystallized phases due to the violent reaction process (Wagh and Jeong, 2003a,b).

Calcium-based CBPC is often termed as "calcium phosphate cement" (CPC) which is used as a bioceramic material in orthopedic industry due to its excellent bioactivity and biocompatibility (Ben-Nissan, 2014). The chemical reactions involved during the CPC setting consist of three stages: dissolution, supersaturation, and precipitation (Zhang J. et al., 2014). In acidic phosphate solution, the calcium donors slowly release the calcium ions and react with (hydro)-phosphate groups to form a supersaturated solution (i.e., paste). When reaching the critical values (e.g., ions concentration or pH condition), the nucleated and precipitated new phase grows along with the continuous dissolution of calcium. According to the composition of final products, the CPC is classified into apatite cement [βtricalcium phosphate, Ca5(PO4)3(OH/Cl/F) or Ca3(PO4)2] and brushite cement (CaHPO4·2H2O) (Tamimi et al., 2012; Ben-Nissan, 2014). In general, according to the recipe (e.g., solubility and Ca/P ratio) and reaction conditions (e.g., pH value and hydrothermal environment), the possible chemical processes involved in the formation of CPC matrix can be described in Equations (8–10).

$$\begin{aligned} \left[3Ca^{2+} + 2[H\_nPO\_4]^{n-3} \rightarrow Ca\_3(PO\_4)\_2 + 2nH^+ \\ \left[5Ca^{2+} + 3[H\_nPO\_4]^{n-3} + H\_2O \rightarrow Ca\_5(PO\_4)\_3\right](OH) + (3n+1)H^+ \end{aligned} \tag{9}$$

$$\text{Ca}^{2+} + \text{[}\text{H}\_{n}\text{PO}\_{4}\text{]}^{n-3} + \text{2}\text{H}\_{2}\text{O} \rightarrow \text{CaHPO}\_{4}.2\text{H}\_{2}\text{O} + (n-1)\text{H}^{+} \tag{10}$$

Fresh CPC is a castable paste which is preferable for bone defects reconstruction with minimum invasive surgery since it can be resorbed under physiological conditions after embedding into the body (Xue et al., 2012; Inzana et al., 2014; Xu et al., 2017). Thus, the moldable properties of calcium-blended SAP geopolymer extend its applications to 3D printing for building arts in civil engineering or other aesthetic fields.

Magnesium phosphate cement (MPC) is another typical CBPC that is widely used in civil engineering because of its excellent performance (Wang and Dai, 2017). In the early research stage, the MPC was fabricated with magnesia and phosphoric acid (Wagh, 2016). The lower pH environment induces a rapid and highly exothermic reaction where the reaction product is water-soluble magnesium dihydrogen phosphate [Mg(H2PO4)2·nH2O]. Subsequently, attempts were made to use less acidic diphosphate. The employment of diphosphate and calcined magnesia further mitigates the reaction rate and heat release during MPC formation. The soluble diphosphates can create an acidic environment in water which facilitates the dissolution of magnesia. When the concentration and pH are approaching ideal value, crystal generation and growth occur in the supersaturated slurry (Chauhan et al., 2008). The general reaction of MPC is described in Equation (11). The resulting phase is crystalline struvite or its analogs (e.g., struvite-K or struvite-Na) with 6–8 crystal waters (Chauhan et al., 2011; Chauhan and Joshi, 2014). Such hydrates have excellent cementitious and mechanical properties (Ding et al., 2012). Finch and Sharp (1989) also studied the reaction of mono-aluminum phosphate and magnesia in terms of molar ratio and final products, showing that the crystalline Newberyite (MgHPO4·3H2O) appeared instead of struvite analog in case of the Mg/P molar ratio of 4. The phosphate species might involve in the formation of amorphous phase as well.

$$\text{MgO} + \text{XH}\_2\text{PO}\_4 + n\text{H}\_2\text{O} \rightarrow \text{MgXPO}\_4 \cdot (n+1)\text{H}\_2\text{O} \text{ (11)}$$

The characteristics of crystalline calcium/magnesium phosphate phases including mineral name, chemical formula and main application fields are listed in **Table 1**. Ca-based apatites are mostly used as biomaterials for dental applications and mimic bones due to their good biocompatibility (Xu and Simon, 2005; Giocondi et al., 2010). The variants of Ca-based apatites can also be employed to immobilize the heavy metal pollutants (e.g., Pb, Cd, Zn, or Cr) by embedding them into the respective apatite minerals (e.g., Ca/Zn replacement) (Saxena and D'Souza, 2006). On the other hand, the final products containing struvite or its analogs (i.e., Mg series) are usually used as rapid repair materials in civil engineering due to their high mechanical performance and adhesive property (Ma et al., 2014).

## Model Comparison of AAS and SAP Geopolymers

Alkali activation process of Ca sources (i.e., C-A-S-H gel formation) and acid-base reaction process of Ca/Mg and phosphate (i.e., amorphous or crystalline phosphates) can promote both rapid setting at room temperature and strength upscaling by forming new phases. Their simplified synthesis models of the two "Hybrid" geopolymers are conceptualized in **Figure 3**.

**Figure 4** shows a comparative flow chart between the formation processes of AAS and SAP geopolymers wherein


the key stages from the raw materials to the final geopolymer binders are outlined. The Ca-containing AAS geopolymer was investigated for many years (as shown in **Figure 4A**), including dissolution, speciation, gelation, reorganization and final polymerization (Duxson et al., 2006). The chemicallydissolved aluminosilicate precursors accessibly form the C-(A)- S-H gels when encountering the free Ca species. This chemical process is relatively faster than that of N-A-S-H gels at ambient temperature. In most cases, the two gels co-exist in the final geopolymer matrix as resulting products.

As a comparison, the synthesis process of SAP geopolymer is simplified in **Figure 4B**. However, the research related to SAP geopolymer is rare, so it should be noted that the synthesis process of such geopolymer is deduced according to the classical interpretations on the CBPC and the AAS geopolymer. The transformation of phosphate species from sol state to gel state, including the dissolution and speciation of (hydro)-phosphate ions, is the first step during SAP geopolymer formation (Perera et al., 2008). Usually, the acid-base reaction between Ca or Mg sources and (hydro)-phosphate ions is initiated before the dissolution of aluminosilicate precursors. The gelation and reorganization among aluminate, silicate and phosphate form S-A-P gels via the condensation processes. Final polymerization of these species gradually synthesizes the SAP geopolymer matrix. Unlike AAS geopolymer, phosphate-bearing crystalline or amorphous phases may occur in some cases when reasonable molar ratios (metal ion to phosphorus) and pH environment are available (Guo et al., 2016).

## ENVIRONMENTAL IMPACT

## Portland Cement and Geopolymer Binders

Due to the vast of mineral resource consumptions, intensive energy usage and high carbon emissions companioned with ordinary Portland cement (OPC) production, it is believed that the geopolymer binders should be strongly proposed as green building materials to partially replace cement in future construction applications (Walkley et al., 2017). Usually, the life cycle assessment (LCA) methodology, also known as the "cradle to grave" approach, is internationally normalized to provide quantitative information for the environmental burden of each product manufacture system (ISO, 2006; Passuello et al., 2017). However, the recent findings related to the LCA-based comparison between geopolymer and OPC products showed a wide variation, even some conflicting results, depending on the goal, scope, the needs, and the targeted audience used in LCA analysis (Ouellet-Plamondon and Habert, 2015).

The aroused controversies regarding the LCA of geopolymer binders are derived from several complexities such as the local resources availability, the process of raw materials, thermal treatment and the mix design parameters (e.g., alkali content, Si/Al molar ratio, water-to-binder ratio etc.), as shown in **Figure 5** (Duxson et al., 2007; Habert et al., 2011; Ouellet-Plamondon and Habert, 2015; Habert and Ouellet-Plamondon, 2016). Another challenge of the LCA application in geopolymer binders is that such binders mainly include two different industrial products (i.e., aluminosilicate precursors and activators) unlike the OPC that deals with a single material industry. When the solid precursor is treated as a waste material in the LCA analysis (e.g., in Europe, Directive 2008/98/EC of the European Parliament and of the Council on waste and repealing certain Directives. L312: 3–30), the energy consumed during the production process (i.e., FA and GGBS from power stations and steel production, respectively, as shown in **Figure 5**) does not negatively impact the CO<sup>2</sup> balance of the geopolymer binders. As a comparison, the cement clinker calcination consumes a large quantity of fossil fuel energy and releases additional CO<sup>2</sup> through the decarbonation reaction. Thus, the geopolymer binders are considered as a groundbreaking solution for the environmental issues caused by Portland cement industry (Duxson et al., 2007). On the other hand, such binders adversely impact the environment when the precursor is considered as a raw material since the thermal input of the solid precursor should be involved in the LCA calculation. Besides, Habert et al. (2011) claimed that the production of a typical AAS

geopolymer concrete presents a slightly lower impact on the global warming compared to standard OPC concrete, mainly due to the energy expenditure of activators (e.g., mining, treatment, manufacture, and transport). Regarding the SAP geopolymer, it was reported that the energy usage during the production of phosphate binders (mainly the phosphate products) accounts for around one-fourth of that consumed in Portland cement industry (Wagh, 2016).

## AAS and SAP Geopolymers

When comparing the two types of geopolymer (i.e., AAS and SAP), the key point lies in the environmental impact of activating solutions. During the preparation of AAS geopolymer, it is necessary to use soluble sodium silicates to address the performance requirements (e.g., mechanical properties; Ding et al., 2016) of geopolymer concrete, especially when aluminosilicate precursors with lower Si/Al ratio is employed (Krizan and Zivanovic, 2002). However, the silicate solutions are always contentious due to their high cost and aggressiveness to the environment during production (e.g., intensive energy consumption and global warming potential) (Fawer et al., 1999). On the other hand, for SAP geopolymer, the phosphate solutions are majorly manufactured through the reaction between phosphate rock and sulfuric acid at <100◦C. Thus, the environmental impacts of the phosphate solutions are largely determined by the sulfuric acid production industry (Kongshaug, 1998). Generally, modest emissions (e.g., carbon dioxide) are linked to the modern sulfuric acid production industry. Besides, the exothermic reactions involved in such production may generate net energy (i.e., heat released) that can be totally employed in other industries (Wood and Cowie, 2004). Further, emission standard of pollutants for sulfuric acid industry is severely implemented all over the world (China Emission Standard of Pollutants for Sulfuric Acid Industry, 2011; EU Best Available Techniques for Pollution Prevention and Control in the European Sulphuric Acid and Fertilizer Industries, 2000; US Guidelines for Limitation of Contact Sulfuric Acid Plant Emissions, 1971). Accordingly, the phosphate solutions seem to be more sustainable and environmentally friendly relative to the alkali-silicate solutions in terms of the gas emission and energy expenditure.

In addition to their moderate carbon emissions and energy consumption, the application of the SAP geopolymer products may form a virtuous eco-system as the phosphorus available in such binders can be absorbed or extracted to form the Pcontaining fertilizers for crops after abandoning (Bartos et al., 1991; Raven and Loeppert, 1996). These biological and chemical actions may further form mineral deposits (e.g., phosphate rocks). Such eco-friendly loop makes the phosphate-based products greener and more sustainable which meets the concept of industrial ecology in the long run (Jelinski et al., 1992). Adversely, the alkali-silicate activators (varying modulus of sodium silicate) used for preparing the AAS geopolymer products are usually non-recyclable. Some researchers are exploring some waste-derived activators (e.g., silica fume and rice husk ash) and non-silicate activators (e.g., carbonate and sulfate)

to replace commercial silicate solutions for AAS geopolymer preparation (Nazari et al., 2011; Bernal et al., 2012; He et al., 2013), which may reduce environmental burden caused by waterglass industry.

## PERFORMANCE, DURABILITY, COST, AND ACCESSIBILITY

## Compressive Strength

The mechanical strengths reported in the published literature varied widely owing to the differences in raw materials (e.g., attribute and recipe), curing condition (e.g., temperature and moisture) and sample preparation technology (e.g., size and age). The 28 days compressive strengths of the geopolymer obtained from literature survey, including low- and high-calcium AAS geopolymers in addition to SAP geopolymers, are shown in **Figure 6** (Davidovits, 2011; Chindaprasirt et al., 2012, 2013b; Nath and Kumar, 2013; Nematollahi and Sanjayan, 2014; Ati¸s et al., 2015; Guo et al., 2016; He et al., 2016; Reddy et al., 2016; Wang et al., 2016; Wang, 2018). Clearly, the statistics showed that the compressive strengths of the AAS geopolymer (i.e., 15– 65 MPa for low-calcium series and 70–110 MPa for high-calcium series) significantly outperformed that of the SAP geopolymer (i.e., 10–50 MPa).

Ding et al. (2016) summarized the mechanical properties of AAS geopolymer concrete, revealing that the compressive strength and elastic modulus can moderately meet the requirements of Portland cement concrete design codes. Reddy et al. (2016) made an attempt to investigate the influences

of oxide composition of precursors on the compressive strength of the AAS geopolymer products. Perera et al. (2008) synthesized (60◦C curing for 24 h) SAP geopolymer with compressive strength approaching 140 MPa which is two-fold that of AAS geopolymer. On the other hand, Wang et al. (2017) reported that the final setting time was more than 48 h and the compressive strength reached 31 MPa at 28 days for the metakaolin (MK)-based geopolymer activated by mono-aluminum phosphate at room temperature curing. Guo et al. (2016) obtained the phosphate-based geopolymer using MK and disused phosphate liquid which resulted in a compressive strength of 62–67 MPa when the matrix experienced 7 days curing at 60◦C. The activating effect of disused phosphate liquid exceeded that of phosphoric acid because of the presence of aluminum ions and thermal curing.

## Thermal Stability, Dielectricity, and Efflorescence

As previously reported, the MK-based geopolymers activated by phosphoric acid exhibited good thermal and volume stability (Liu et al., 2012). Phase transition (from aluminum hydrogen phosphate to berlinite) was observed when the geopolymer sample was subjected to elevated temperatures (900–1,550◦C). In the AAS geopolymer, the presence of calcium source could reduce the curing cost and increase early strength as previously explained. The early compressive strength of such hybrid geopolymer matrix is augmented along with the increased use of GGBS (Guerrieri and Sanjayan, 2010). However, such practice lowers the ability of geopolymer to resist fire and high temperatures (Buchwald et al., 2007) due to the existence of the non-fire-resistant C-(A)-S-H gels.

Cui et al. (2011) compared the dielectric loss of two types of geopolymers. It was found that the SAP geopolymer held a very low dielectric loss (around 0.01 for frequencies more than 300 MHz) at 300◦C dried for 2 h. On the other hand, the AAS geopolymer was temperature-dependent due to the availability of free metal and hydroxyl ions; thus, it was difficult to decrease the dielectric losses resulting from ion transfer.

Besides, the efflorescence is also a serious issue with relatively little attention in research (Najafi Kani et al., 2012; Allahverdi et al., 2015; Pacheco-Torgal, 2015). Efflorescence in AAS geopolymer is the result of the incomplete consumption of alkaline and/or soluble silicates that leads to form sodium carbonate in the pores or on the surface of geopolymer matrix (Xue et al., 2018). This can be because the presence of water weakens the bond of sodium in geopolymer; besides, the efflorescence rate is strongly activation-dependent. In case of the same alkali content and curing temperature, NaOHactivated geopolymers show less and slower efflorescence relative to geopolymers activated by sodium silicate. Geopolymers with high Na2O/Al2O<sup>3</sup> ratio may suffer from unsightly efflorescence as well (Najafi Kani et al., 2012). The use of calcium blended precursors (i.e., FA/GGBS-based geopolymers) in addition to the hydrothermal curing process usually generates geopolymer binders with relatively weaker efflorescence, yet the GGBS addition appears to delay rather than mitigate the effect of efflorescence (Zhang Z. et al., 2014). Najafi Kani et al. (2012) stated that the alumina-rich admixtures can reduce the efflorescence. Therefore, the well-designed mix proposition and chemical composition of the geopolymer binders can mitigate or even eliminate the efflorescence, e.g., providing hydrothermal curing, sufficient Al-rich mineral admixtures for gel formation and alkali consumption (Allahverdi et al., 2015). On the other hand, SAP geopolymers have high resistance to efflorescence relative to the AAS geopolymers; thanks to the acidic synthesis environment.

## Cost and Availability

In addition to the performance comparison, the cost and the availability of both types of geopolymers (i.e., AAS and SAP) are also crucial to practical engineering. Although quotations form regional suppliers (e.g., Alibaba) indicate that the price of phosphoric acid is more than three-fold that of sodium silicate, the SAP geopolymer can act as a complement to AAS geopolymer in the field of geopolymer category. In other words, the SAP geopolymer shall be further explored as small-scale or specifically-required functional materials (e.g., as fire-proof material, dielectric ceramics, coating materials, fast-rapid materials, and 3D printing materials), while the goal of AAS geopolymer is to reduce the traditional cement consumption.

The complexity of aluminosilicate sources, especially aluminosilicate-containing industrial byproducts, is a major concern for geopolymer production. Fly ash as a typical aluminosilicate source for geopolymer preparation is an anthropogenic material with polymorphic, polycrystal, and multi-component system (Blissett and Rowson, 2012). According to different regions, more than 316 individual minerals and 188 mineral groups in coal fly ash have been identified and characterized, which may cause even more differences in non-crystalline phases and particle characteristics (Vassilev and Vassileva, 2005). The diversity of raw materials results in wide differences in the geopolymer mechanical properties, even when using the same formula and mix design. Therefore, if the aluminosilicates are properly characterized, classified and pretreated, the properties of geopolymer products can be tailored for certain applications and performances. The geopolymer turns out to possess multi-phase, multi-component, and multi-morphology; however, its performance is strongly reliable on the types of raw materials. In order to achieve a high performance geopolymer matrix, the following three assessment indices of aluminosilicate are required for preparing different geopolymers (Luxán et al., 1989; De Rojas and Frías, 1996; Agarwal, 2006).

(1) Material characterization: The characterization indices of aluminosilicates include their compositions (oxide and phase) and physical features (particle size, specific gravity, surface, and loss of ignition), which can greatly influence the performances of geopolymer products.

(2) Reactivity test: The reactivity of the aluminosilicate precursors is closely associated with their phases. The combined use of Rietveld analysis (Hill and Howard, 1987) and PONKCS (partial or no known crystal structure; Scarlett and Madsen, 2012) are relatively useful for initial gauging of the crystalline and amorphous phase contents. Glassy phase composition of some aluminosilicates can be captured by SEM-EDXS with multispectral image analysis, to differentiate the inert crystalline and the reactive amorphous phases in case of similar elemental composition (Chancey et al., 2010; Durdzinski et al., 2015). The participation degree of aluminosilicates during geopolymerization heavily determines the properties of geopolymer products.

(3) Performance enhancements: For improper aluminosilicate sources, particle size reduction and minerals addition seem to be valid. However, the thermal and chemical pretreatments can be more effective because such pretreatments increase the amorphous proportion and decreases the loss of ignition (Handoo et al., 2002; Škvára et al., 2009; Bentz et al., 2012).

## POTENTIAL APPLICATIONS

## Application in Construction

One of the large-scale applications of geopolymers (mainly AAS geopolymer) is in construction due to their good mechanical behavior and durability. Some modified AAS geopolymer binder has been commercially branded for infrastructure projects, such as PYRAMENT cement (Davidovits, 2011). Currently, the SAP geopolymer is being researched in labs, especially with the focus on the unclear chemistries between phosphate and aluminosilicate. In cases of high-volume content of FA-blended phosphate binder, FA particles are conventionally considered as a diluent or inert filler that modifies the aesthetics of the formed products (e.g., change color and reduce porosity) to be comparable with Portland cement as a rapid healant. Recently, some reports have conversely indicated that the high mechanical strength of the FA-blended phosphate binders is also attributed to the formation of secondary amorphous phase (e.g., phosphosiloxonate geopolymeric phases; Gardner et al., 2015).

## Application in Waste Management

The waste production is directly related to population increase. In order to stop the contaminants from flowing into the food chain of human, some effective and low-cost amendments are applied to stabilize/solidify the hazardous or radioactive waste streams (Morrissey and Browne, 2004; Temuujin et al., 2014). According to the definition of stabilization/solidification, stabilization process is defined as the chemical converting of contaminants into their insoluble, immobile or nontoxic forms without spillage, leakage, and disintegration in a long term, while solidification is defined as changing the physical consolidating sludges, liquids, or powders into solid forms so that they will not be dispersed during transportation, long-term storage or disposal. Recently, AAS geopolymer binders are used for the remediation of heavy metals-contaminated soils through stabilization/solidification (S/S) (Luna Galiano et al., 2011). It is claimed that the heavy metals can be fixed in the molecular network during geopolymerization, either chemically through embedment into the molecular structure for charge balance or physically by being trapped in the surrounding 3D network (Van Jaarsveld et al., 1997). In SAP geopolymer, the phosphate ions are also chemically combined with detrimental metal species to form highly insoluble phosphates (e.g., pyromorphite during Pb S/S; Wang et al., 2018a).

## Application in Biomaterials

Considering the biocompatibility, resorbability, and customization of living body, the SAP geopolymer with Ca-containing setting agent is considered as an excellent suitable candidate. The tailor-made shape can be pre-prepared using 3D printing technique due to the good injectability of such materials (Chia and Wu, 2015). Except for precast skeletal structure, SAP geopolymer can be injected as a paste into selected regions in the biological body. Not only that paste cements the adjacent tissues, but also has less intrusion properties compared to the implants made of hardened biomaterials. Since bones and teeth mainly contain calcium phosphate compounds, Ca-containing SAP geopolymer can provide the necessary compositions for better biocompatibility. The zinc and magnesium species enable the performance enhancement (e.g., strength) in calcium enriched mixture (CEM) matrix as blended phosphate biomaterials (Utneja et al., 2015). Besides, the viable biocompatible AAS geopolymers have been demonstrated (MacKenzie et al., 2010).

## Application in 3D Printing

3D printing is a promising technique that can rapidly manufacture products with spatial complicity and precise distribution (Bassoli et al., 2007; Rengier et al., 2010; Lim et al., 2012). Such attractive rapid prototyping (RP) is based on the additive manufacturing (AM), defined by ASTM International as "the process of joining materials to create objects from 3D model data, usually layer upon layer" (ASTM F2791-10). The customized imaging database (e.g., computed tomography scans) can be visualized via 3D printing system. Nevertheless, how to select the appropriate material with appropriate rheology, setting time, temperature, strength, durability, and other special requirements, is the major problem. Recently, the 3D printing of biomaterials has been applied for surgical planning and prosthetics using phosphate materials due to its biocompatibility in human body (Guvendiren et al., 2016). The "two-part" geopolymer (aluminosilicate and activator) is also potentially applicable for 3D printing in construction (Xia and Sanjayan, 2016; Zhong et al., 2017).

## CONCLUDING REMARKS

This paper provides a comprehensive review and comparison between alkali and phosphate activations of aluminosilicate materials. According to the activation mode, the chemically activated aluminosilicates can be termed as alkali-aluminosilicate (AAS) geopolymer or silico-aluminophosphate (SAP) geopolymer, where both of geopolymers need extra promotions (e.g., thermal curing or setting agent) to obtain reasonable engineering properties. The thermal curing is usually inappropriate for in-situ applications, while the incorporation of fast-setting chemistries (induced by setting agents) during geopolymerization can easily facilitate an ambient-temperature setting mechanism in geopolymer preparation. Just like the formation of C-(A)-S-H gels in AAS geopolymer, the SAP geopolymer can produce crystalline or amorphous phosphate phases during geopolymerization when metal ions are incorporated with the aluminosilicates. The different activation modes generate various potential values of these "hybrid" geopolymers in a wide range of research fields. According to the performance, cost and accessibility, the two types of geopolymers are mutually supplemented in terms of function and application fields.

## AUTHOR CONTRIBUTIONS

Y-SW collected the data and wrote the paper. YA wrote some parts of the paper besides linguistic and grammar editing. J-GD supervised and finalized the paper.

## REFERENCES


## ACKNOWLEDGMENTS

The authors would like to acknowledge the financial support received from the Hong Kong-Guangzhou Technology and Innovation Partnership Program (Project No. 201807010055), National Science Foundation of China (NSFC) (Projects No. 51638008 and 51478406), HKSAR Innovation Technology Fund (Project code: ITS/009/17) and the Hong Kong Ph.D. Fellowship Scheme (HKPFS) awarded to the second author. The authors would also like to thank Prof. John L. Provis for his advice toward the paper during the study visit of the first author at the University of Sheffield.


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106th Annual Meeting of the American Ceramic Society (Indianapolis, IN), 107. doi: 10.1002/9781118408353.ch10


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Wang, Alrefaei and Dai. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Waste Stream Porous Alkali Activated Materials for High Temperature Application

#### Diana Bajare, Laura Vitola\*, Laura Dembovska and Girts Bumanis

*Department of Building Materials and Products, Riga Technical University, Riga, Latvia*

The growing interest of using industrial waste as recycled raw materials for the production of new, innovative materials is associated with effective use of natural resources and circular (zero-waste) economy. The research object is waste stream materials coming from chemical and processing industries, such as aluminum scrap recycling waste, chamotte-like precursor, firebricks sawing residues, and their use in production of high-temperature resistant, porous insulation materials by using alkali activation technique with 6 M NaOH solution. Adding aluminum scrap recycling waste to the composition of the tested alkali activated materials (AAM) contributed to the porous structure of the material with the pore size ranging from 1,000 to 5,000µm (detected by Micro-XCT, SEM). Lightweight (350–850 kg/m<sup>3</sup> ) and heat-resistant (up to 1,000◦C) AAM with compressive strength from 1.0 to 3.0 MPa was obtained. The mineralogical composition of the obtained AAM was detected (XRD) and the heat resistant minerals in the structure of AAM were identified. It was concluded that the increased amount of Al2O<sup>3</sup> in the raw material composition resulted in improved thermal stability of the AAM. In case where SiO2/Al2O<sup>3</sup> ratio is <2, the formation of high-temperature resistant minerals, such as carnegeite and nepheline, was observed. The obtained AAM could resist up to 8 thermal shock cycles and it could be easily adapted to the industrial production and application such as thermal insulation layer in laboratory furnaces.

Keywords: alkali activated material, porous materials, heat resistance, application in high temperature, aluminum scrap recycling waste

## HIGHLIGHTS


### Edited by:

*Maria Juenger, University of Texas at Austin, United States*

#### Reviewed by:

*Sriramya Duddukuri Nair, Cornell University, United States Ru Mu, Hebei University of Technology, China*

> \*Correspondence: *Laura Vitola laura.vitola\_1@rtu.lv*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *18 January 2019* Accepted: *12 April 2019* Published: *22 May 2019*

#### Citation:

*Bajare D, Vitola L, Dembovska L and Bumanis G (2019) Waste Stream Porous Alkali Activated Materials for High Temperature Application. Front. Mater. 6:92. doi: 10.3389/fmats.2019.00092*

## INTRODUCTION

Low calcium alkali activated materials (AAM) or "Geopolymers" in some cases are the terms used to describe aluminosilicatebased inorganic polymers produced by mixing pozzolanic compounds or aluminosilicate source precursors and highly alkaline solutions (Davidovits, 1991). Production of AAM provide an opportunity to convert a variety of waste stream materials and by-products into useful, high value materials (Singh et al., 2015). Many factors can contribute to the alkali activation reaction mechanism and the properties of the final products, including (but not limited to) the aluminosilicate source precursors (Bumanis et al., 2017), the type of activator solution (Criado et al., 2007), the calcium content in the raw material mixture composition (Canfield et al., 2014), and the mixing and curing conditions (Moukannaa et al., 2018) as well as pore forming process (Bajare et al., 2014; Strozi et al., 2014; Seabra et al., 2016). AAM have attracted an increasing interest as fireproof materials associated with circular economy, namely, their production have low energy requirements and the final product is characterized by good mechanical performance, thermal behavior, and durability (Rashad et al., 2016). The existing publications on fire resistance and performance of AAM under high temperature conditions are promising.

The stability at high temperatures of binders made by the alkaline activation of aluminosilicate precursors is clearly distinct from the behavior observed in Portland cement at high temperatures (Martin et al., 2015). The shrinkage of AAM up to 600◦C temperature occurs due to evaporation of the adsorbed and weakly bound water. Several reports indicate that the strength of AAM increases by the treatment at temperature between 600 and 800◦C as densification occurs on the contrary with Portland cement-based materials (Fernández-Jiménez et al., 2010). In many cases melting process of AMM starts up to the temperature of 1,100◦C (Kim and Kim, 2017). The resistance to high temperatures of AAM strongly depends from the composition of raw materials, i.e., Dupuy et al. has reported that all argillite-based AAM samples show low resistance to thermal shock (800◦C, 10 min) and strength decreased up to 10 times (Dupuy et al., 2018), while other report by K. Hemra and P. Aungkavattana states that thermal shock resistance of metakaolin-based AAM at temperature of 800◦C could reach up to 15 cycles (Hemra and Aungkavattana, 2016). In comparison with traditional refractory brick performance these results seems promising as their thermal shock resistance varies from 10 to 30 cycles (Philip, 2014).

One of the basic approaches to improve thermal properties of the AAM is to use fillers based on refractory materials. Bernal et al. reported that thermal behavior of metakaolinbased AAM with ground refractory brick filler (10–15 vol.%) could reduce the volumetric contraction and increase the residual compressive strength up to three times after exposure to high temperature in comparison with the AAM without fillers (Bernal et al., 2012).

Type and concentration of alkali activation solution play a dual role which calls for a certain compromise in the AAM design. On one hand, high alkali content in activation solution (8 to 12 M NaOH solutions) is necessary to accelerate the activation reactions of precursors and to obtain AAM with high initial strength, but at the same time increased concentration of alkalis could act as fluxes at high temperatures. Therefore, the content of alkalis in the alkali activation solution must be limited to generate thermally stable materials (Martin et al., 2015). It has been reported that the alkali concentration in the activation solution could be reduced event to 2 M (Huiskes et al., 2016). The results of previous research show that increasing the SiO2/Na2O molar ratio of the activation solution the setting time of AAM mixture could also increase (Tchakoute et al., 2013). It has been reported that in cases, when ratio of SiO2/Na2O of activation solution is beyond optimum (>3.0), physical properties of the AAM can draw down quickly—water absorption increases while the bulk density does not change (Lemougna et al., 2011).

The aim of the present study is to assess the thermal stability of chamotte-like precursor based AAM with refractory brick byproduct additive in order to achieve high mechanical properties after heat treatment and thermal shock resistance of the obtained material. The study includes development and demonstration of AAM production and application methods.

## MATERIALS AND METHODS

## Raw Materials

Commercially produced chamotte-like precursor (CH), where 36.9% of particles were with size lower than 45µm was used in the research (Keramserviss Ltd). According to the XRD analysis CH contains the following minerals: cristobalite (SiO2), quartz (SiO2), and mullite (Al6Si2O13). Also minor amorphous phase as background hump was identified (**Figure 1**).

Firebrick sawing residues (K26) (Morgan Thermal Ceramics Ltd.) is classified as a by-product from the sawing process of insulation material typically used for furnace production. The K26 is aluminosilicate material and its chemical composition is given in **Table 1**. K26 was ground for 30 min by using planetary ball mill Retsch PM 400. 96.3% of the particles were with the size lower than 45µm. According to the technical data sheet, working temperatures of the K26 range from 1,260 to 1,790◦C. The K26 is made from high-purity refractory clay with high amount of Al2O<sup>3</sup> (up to 60%). According to the XRD analysis, mullite (Al6Si2O13) and corundum (Al2O3) as well as small background hump indicating amorphous phase were identified (**Figure 1**). K26 was considered as additional aluminum silicate source and as filling material which could improve thermal stability of the AAM.

Aluminum scrap recycling waste (ASRW) was received from the aluminum scrap recycling facility (Dilers Ltd). ASRW is the final waste, which is obtained from black dross in aluminum recovery process. ASRW was ground for 30 min by using planetary ball mill Retsch PM 400. Chemical composition of ASRW is given in **Table 1**; main oxides are Al2O<sup>3</sup> (63.2 wt.%), SiO<sup>2</sup> (7.9 wt.%), CaO (2.6 wt.%), and Na2O+ K2O (7.7 wt.%). According to the XRD analysis, the ASRW contains quartz (SiO2), spinel (MgAl2O4), iron (III) oxide (Fe2O3), calcium iron aluminum oxide (CaAlFe4O10), magnesium aluminum silicate (Mg2Al2Si3O12), aluminum iron oxide (FeAlO3), and gibbsite (Al(OH)3) (**Figure 1**). Ground ASRW with fraction <1 mm was

FIGURE 1 | X-ray analysis of solid raw materials, (A) CH, (B) K26, (C) ASRW; Q / q, quartz (85-0798 / 83-2187); Cr, cristobalite (03-0267); M, mullite (01-0613); C, corundum (71-1124); s, spinel (75-1799); f, iron oxide (32-0469); , calcium aluminum iron oxide (21-0830); , magnesium aluminum silicate (30-0788); •, aluminum iron oxide (18-0633); g, gibbsite (70-2038).

TABLE 1 | Chemical composition of solid raw materials, (weight %).


TABLE 2 | Compositions of studied samples and amount of main oxides in pastes.


used as the pore forming agent due to gas emissions, when it reacts with an activation solution (Bajare et al., 2012).

6 M NaOH solution, which was prepared by using sodium hydroxide flakes with 99% purity (Tianye Chemicals Ltd., China), was used as alkali activating solution.

## Mixture Design and Sample Preparation

The basic AAM mixture compositions consisted of CH and ASRW with the mass ratio of 10:1. The mass ratio of the alkali activation solution to total solid raw materials ratio was 0.33. CH was partially substituted with sawing residues K26 up to 50 wt.% (**Table 2**). The theoretically calculated oxide ratios by mass are given in **Table 2**. The SiO<sup>2</sup> / Al2O<sup>3</sup> ratio for studied AAM is in the range from 1.3 to 3.1, SiO<sup>2</sup> / Na2O ratio—from 8.6 to 11.1 and Al2O<sup>3</sup> / Na2O−5.0 to 9.2. According to the report of Peigang He, the increased ratio of SiO2/Al2O<sup>3</sup> can increase mechanical properties of the AAM due to the increased amount of Si-O-Si bonds, which could appear in the polymerization process and presence of residual silica, which can operate as reinforcement (He et al., 2016). The decrease of Na2O/Al2O ratio could cause less dissolution of precursors due to lower content of NaOH resulting in weaker active alkali activation reaction (Siyal et al., 2016).

The solid raw materials were mixed together and activation solution was added. After mixing with electrical twin shaft mixer for 2 min, the paste was immediately poured into prismatic molds (40 × 40 × 160 mm) sealed with plastic film and after formation of porous structure molds were covered and samples were cured at 80◦C for 24 h. For thermal conductivity test and industrial application AAM plates of 45 × 45 × 5 cm were prepared. To characterize the effect of heat treatment on AAM properties prepared samples were treated at 800 and 1,000◦C temperatures with the heating rate of 10◦C/min and kept for 3 h at the max temperature. Three series of samples were prepared: Ref, 0.1FB, 0.3FB, and 0.5FB without heat treatment (i), Ref-800, 0.1FB-800, 0.3FB-800, and 0.5FB-800 with heat treatment at temperature 800◦C (ii) and Ref-1000, 0.1FB-1000, 0.3FB-1000, and 0.5FB-1000 with heat treatment at temperature 1,000◦C (iii). The obtained AAM samples were kept at ambient temperature (20 ± 2 ◦C) for further testing.

## Techniques

Chemical composition of the raw materials was determined according to LVS EN 196-2 with sensibility ±0.5 wt.%. The elemental analysis of ASRW was carried out with inductively coupled plasma optical emission spectrometry (ICP-OES) for Al and Si, atomic absorption spectrometry (AAS) for Na, Mg, K, Ca, Fe, Cu, Zn, Pb. Potentiometer titration analysis (PTA) was used to determine Cl and combustion process—to determine S.

Thermogravimetric-differential thermal analysis (DTA/TG) was carried out by using Netzsch STA 409 PC from 0 to 1,200◦C with the heating rate of 10◦C/min. High temperature microscopy (HTOM) EM201, HT163 was used to determinate heat resistance and shrinkage of the AAM in temperatures up to 1,400◦C. Samples were tested by heating rate 80◦C/min up to 500◦C temperature and then switched to 15◦C/min, while reaching 1,400◦C temperature.

The material density of AAM was measured by the geometric method using 40 × 40 × 40 mm specimens. The Archimedes principle (using water as immersion fluid) was employed to evaluate the specimens' water absorption, and combined with geometric measurements to determine the open porosity. The true density was determined for powdered AAM by using Le Chatelier flask and then the total porosity was calculated. Xray computed tomography (XCT, Xradia µCT-400, XRadia, Concord, USA with Avizo Fire 3D software) was employed to obtain 3D images of the AAM. Pore structure and porosity were assessed by using the XCT data and comparing the results to the traditional methods.

The AAM macro and microstructure was observed by using optical microscope and scanning electron microscope ("TESCAN Mira / LMU Field-Emission-Gun"). The thermal conductivity was measured with heat flow meter instrument LaserComp FOX 660.

Flexural and compressive strength of the prepared AAM samples 9 aged 28 days) were tested according to the LVS EN 196-1.

To determine the mineralogical composition of solid raw materials and AAM before and after heat treatment the Xray diffraction (XRD) patterns of the powdered samples were recorded on a "RIGAKU ULTIMA+" diffractometer using CuKα radiation, the test was run in a 2-Theta range of 5–70◦ step-scan and 10 s/step.

The thermal cycling of the prismatic samples 40 × 40 × 160 mm was conducted in accordance with testing of refractory concrete (GOST 20910-90 Refractory concretes, 1991). The samples were kept for 60 min in a furnace heated to 800◦C, then the samples were cooled for 20 min by a stream of air from a fan (700 W, 1,000 rpm). After each thermal cycle the speed of an ultrasonic pulse (SUSP) in the samples was determined and its decrease with respect to the SUSP values after heat treatment at 800◦C was determined (Antonovich et al., 2011).

## RESULTS

## Thermal Analysis DTA/TG

The changes of AAM during heat treatment up to 1,200◦C could be described in four distinct phases (**Figure 2**). The first phase consists of evaporation of physically adsorbed water from the material matrix. Exothermic effect is appearing in DTA curve before heating temperature rises to 100–150◦C (Djobo et al., 2016). The second phase begins at 150◦C and proceeds up to 600◦C involving evaporation of the chemically bound water. It comes from dissolution of aluminosilicate raw materials, when aluminate and silicate species are formed. In this phase significant processes contributing to the knowledge base about the extent of alkali activation reaction can be observed (Firdous et al., 2018). Namely, the TGA result showed similar behavior of AAM, when between 150 and 600◦C the mass loss was 2.91 and 2.67% respectively. The mass loss ∼5% in the temperature between 500 and 700◦C as well as start of another mass loss region at 900◦C along with minor features on the differential scanning calorimeter represented various phases of crystallization (Canfield et al., 2014). The densification of AAM occurs at 820–920◦C. Onset temperatures decrease slightly and sharply appear due to softening and viscous sintering (Barbosa and MacKenzie, 2003; Villaquirán-Caicedo et al., 2017). Another minor weight loss of 0.4% at the temperature around 800◦C is associated with the decomposition of calcite which in minor parts was detected in ASRW. Finally, rapid densification and weight loss was observed above 1,000◦C.

## High Temperature Microscopy (HTOM)

HTOM results can be divided in three relative intervals which correlate well with DTA/TG results: (a) specific area changes at the temperature up to 800◦C; (b) from 800 to 1,000◦C; and (c) from 1,000 to 1,200◦C (**Figure 3**; **Table 3**). Shrinkage of the AAM with the intensity from 3.2 to 4.9% was observed at the temperature up to 1,000◦C.

## Physical and Mechanical Properties

The physical properties of the porous AAM are given in **Figure 4**. The density of porous AAM was in the range from 670 to 755 kg/m<sup>3</sup> . The total porosity of the obtained materials was between 71 and 75 vol.%. The open porosity was from 27 to 30 vol.%. The pore structure of the obtained materials remain open which corresponds to the water absorption. The water absorption of the AAM was between 45 and 48 wt.%.

The thermal conductivity of prepared porous AAM plates (45 × 45 × 5 cm) was measured. The measured thermal conductivity was from 0.14 to 0.15 W/m·K and it was not affected by the mixture composition of AAM but correlated well with total porosity of the material.

The porosity measured by the micro-XCT indicated the total pore value ranging from 41.2 to 52.7 vol.% with the pore structure of the AAM, pore distribution as well as the measurement technique being the factors contributing to these differences. Three-dimensional internal pore structure of the specimens was made and pore size distribution was reconstructed using 3Dimage analysis software (Avizo Fire) (**Table 4**). According to the research results the size of most of the pores in the material (90.2 vol.%) was between 1,000 and 5,000µm. It can be clearly seen in the pore sieving images obtained by using the Avizo Fire 3D software. The pores within the range from 7E+9 to 9E+9 µm<sup>3</sup> have a blue color, whereas the pores having a size greater than 9E+9 µm<sup>3</sup> have a green color (**Figure 5**).

The compressive and flexural strength results of the AAM are presented in **Figure 6**. The compressive strength was 1.1 MPa for Ref and increased to 2.0 MPa for 0.5FB. After sample


TABLE 3 | Relative area changes of obtained porous AAM.

TABLE 4 | Volumetric pore size distribution (vol. %) of 0.5FB determined by Micro-XCT.


treatment at 800◦C, compressive strength was 0.7 for Ref and 0.9 MPa for 0.1FB. Slight strength increase was for 0.3FB, while for the 0.5FB strength increased to 2.5 MPa. The heat treatment at 1,000◦C slightly increased strength comparing to samples treated at 800◦C for the mixture compositions Ref, 0.1FB and 0.3FB. The composition 0.5FB had compressive strength 2.7 MPa.

The flexural strength for all samples except for 0.5FB before heat treatment was in a range from 0.5 to 0.6 MPa, while the strength of 0.5FB samples was 0.8 MPa. The flexural strength decreased to 0.3–0.5MPa for samples with Ref, 0.1FB, 0.3FB but the increase of flexural strength was observed for 0.5FB—reaching 1.0 MPa. Similar tendency was observed for the heat treated samples at 1,000◦C. Structural integrity was preserved and insignificant or no strength loss was observed compared with specimens treated at 800◦C temperature.

## X-Ray Diffraction

The XRD was determined for blends of solid raw material (**Figure 7** with index raw), prepared porous AAM before and after heat treatment at 800 and 1,000◦C for mixture compositions Ref and 0.5FB. After alkali activation slight phase change was observed for Ref but almost all minerals detected in mix of raw materials were still present in the AAM. The composition Ref had crystalline phases, which have remained from precursors (CH) like mullite (Al4.56Si1.44O9.72), quartz (SiO2), and cristobalite (SiO2). After alkali activation low intensity peaks of analcime appeared, but after heat treatment at 800 and 1,000◦C crystalline phase of analcime disappeared and new crystalline phases have not been detected. Meanwhile composition 0.5FB had crystalline phases, which had remained from precursors like mullite (Al4.56Si1.44O9.72), quartz (SiO2), corundum (Al2O3), and cristobalite (SiO2). Hydroxy sodalite (1.08Na2O·Al2O3·1.68SiO2·1.8H2O; sodalite group zeolites) and N-A-S-H gel which are formed during the activation process were also detected. The hydroxy sodalite was not detected after heat treatment of specimens at 800◦C temperature, but new crystalline phase carnegeite (NaAlSiO4) was present. After heat treatment at 1,000◦C phases such as quartz, mullite, corundum, and crystabolite remained, but new crystalline phase—nepheline (Na6.65Al6.24Si9.76O32) was detected.

## Cycling Performance

The cyclic performance of the prepared AAM was evaluated (**Figure 8**). Pre-treatment of separate batch of samples was done to reduce first thermal shock induced by the formation

of new minerals inside the AAM previously detected by the XRD. The ultrasonic pulse velocity was measured to evaluate the performance. The cycling results are given in **Table 5**. The first thermal cycle at the temperature of 800◦C reduced the initial ultrasonic pulse velocity by 30 to 50%, while the critical value of 300 m/s was not reached. The 2nd high temperature cycle was crucial for the samples Ref, Ref\_800 and 0.1FB with the ultrasonic pulse velocity decrease below 300 m/s. With the increasing proportion of K26 in the mixture composition of AAM the thermal cycle performance improved and up to 7 cycles were achieved for the specimens 0.5FB\_800.

## DISCUSSION

The mass loss intensity during TGA analysis can be used to measure the extent of alkali activation (**Figure 2**). Higher overall mass loss relates to higher extent of alkali activation as more water is attracted to hydrated NASH gel (Vogt et al., 2016). Increased weight loss was observed for 0.5FB as higher amount of K26 filler was incorporated in the mixture composition of AAM, showing the positive role of K26 in alkali activation process.

FIGURE 7 | XRD analysis of obtained porous AAM and solid raw material mixtures, (A) porous AAM based on chamotte-like clay (Ref), (B) porous AAM based on chamotte-like clay and firebricks sawing residues (0.5FB).

The DTA/TG results showed that at the interval between 800 and 1,000◦C the structural changes could take place; therefore, further investigation was done at selected temperatures. Since the DTA/TG results showed similar performance between Ref up to 0.5FB with different K26 concentration in mixture composition, for the ease of result interpretation compositions Ref and 0.5FB were compered except in cases, where noticeable differences were detected.

Thermal expansion can be limited by highly porous structure of the AAM (**Figure 3**). The deformations in specimen can take place inside the pores and true expansion value could be even greater. Some expansion was attributed to the transformation of α-quartz to the β-quartz which follows the volume increase of specimens at the temperature 573◦C. Due to transformation of quartz and expansion of specimens the microstructure of AAM could be damaged, which could have negative influence on the

FIGURE 9 | Porous structure of 0.5FB, (A) optical microscope image with magnification x40, (B) SEM image with magnification x200.

mechanical properties. According to the literature the second interval from 800 to 1,000◦C is associated with melting of alkalis. At temperatures over 800◦C solid raw materials are attacked by alkalis, which causes partial sintering in the contact zone between the geopolymer matrix and the aggregate, which in this case is unreacted crystalline phases remained from raw materials. This effect plays a substantial role in the improvement of residual mechanical properties after exposure to 1,000◦C and helps AAM to resist higher temperatures without damage (Mandal et al., 2011; Rovnaník and Šafránková, 2016).

The physical and mechanical properties were slightly affected by the selected mixture compositions. Porosity within the material structure, as well as the material density, depends on the gas supply in the fresh paste. The chemical reaction between the ASRW and the alkali solution as well as generated heat flux leads to gas formation in the fresh paste. Low viscosity of the fresh paste and the pressure of evolved gasses in the process of pore formation are the factors, which cause breaking down of the walls and coalescence of pores, thus small bubbles diffuse into the bubbles, which are larger (**Figures 9**, **10**). The above mentioned observations show the differences between the pore formation using ASRW as pore formation agent and other possible solutions

FIGURE 10 | The structure of obtained AAM detected by Micro-XCT.

such as H2O<sup>2</sup> (Ducman and Korat, 2016). Structure formation, gas release, and the evaporation of free water during initial setting of AAM are the factors contributing to high open porosity (up to 33%). In the SEM microphotographs it is possible to observe micropores not exceeding few micrometers (**Figure 9B**). As these small micropores were not detected by the Avizo Fire 3D software due to insufficient resolution of X-ray computed tomography, it is possible that some of the pores are not included in the calculations of porosity by the micro-XCT (Dembovska et al., 2017). The results of the pore analysis obtained using Avizo Fire 3D software include only the pores being larger than 48µm (identified as macropores), whereas the difference between the total porosity detected by using the Le Chatelier flask and the porosity detected by the Avizo Fire 3D software show both micro and capillary pores.

The amount and structure of pores had influence on compressive strength of the AAM giving characteristically low compressive strength associated with highly porous materials. The increase of K26 amount in the composition of AAM increased the initial compressive strength. The heat treatment of samples at the temperature of 800◦C reduced compressive strength for Ref and 0.1FB to 0.7 and 0.9 MPa, respectively, indicating strength reduction by 36 and 32% while for the 0.5FB strength increase was 24% reaching 2.5 MPa. The heat treatment at 1,000◦C led to slight strength gain comparing to the samples treated at 800◦C for the mixture compositions Ref, 0.1FB and 0.3FB, while it was still lower comparing to the samples without heat treatment. The composition 0.5FB had further strength increase up to 2.7 MPa (8% increase


TABLE 5 | Ultrasonic pulse velocity after high temperature cycling.

*Acceptable results has been highlighted green, i.e. values over 300.*

comparing to the strength at 800◦C) indicating the improved structure performance with the incorporation of higher proportion of K26.

The XRD results indicate the N-A-S-H gel formation during alkali activation process and it was detected by the shift of halo at 2-Theta = 10−30◦ to higher degrees (Zhang et al., 2012) in both cases (for Ref and 0.5FB before heat treatment respectively). The phase transformation was identified during heat treatment of AAM. It could be explained by transformation of initially detected hydroxide sodalite into zeolite X, which melts at 760◦C temperature and becomes amorphous but later it recrystallizes again at 800◦C and crystallizes as carnegeite (Jacobs, 1984). After further temperature increase the nepheline transforms from the carnegeite at the temperature around 900– 1,000◦C. Considerable change of the N-A-S-H gel after heat treatment at different temperatures was not detected by Xray diffraction.

Also phase changes was detected by XRD and the first controlled pre-treatment at the temperature of 800◦C was applied for porous AAM, it did not improve the overall cyclic performance of AAM and only initiated the early deterioration of AAM structure by showing possible relationship between pre-treatment and first cyclic load for untreated samples. Each additional cycle continued to damage structure of the AAM.

## HIGH TEMPERATURE CYCLIC PERFORMANCE AND INDUSTRIAL APPLICATION

## Technological Perspective of High Temperature AAM Production: Production Process

The schematic diagram of the AAM production process is given in **Figure 11**. The estimation of AAM mixture composition regarding raw material chemical composition should be calculated to cover the optimum range of main oxide ratios. The production process consists of dosage of raw materials and pre-mixing to obtain homogenous precursor for alkali activation. It should be noted that ASRW and K26 components must be pre-treated by milling to obtain fine, powder-like material. After adding alkali activation solution and mixing, paste should be molded instantly to allow blowing agent to react and form porous structure of the AAM. If prolonged mixing or molding time is necessary, the cooling process of raw materials should be considered (reaching −20◦C) to delay the initial exothermic reaction provided by the ASRW, which determines rapid setting of the AAM. The further process control related to the curing regime should be maintained at 80◦C for 24 h. The ventilation and filtering of gasses emitted by the ASRW during blowing of AAM should be ensured to provide safe production process and to avoid environmental pollution. After the AAM are demoulded, high temperature pre-treatment at 800◦C could be applied to the allow phase change transformations take place to ensure further dimensional stability before cutting the material and creating final shape of the detail.

## Technological Perspective of High Temperature AAM Production: Industrial Approbation

Following the previously described technological process, AAM plates for industrial application were prepared. 45 × 45 × 5 cm porous AAM plates were made in the laboratory conditions and provided to the producer of laboratory furnaces. During further processing of the prepared AAM plates they proved to be workable with traditional tools (**Figure 12**). The high temperature oven prototype was created by combining prepared porous AAM insolation plates with traditional refractory insulation materials. The prepared AAM was placed as outer insulation layer of the oven. The possible limitations of the AAM described is the relatively low strength of the material which is associated with high porosity; therefor the application of the material is suggested to be used as the interlayer material in high temperature furnaces which main function would be heat insulation. Other limitation which is related with production process is associated with the availability of the raw materials and its nature, as most of them are industrial wastes and by-products so the production of the material must be carefully controlled.

## CONCLUSIONS

The investigated porous alkali activated material (AAM), which is based on chamotte-like precursor, firebrick sawing residues and aluminum scrap recycling waste, proved to be suitable for industrial production and applications.

The tested alkali activated material had high porosity (open porosity 25–30 vol.%, total porosity 71–75 vol.%), reasonable strength (before heat treatment from 1.1 to 2.1 MPa and after treatment at 1,000◦C 0.8–2.6 MPa, respectively) and effective working temperature up to 800 (shrinkage from −0.4 to 1.0%) and 1,000◦C (shrinkage from 3.2 to 4.9%).

High thermal stability of porous AAM reaching up to 8 thermal shock cycles was obtained. It is suggested that the ultrasonic pulse velocity lower limit of 300 m/s could be acceptable for non-destructive testing method for porous AAM, as samples withstand thermal cycling load and remain solid appearance and still reasonable strength.

The macropore (>48µm) structure of AAM could be effectively determined by the X-ray computed tomography with Avizo Fire 3D software giving possibility to test structure with non-destructible method i.e., for quality control during the production. Though to characterize the total porosity test methods based on the Archimedes principle should be used.

The obtained material has a simple, reproducible and low-cost manufacture ensuring maximal valorization of waste into a heatresistant material with high added value. The firebrick sawing residue incorporation in the composition of AAM improved the performance of heat resistance as the increase of Al2O<sup>3</sup> content was achieved. This innovative waste-based material was treated at temperature 1,000◦C and proved to be suitable for application at the temperature up to 1,000◦C.

It was proved that the developed material is effective for industrial production and could be a promising for using as outer insulation layer in high temperature furnaces.

## AUTHOR CONTRIBUTIONS

All co-authors are working together on porous AAM for different kind of applications since 2013. DB is the main author of this manuscript and supervisor of all other co-authors. LV is a Ph.D. student and her responsibility at this research was to analyze XRD

## REFERENCES


results, as well as other obtained data. LD is the last year Ph.D. student and this research is included in her thesis, so most of the tests have been done by LD. GB is leading researcher, as his Ph.D. thesis was about porous AAM and he has the experience of the work with aluminum scrap recycling waste usage in AAM production he was helping with data interpretation.


Zhang, Z., Wang, H., Provis, J. L., Bullen, F., Reid, A., and Zhu, Y. (2012). Quantitative kinetic and structural analysis of geopolymers. Part 1. The activation of metakaolin with sodium hydroxide. Thermochim. Acta 539, 23–33. doi: 10.1016/j.tca.2012.03.021

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Bajare, Vitola, Dembovska and Bumanis. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

## Inorganic Polymers From CaO-FeOx-SiO<sup>2</sup> Slag: The Start of Oxidation of Fe and the Formation of a Mixed Valence Binder

Arne Peys <sup>1</sup> \*, Alexios P. Douvalis <sup>2</sup> , Vincent Hallet <sup>1</sup> , Hubert Rahier <sup>3</sup> , Bart Blanpain<sup>1</sup> and Yiannis Pontikes <sup>1</sup>

*<sup>1</sup> Department of Materials Engineering, KU Leuven, Leuven, Belgium, <sup>2</sup> Department of Physics, University of Ioannina, Ioannina, Greece, <sup>3</sup> Department of Materials and Chemistry, Vrije Universiteit Brussel, Brussels, Belgium*

#### Edited by:

*Andrew C. Heath, University of Bath, United Kingdom*

#### Reviewed by:

*Hongyan Ma, Missouri University of Science and Technology, United States Zuhua Zhang, University of Southern Queensland, Australia Rupert Jacob Myers, University of Edinburgh, United Kingdom*

#### \*Correspondence:

*Arne Peys arnepeys@hotmail.com*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *22 February 2019* Accepted: *14 August 2019* Published: *28 August 2019*

#### Citation:

*Peys A, Douvalis AP, Hallet V, Rahier H, Blanpain B and Pontikes Y (2019) Inorganic Polymers From CaO-FeOx-SiO2 Slag: The Start of Oxidation of Fe and the Formation of a Mixed Valence Binder. Front. Mater. 6:212. doi: 10.3389/fmats.2019.00212* Belonging to the family of alternative cementitious materials, inorganic polymers are rising in importance because of the drive to decrease CO<sup>2</sup> emissions of concrete production. A synthetic Fe-rich slag resembling industrial copper or lead slags, was mixed with a sodium silicate activating solution. <sup>57</sup>Fe Mössbauer spectra analyses indicate that the oxidation reactions are taking place simultaneously with the polymerization reactions. The slag contains Fe2<sup>+</sup> states and a small amount of Fe3+. During polymerization a new octahedral Fe2<sup>+</sup> state is formed, while oxidation is manifested through the appearance of an additional Fe3<sup>+</sup> state. The reactions continue after setting, lowering the relative contributions of the slag in the Mössbauer and FTIR spectra of the samples. The Na+/Fe3<sup>+</sup> molar ratio in the mixture that makes up the binder after 28 days is ∼1, suggesting the participation of tetrahedral Fe3<sup>+</sup> in the silicate framework, charge balanced by Na+.

Keywords: alkali activated materials, geopolymers, <sup>57</sup>Fe Mössbauer spectroscopy, non-ferrous slag, ex-situ kinetics, FTIR spectroscopy

## INTRODUCTION

In search for alternative cementitious materials, geopolymers (Davidovits, 2008), and alkali activated materials (AAM) (Provis and van Deventer, 2014) have received wide attention. More recently, a push for circularization of the economy encouraged metallurgical companies to find alternatives for their residues. In light of this, research in inorganic polymers (IPs) is expanding to ferrosilicate precursors, for instance slags from the non-ferrous metal industries. It was shown that these residues can be increased in reactivity by rapid cooling (Pontikes et al., 2013) or hot stage slag engineering (Van De Sande et al., 2017), while the IPs thereof obtain mechanical properties comparable to geopolymers, i.e., compressive strengths reaching up to 80 MPa (Kriskova et al., 2015). They have shown to be environmentally important, reducing environmental impacts down to 17% of that of ordinary Portland cement (OPC) (Peys et al., 2018a).

Despite these interesting properties and because of the novelty of the understanding that ferrosilicate precursors can be used in alkali-activation, fundamental knowledge on the chemistry of the system and the structural rearrangements during formation is scarce. Using isothermal calorimetry it was shown that, chemically, iron-rich inorganic polymers seem to follow a two stage reaction pattern: dissolution and polymerization (Kriskova et al., 2015; Onisei et al., 2015). As such,

**101**

they are comparable to their aluminosilicate counterparts, where calorimetric techniques are often used to distinguish these two reaction stages (Rahier et al., 2007; Zhang et al., 2013) or to compare the reactivity of different formulations (Kumar et al., 2007; Lemougna et al., 2014; Peys et al., 2016). More thorough analysis of calorimetry results also enables the calculation of the reaction enthalpy and subsequent determination of the stoichiometry of the chemical system (Rahier et al., 1997) or the determination of the rate limiting step in the formation (Rahier et al., 2003). In terms of the evolution of the molecular structure, in-situ Fourier transform infrared (FTIR) spectroscopy is probably the most used technique, mainly because of its accessibility. Although success is seen for aluminosilicates (Rees et al., 2007; Hajimohammadi et al., 2010), in-situ FTIR by itself did not show the evolution of the structure (Onisei et al., 2015) sufficiently clearly and the exact reaction mechanism therefore remains unclear. However, more recent developments (Peys et al., 2017; Onisei et al., 2018) have increased the understanding of the infrared spectrum of ferrosilicate precursors and their alkali-activated products. The main interesting feature in the spectra is the Si-O stretching band. The position of this band in the precursors varies between 950 and 850 cm−<sup>1</sup> , depending on the amount of network modifiers per silicon atom in the silicate phases of the slag (Peys et al., 2017). This band shifts to ∼950 cm−<sup>1</sup> after alkali-activation (Kriskova et al., 2015; Peys et al., 2017; Onisei et al., 2018). This was attributed to a higher degree of polymerization observed in recent work (Peys et al., 2018b), because of the (at least partial) incorporation of the Fe3<sup>+</sup> species in the network.

The aforementioned oxidation was experimentally observed using <sup>57</sup>Fe Mössbauer spectroscopy (Lemougna et al., 2013; Onisei et al., 2018). In these works the success in providing information for the oxidation state of iron (Fe3+) in the IP binder, did not expand to the level of the exact description of the coordination of iron (Lemougna et al., 2013; Onisei et al., 2018). Both referred papers suggest an average 4- or 5-fold coordination. Similarly, X-ray absorption spectroscopy (XAS) measurements (Simon et al., 2018) confirm the oxidation state of the binder phase and suggest an average coordination number of 5. An important characteristic referred to the formation of the binder is that the evolution over time of the oxidation reaction is unknown. Mössbauer spectroscopy measurements can be performed on solutions after freezing in liquid nitrogen temperatures (Silver et al., 1979; Douglas et al., 1984). As this practically decelerates the reaction kinetics toward negligible speeds, it enables their ex-situ study (Krebs et al., 2005; Kamnev et al., 2014). In-situ kinetic studies have also been performed using synchrotron radiation techniques, e.g., X-ray/neutron total scattering on the alkali-activation of metakaolin (Bell et al., 2008; White et al., 2011, 2013a), blast furnace slag (White et al., 2013b; Garg and White, 2017) and on conventional cements (Grangeon et al., 2017; Bae et al., 2018; Cuesta et al., 2019). With these techniques, the short and intermediate range ordering can be investigated in great detail. The results of these studies are often interpreted with the help of theoretical modeling techniques (White et al., 2010; Yang and White, 2016; Geng et al., 2017; Zhou et al., 2017), which increases the understanding of the molecular structure or formation mechanism to a great extent.

In the present work, the reaction mechanism of Fe-silicate IPs is clarified with respect to the behavior of Fe in the first 28 days after mixing. A combination of <sup>57</sup>Fe Mössbauer spectroscopy and attenuated total reflectance Fourier transform infrared (ATR-FTIR) spectroscopy is performed ex-situ to provide details on the evolution of the local environment of Fe and the connectivity of the silicate network throughout the reaction. These observations are combined to provide new information on our binder system, such as the timing of oxidation reactions. Additionally, they provide detailed information on the local atomic environment of Fe in potential intermediate stages of the material, formed between the initial slag and the final IPs studied in previous works. As such, a step is taken in unraveling the mechanism of formation of Fe-silicate IPs.

## EXPERIMENTAL METHODS AND MATERIALS

The precursor slag with chemical composition given in **Table 1** was synthesized by melting iron ore, limestone, and sand in a pilot scale furnace. After a complete melt was obtained, the slag was tapped and granulated by spraying the slag with pressurized water jets into droplets to maximize the amount of glassy phase in the precursor. The precursor hereby replicates a composition often used in non-ferrous metallurgy, but including only the major elements FeO, SiO2, CaO, and Al2O<sup>3</sup> (a minor content of MgO was detected due to impurities in the raw materials). This methodology to produce synthetic glasses at pilot scale is described in more detail in Machiels et al. (2017). Milling was performed to reach a specific surface of 4,000 ± 200 cm²/g measured by the Blaine method. The chemical composition was measured standard-less with wavelength dispersive X-ray fluorescence (XRF) on a Philips PW 2,400. The glassy phase in the precursor was quantified by X-ray diffraction (XRD) after mixing with 10 wt.% of internal zincite standard in a McCrone mill. A Bruker D2-Phaser with Cu K<sup>α</sup> radiation was used for these measurements under 30 kV−10 mA, a slit size of 0.6 mm and an anti-scatter slit of 1 mm. With these settings, no crystalline peaks could be observed apart from those of zincite; the amount of glassy phase is estimated to be > 97 wt.%. The diffractogram is added in the supplementary information (**Figure S1**) for interested readers. The IP mixture consisted of the above described precursor and a sodium silicate solution with molar ratios SiO2/Na2O of 1.6 and H2O/Na2O of 20. The solution/slag mass ratio was 0.45. Pastes using these ratios were mixed by hand and used in all experimental techniques below.

TABLE 1 | Chemical composition of precursor slag in wt.%, obtained from XRF (estimated relative error 10%).


Isothermal calorimetry was performed using a TAM Air at 20◦C. A sample with 4 g of slag was mixed for 90 s and transferred into the microcalorimeter. Data are presented after 53.5 min of stabilization in the machine to avoid the influence of the temperature difference between the calorimeter and sample (+26.5 min preparation). The first dissolution and wetting heat is not monitored, i.e., no admix ampule was used, as this information was not needed for the present study. ATR-FTIR spectroscopy was performed on a Bruker Alpha-P, which uses a diamond ATR crystal. A resolution of 4 cm−<sup>1</sup> was selected and a measurement considered 32 scans. Before setting, for measurements after 6 h or less, a drop was measured without applying pressure on the paste in the ATR-cell. Because the powder is not pressed against the crystal, a relatively higher intensity will come from the solution, where the changes are anticipated to be the most interesting. For later reaction stages (>6 h), the reported spectrum is the average of 2 measurements on powdered samples (these were pressed against the crystal). The powdering of the samples was performed right before the FTIR (and Mössbauer) measurements, before which the pastes were stored in closed bottles at 21 ± 1 ◦C. The FTIR spectra shown in the plots in the results section are rescaled in the y-axis to show the same maximum absorbance of the Si-O stretching band. <sup>57</sup>Fe Mössbauer spectroscopy measurements were performed to investigate the evolution of the state and environment of the Fe atoms during reaction. Gamma rays were emitted from a source using the decay of <sup>57</sup>Co in Rh matrix kept at room temperature. The samples were measured in transmission geometry on a constant acceleration spectrometer at 77 K. The reported isomer shift (IS) values are given relative to α-Fe at room temperature. The IMSG software (Douvalis et al., 2010) was used for the fitting of the spectra. Before setting, the samples were prepared by adding a drop of paste in the holder and freezing this in liquid nitrogen to solidify the material and stop the reactions. At later stages of the reaction after hardening, the materials were powdered and pressed in the sample holder before freezing and loading into the spectrometer. The effectiveness of this freezing to stop the reactions was tested by measuring a sample that reacted for 3 h and for 3 h + 14 days in liquid nitrogen. These two measurements resulted in spectra with no experimental difference and led to the conclusions that: (i) the reactions are halting after freezing the samples at 77 K and (ii) the samples do not change through the duration of the measurement, which has a typical length of 24 h.

## RESULTS AND DISCUSSION

Isothermal calorimetry measures the heat flow, proportional to the reaction rate, coming from the reactions in the sample. It can thereby provide the time-frame in which the different reaction stages occur. The heat flow, **Figure 1**, shows the exotherm which in geopolymers is often used to indicate the timing of the polymerization reactions (Rahier et al., 2007). The maximum heat flow in **Figure 1** is observed around 4 h, which was also approximately the time of setting; the main part of these polymerization reactions seems to occur between 3 and 24 h

after mixing. It should be mentioned that in the Fe-rich IPs the Fe2<sup>+</sup> in the presented slag is oxidized to Fe3<sup>+</sup> during and/or after the synthesis (Onisei et al., 2018; Simon et al., 2018). A contribution from this reaction is not distinguished here or in previous studies (Kriskova et al., 2015; Onisei et al., 2015). The oxidation most probably occurs simultaneously with the dissolution or polymerization, or occurs on a longer timescale and it is not noticed as a separate phenomenon in the calorimetry measurements. Important to note is that no crystalline components are consumed or formed by the reactions; previous work showed that no changes are observed in X-ray diffraction, apart from the appearance of carbonate phases in certain mixtures (Machiels et al., 2014; Iacobescu et al., 2017).

The kinetics of the oxidation of iron are elucidated with <sup>57</sup>Fe Mössbauer spectroscopy. **Figure 2** presents the raw spectra of the slag and IPs with curing times up to 28 days. There it can be seen that the slag spectrum is dominated by contributions with IS values of ∼1.0 mm/s and quadrupole splitting (QS) values of ∼2.5 mm/s. This is an indication that the main part of the iron in this sample is of Fe2<sup>+</sup> character (White et al., 2010; Cuesta et al., 2019), however the asymmetry of the lower and higher velocity resonant line intensities suggests that there must be at least two Fe2<sup>+</sup> contributions that should have different characteristics, and in particular different QS values. Moreover, a minor peak at ∼1.0 mm/s suggests the presence of an additional minor Fe3<sup>+</sup> state. Thus, the lowest number of components that could fit the slag spectrum was three, two different Fe2<sup>+</sup> (S1 and S2) and one Fe3<sup>+</sup> (S3). The naming of these components is composed of "S" referring to Slag and a number. Similarly, new components in the IPs will be called IP1 and IP2. The fitting result using this model for the slag, in which components S1, S2, and S3 were allowed a Gaussian-type symmetric spreading of their QS values (1QS) to simulate the observed broadening of the absorption resonant lines, is given in **Figure 3** and the Mössbauer hyperfine parameters are presented in **Table 2**. The IS and QS values of S1 and S2 correspond to high spin (S = 2) Fe2<sup>+</sup>

in oxides or silicates (Greenwood and Gibb, 1971; Coey, 1984), while the corresponding values for the minor S3 lie at the upper limit of high spin (S = 5/2) Fe3<sup>+</sup> ions found in such compounds (Greenwood and Gibb, 1971; Coey, 1984). This result shows that the particular slag contains a majority of Fe2<sup>+</sup> combined with a minority of Fe3<sup>+</sup> species.

Moving to the IP sample spectra, a more detailed observation of **Figure 2** (insert dashed rectangles) reveals the appearance and increase of the intensity of one additional absorption contribution at ∼0.75 mm/s and a second at ∼2.50 mm/s as the reaction time increases. These are evidences for the appearance of two additional components, one of Fe3<sup>+</sup> and one of Fe2<sup>+</sup> character respectively, during the reaction of the slag with the activating solution. Consequently, the fitting of the IP spectra could be performed by keeping the three contributions of the slag (S1, S2, S3) and adding one new Fe2<sup>+</sup> (IP1) and one new Fe3<sup>+</sup> (IP2) contribution to the fitting model, which represent the iron ions in the binder part of the newly formed IP. Gaussian spreading of the quadrupole splitting was allowed for components IP1 and IP2, i.e., the 1QS was refined by the fitting procedure. The parameters of the slag components (S1, S2, and S3) were fixed to the values resulting from fitting the

slag spectrum, the intensities and absorption areas (AAs) of components S1 and S3 were made dependent on the major S2 component to result in the same ratio of AAs as in the slag and only the intensity of component S2 was left to vary in the fits. As such, the IP spectra are fitted with a combination of the slag spectrum and the new Fe2<sup>+</sup> and Fe3<sup>+</sup> components, IP1 and IP2 respectively. This model thus embeds the idea that the part of the slag that has not reacted with the activating solution to form the binder phase of the IP, is conserved in the configuration it had before the start of the reactions. The described model was used to fit the IP spectra with reaction times equal or longer than 6 h, as for the frozen solutions with reaction times shorter than 6 h the spectra are essentially the same as the slag spectrum within the experimental tolerance limits and according to their fitting results. Example fitting results are shown in **Figure 4**.

The Mössbauer parameters of the two Fe2<sup>+</sup> sites in the slag, S1 and S2, show similarities with the slag components in our previous work (Onisei et al., 2018). However, the composition of the slag there was found to be close to fayalite, while the IS and QS values here are close to those obtained for pyroxenes toward the ferrosilite composition, in which two sites with varying QS values are observed (Greenwood and Gibb, 1971; Coey, 1984). The spread in the QS values for components S1 and S2 is quite large, showing that this is representing the glassy phase observed with XRD. Component S3 shows relatively large IS and QS values for an Fe3<sup>+</sup> state as well as a large 1QS, it is thus another component in the glassy phase of the slag. Based on the absorption areas, the Fe3+/6Fe ratio of the slag was 0.14.

The Fe3<sup>+</sup> in the IP binder (component IP2) shows IS and QS values similar to the Fe3<sup>+</sup> component of the binder in previous

binder gradually appear.



*IS is the isomer shift relative to* α*-Fe at room temperature, QS is the central value of the quadrupole splitting,* ∆*QS is the total Gaussian-type spreading of the quadrupole splitting, and AA is the relative absorption area. The half line-width was fixed to* Γ */2* = *0.14 mm/s to avoid interference with the* 1*QS. Typical errors include* ±*0.02 mm/s for IS and QS and* ±*5% for AA.*

work on final samples (Onisei et al., 2018) and thus suggest that an average of 5-fold or 4-fold coordination with rather high Fe-O distances is present. The low QS with respect to other glasses (Mysen, 2006) shows, for this high spin Fe3+, that the IP binder contains a higher symmetry of the ligands, i.e., a higher similarity of the species connected to Fe3<sup>+</sup> or more symmetrically distributed around it (Greenwood and Gibb, 1971; Coey, 1984). This might be an indication of the participation of these Fe3<sup>+</sup> ions in the silicate network as network formers, charged balanced by Na+, as the network forming character would mean (for the example of a tetrahedron) that Fe3<sup>+</sup> is surrounded by four bridging oxygen atoms, instead of three bridging and one nonbridging. The new Fe2<sup>+</sup> site, component IP1, has not been observed in previous measurements (Onisei et al., 2018; Simon et al., 2018), although similar Fe3<sup>+</sup> sites as component IP2 have been found. This Fe2<sup>+</sup> state thus seems to have been oxidized in later reaction stages or after milling these samples to a powder. Component IP1 shows larger IS and QS values than the Fe2<sup>+</sup> components of the slag (S1 and S2), suggesting the presence of distorted, asymmetric octahedra, comparable with fayalite minerals (Coey, 1984; Vandenberghe and De Grave, 2013; Onisei et al., 2018), Fe-rich smectites (Chemtob et al., 2015), micas (Ferrow, 1987), or Fe(II) hydroxides (Génin et al., 1996). The spreading of the quadrupole splitting of IP1 is lower compared to the other components, suggesting a more ordered environment. The presence of crystalline phases has however not been observed in this material with XRD (Machiels et al., 2014; Iacobescu et al., 2017), the IP2 component is thus rather present in a nano-crystalline phase or in very small crystallite clusters.

The evolution of the absorption areas of the new contributions (IP1 and IP2) are illustrated in **Figure 5**. Changes in the Mössbauer spectra do not happen immediately upon dissolution. This is not because of an insignificant extent of dissolution: The immediate start of the dissolution reactions was validated by measuring the concentration of the ions with induced coupled plasma optical emission spectroscopy (Pontikes et al., 2013). No rearrangement of the Fe environment occurs during the first 3 h, there is no oxidation or coordination change. The first changes are observed from 6 h on, meaning reactions have started between 3 and 6 h. This corresponds to the time where according to calorimetry the polymerization reactions start. The Fe2<sup>+</sup> (IP1) and Fe3<sup>+</sup> (IP2) components both emerge at this time stage. The largest increase in absorption area of components IP1 and IP2 is observed in the first 24 h, an observation which is analogous to the heat evolution. After the first 24 h, IP1 and IP2 keep increasing in area, although at a slower rate. After 28 days, an AA of 21% is observed for the new Fe2<sup>+</sup> state (IP1) and 21% for the Fe3<sup>+</sup> (IP2). At this point it is interesting to check the Na/Fe3<sup>+</sup> ratio to confirm the possibility of having Fe3<sup>+</sup> in the silicate network of the binder, like Al3<sup>+</sup> in geopolymers. This ratio is ∼1 for the mixture presented here, making it possible that Na provides the charge balance for Fe3+. The molar ratio Si/Fe3<sup>+</sup> provides information on the environment of Si in the network. This ratio is ∼3; the dominant silicate species [using Engelhardt notation (Engelhardt et al., 1982)]

are thus Q<sup>4</sup> (1Fe) and Q<sup>4</sup> (2Fe). Previous work attempted to provide such compositional information using an electron probe microanalyzer (Kriskova et al., 2015; Iacobescu et al., 2017; Onisei et al., 2018), but because of experimental limitations (e.g., interaction volume, sodium evaporation, etc.) and the inhomogeneity at the nanoscale (see for instance the two different states of Fe), no relevant compositional information could be derived from these experiments. The ratios calculated in the present paper are thus not confirmed by another experimental technique. Future work, for instance considering transmission electron microscopy coupled with electron dispersive X-ray spectroscopy might provide a more accurate and detailed view on the binder phase composition.

The kinetics of the evolution of the AA of components IP1 and IP2 follows an interesting sequence. IP1 seems to form faster initially (**Figure 5A**), which might suggest that the structure associated with IP2 needs IP1 as nucleation site. Between 6 h and 3 days the AA of component IP1 stays approximately constant, while IP2 evolves toward its maximum (**Figure 5B**). Afterwards, also the AA of IP1 gradually goes to its maximum at 28 days. It seems that after the start of formation of IP2, the structure that contains this Fe3<sup>+</sup> component forms until the Na/Fe3<sup>+</sup> ratio of 1 is reached, after which the reactions switch again to produce the Fe2<sup>+</sup> component IP1.

The same time stages and samples are investigated with ATR-FTIR, **Figure 6**. The spectra are focused on the Si-O stretching band (700–1,200 cm−<sup>1</sup> ). The "initial" spectrum, obtained after 8 min, shows two main features. The shoulder at 850 cm−<sup>1</sup> is associated with the precursor slag, as can be derived from the dotted line representing the spectrum of the slag. The band at 984 cm−<sup>1</sup> corresponds to the sodium silicate activating solution. The dissolution of the slag is reflected by the decrease in intensity of the shoulder around 850 cm−<sup>1</sup> . This is not only observed in the initial stages (**Figure 6A**), but extends to days after mixing (**Figure 6B**). The formation of the IP network is seen in the FTIR as the growth of a shoulder around 918 cm−<sup>1</sup> and the shift of the band that was previously at 984 cm−<sup>1</sup> (**Figure 6A**, notice the transition from 3 to 6 h). Simultaneously, the shoulder at ∼1,100 cm−<sup>1</sup> disappears, indicating the consumption of the highly connected silicate species from the activating solution. The difference plot in **Figure 7** shows the changes between the starting point and the spectrum after 6 h more clearly. The emergence of the shoulder at 918 cm−<sup>1</sup> and the shift of the band at 984 cm−<sup>1</sup> can both be linked with the emergence of a new band around 950 cm−<sup>1</sup> . These transitions are associated with the observation from Mössbauer spectroscopy: the incorporation of Fe3<sup>+</sup> in the silicate network (see the calculations in previous paragraph suggesting the presence of Q<sup>4</sup> (1Fe) and Q<sup>4</sup> (2Fe) Si) and thus the change from Si-O-Si linkages in the activating solution toward Si-O-Fe linkages in the IP binder results in a lower wavenumber of the Si-O stretching band in the IPs with respect to the activating solution. The position of the maximum of the Si-O stretching band after 28 days is observed at 959 cm−<sup>1</sup> , which is similar to what is observed for Si-O-Al linkages in aluminosilicate IPs (950–1,000 cm−<sup>1</sup> ) (Rahier et al., 2007; Rees et al., 2007) and in previous work which suggested that Fe has a similar impact on the band (Peys et al., 2017; Onisei et al., 2018). In general, this shift of the stretching band to lower wavenumbers is due to the widening of the Si-O-T distance, compared to Si-O-Si, and is observed in a wide variety of materials (Henderson and Taylor, 1979).

Using the information from the above experiments it is now possible to suggest a reaction scheme for the Fe-rich IPs. A schematic representation of the reactions is shown in **Figure 8**. After dissolution, Fe-silicate species are present in the pregnant activating solution. The concentration of these species rises until polymerization takes place into the partially oxidized binder, which consists of tetrahedral Fe3<sup>+</sup> and octahedral Fe2<sup>+</sup> of which at least one should be present in the silicate framework. For the mixture studied in the present paper the starting time of polymerization-oxidation should be between 3 and 6 h at 20◦C. The dissolution of the slag and polymerization and oxidation is

FIGURE 5 | Evolution of the absorption areas of components IP1, IP2, and the sum of absorption areas of components S1, S2, and S3. (A) Shows the short term (≤12 h) evolution of the absorption areas of the IP1 and IP2 components. The IP binder components only start forming after 3 h (0.125 days). (B) Shows the evolution of the absorption areas up to 28 days of curing.

FIGURE 6 | The FTIR spectra during the first 28 days of reaction. The spectra before setting are shown on the left in (A), these thus consider measurements in the liquid state, with the liquid in direct contact with the ATR crystal. The measurements on powdered samples after setting, pressed against the ATR crystal, are shown on the right in (B).

ongoing until at least 28 days after mixing, however, it is very likely that the time-scale for finalizing the reactions is longer than the 28 days presented here. A comparison with previous work (Onisei et al., 2018; Simon et al., 2018), where the contribution of the Fe3<sup>+</sup> binder component was much larger and no Fe2<sup>+</sup> binder component was observed, shows that a more complete oxidation happens in later stages of the reaction or during intensive milling.

Future work will study in detail the transition from the 28 days study presented here, revealing a combination of IVFe3<sup>+</sup> and VIFe2<sup>+</sup> in the binder, toward the fully oxidized structure observed in previous work (Onisei et al., 2018; Peys et al., 2018b; Simon et al., 2018).

## CONCLUSIONS

The evolution of the molecular structure during synthesis of Ferich slag based IPs has been monitored with <sup>57</sup>Fe Mössbauer spectroscopy and infrared spectroscopy. Fe-rich slags mainly consist of Fe2<sup>+</sup> (85% absorption area in Mössbauer spectrum), whereas an IP binder was shown to contain both Fe2<sup>+</sup> and Fe3+. The first appearance of a new Fe3<sup>+</sup> band in Mössbauer spectroscopy occurs after 6 h of reaction at 20◦C, and the

main oxidation happened during the first 24 h. It is suggested that this Fe3<sup>+</sup> is situated in tetrahedral coordination in the silicate framework. At the same time, a novel octahedral Fe2<sup>+</sup> component is found to form in the binder phase with Mössbauer parameters similar to fayalite or clay minerals and which approximately follows the same kinetics of formation as the Fe3<sup>+</sup> state. The fact that these new components only develop from 6 h after mixing and are mostly occurring within the first 24 h is in line with the onset of the polymerization exotherm in calorimetry, between 3 and 6 h, and the fact that the largest part of the heat from the reactions evolved in the first 24 h. ATR-FTIR spectroscopy suggested the incorporation of Fe in the silicate network. At 28 days, a total of 42% of the absorption area (AA) in the Mössbauer spectrum was attributed to the binder phase, comprised 21% of Fe2<sup>+</sup> and 21% of Fe3+, compared to 86% Fe2<sup>+</sup> and 14% Fe3<sup>+</sup> in the precursor slag. This results in a Na/Fe3<sup>+</sup> ratio of ∼1 in the mixture from which the binder phase is formed, revealing an important similarity with geopolymers where a Na/Al3<sup>+</sup> ratio of 1 is observed. The Fe2<sup>+</sup> binder state was not observed previously, whereas the Fe3<sup>+</sup> binder state always had a higher absorption area in previous work. This is probably associated with oxidation at later stages or/and with the milling procedures performed in previous work which increase the surface of the material exposed to air and consequently to oxidation.

## REFERENCES


## DATA AVAILABILITY

The datasets generated for this study are available on request to the corresponding author.

## AUTHOR CONTRIBUTIONS

AP lead all parts of the paper making process. AD supported in designing, performing, analyzing, and discussing the <sup>57</sup>Fe Mössbauer spectroscopy measurements. VH helped in designing, analyzing and discussing the FTIR, and performed the measurements. HR, BB, and YP are the supervisors of AP and provided feedback, and inspiration in all aspects.

## FUNDING

AP is grateful to the Research Foundation–Flanders (FWO) for the Ph.D. grant. VH is thankful to VLAIO for the Baekeland scholarship.

## SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fmats. 2019.00212/full#supplementary-material

ye'elimite-based pastes. Adv. Cem. Res. 31, 138–146. doi: 10.1680/jadcr.1 8.00097


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Peys, Douvalis, Hallet, Rahier, Blanpain and Pontikes. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Microstructure and Composition of Hardened Paste of Soda Residue-Slag-Cement Binding Material System

Rongjie Song1,2, Qingxin Zhao1,2, Jinrui Zhang<sup>3</sup> \* and Jizhong Liu1,2

*<sup>1</sup> State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, China, <sup>2</sup> Key Laboratory of Green Construction and Intelligent Maintenance for Civil Engineering of Hebei Province, Yanshan University, Qinhuangdao, China, <sup>3</sup> State Key Laboratory of Hydraulic Engineering Simulation and Safety, Tianjin University, Tianjin, China*

In order to utilize industrial waste soda residue, a kind of binding material system based on soda residue-slag-cement has been developed. The microstructure and composition of hardened paste of the binding material system after curing for 28 days was analyzed by means of scanning electron microscope (SEM), energy dispersive X-ray spectroscopy (EDS), X-ray diffraction (XRD), and thermogravimetric-differential scanning calorimetry (DSC-TG). The main hydration products of binding material system after curing for 28 days were C-S-H gel and 3CaO·Al2O3·CaCl2·10H2O crystal (Friedel's salt) with a needle bar and irregular shape. The Ca/Si ratio of C-S-H gel in hydration products varied from 0.7 to 1.4, the ratio of n(Cx-S-Hy)/n(3CaO·Al2O3·CaCl2·10H2O) varied in the range of 1.8–6.4. The C-S-H gel and Friedel's salts that densely arranged and closely bonded with pores and cracks constituted the main strength source of the whole soda residue-slag-cement binding material system.

#### Edited by:

*John L. Provis, University of Sheffield, United Kingdom*

### Reviewed by:

*Ionut Ovidiu Toma, Gheorghe Asachi Technical University of Ia ¸si, Romania Marija Nedeljkovic, Delft University of Technology, Netherlands*

> \*Correspondence: *Jinrui Zhang jinrui.zhang@tju.edu.cn*

#### Specialty section:

*This article was submitted to Structural Materials, a section of the journal Frontiers in Materials*

Received: *09 October 2018* Accepted: *14 August 2019* Published: *28 August 2019*

#### Citation:

*Song R, Zhao Q, Zhang J and Liu J (2019) Microstructure and Composition of Hardened Paste of Soda Residue-Slag-Cement Binding Material System. Front. Mater. 6:211. doi: 10.3389/fmats.2019.00211* Keywords: soda residue, slag, binding material, hydration products, Friedel's salt

## INTRODUCTION

The rapid development of industry has driven the world economy. However, the resulting industrial waste occupied the land and polluted the soil, water and air, which caused pollution to the global environment (Chen, 2018). Researchers around the world have conducted many studies on the reuse of industrial wastes. There are various types of industrial by products, such as calcium carbide slag, blast furnace slag, fly ash, and so on. The technique of using calcium carbide slag as construction materials and alkaline activator has been successfully used by researchers (Bilondi et al., 2018; Hanjitsuwan et al., 2018; Siddiqua and Barreto, 2018). Zhang et al. (2017) used fly ash and high-magnesium nickel slag as solid materials to manufacture geopolymer cement under the room temperature conditions. The results showed that the compressive strengths of the geopolymer cements can achieve 60 MPa, which is comparable to hardened Portland cements and adequate for construction purposes. Wiemes et al. (2017) conducted an experiment with the incorporation of different types of industrial waste in brick manufacturing process. Three types of wastes were mixed with clay: automotive industry waste sludge, glass waste and wood ash. The results showed that brick production can be obtained with different types of industrial waste. Amin et al. (2017) researched chloride resistant cement from industrial waste through geopolymerization, the test showed that the highest compressive strength of cement was 18.85 MPa and resistance to chloride attack was recorded for geopolymer with silica to alumina ratio of 2.7. Meanwhile, the resistance of geopolymer to magnesium chloride is greater than that of sodium chloride, and the attack of chloride cannot affect the geopolymer network when having optimum silica to alumina ratio. Self-compacting concrete was prepared by El Mir and Nehme (2017) utilizing industrial waste perlite powder, the results indicated that waste perlite powder had a significant pozzolanic effect on the concrete microstructure, resulting in a positive impact on the compressive strength of the concrete.

Soda residue was an industrial waste produced in the process of soda ash production by ammonia-soda process, whose chemical compositions are CaCO<sup>3</sup> and soluble salt chloride. The storage and emission of soda residue had caused serious environmental problem (Yan, 2015; Gomes et al., 2016). The accumulation of soda residue will change the acidity and alkalinity of the soil in the region, cause soil salinization and affect the normal growth of vegetation. As time goes by, chloride ions in soda residue will gradually infiltrate into the ground, causing different degrees of pollution to underground water sources and seriously threatening the safety of surrounding residents' drinking water. The soda residue mountain formed by the accumulation of soda residue will occur landslides under the action of earthquake and other external forces, endangering the safety of people's lives and property. How to use soda residue has become the focus of attention in the engineering field. Therefore, a great deal of research work has been carried out by scholars. Uçal et al. (2018) investigated hydration characteristics of alinite cement, which was produced by using soda waste sludge as raw material. It was found that induction period of alinite cement was around 15–20 min and unlike Portland cement, hydration did not practically stop during this period. Up to 12% gypsum addition to alinite cement resulted in increased compressive strength. Zhao et al. (2019) had studied the crucial properties of fly ash-based geopolymers incorporating soda residue. The research showed that the compressive strength of the alkaliactivated fly ash-based samples incorporating about 20% soda residue curing for 60 and 180 days was 13.5 and 18.0 MPa, the geopolymers had low shrinkage and good thermal stability. Zhang et al. (2013) had prepared the concrete with some soda residue replacing the cement by mass. The mix proportion of the concrete was 45–49% of stone, 25–29% of sand, 10–14% of cement, 5–10% of soda residue, and 5–8% of water. The research showed that the concrete had good frost resistance and wear resistance. Liu et al. (2017) had researched the property of polymer after mixing soda residue and fly ash. The research showed that when the mass ratio of soda residue and fly ash was <0.36, the flexural strength and compressive strength of the fly ash-NaOH system increased with the proportion of soda residue, the addition of soda residue had a significant modification effect on the development of the system strength. Sun and Gu (2014) prepared a new soda residue curing agent by using soda residue, slag and fly ash as the main raw materials and a proper amount of desulphurization gypsum and compound activator were added to these main materials. The physical and mechanical property of the curing agent basically met the technical parameters of the P.C 32.5 composite cement. Liu et al. (2016) researched the baking-free brick that were prepared with soda residue as the main material. It was found that when soda residue, fly ash, steel slag, and stone powder were prepared according to the mass ratio of 4:3:2:1, the compressive strength of baking-free brick curing for 7 days was 14.0 MPa and was 20.6 MPa curing for 28 days. Sun et al. (2012) improved the expansive soil with soda residue as an additive. The research showed that the relative density, liquid limit, plastic index, free expansion rate, and load expansion of the expansive soil decreased noticeably with the increase of soda residue content, which indicated that the soda residue had obvious improvement effect on expansibility of expansive soil.

However, the consumption of soda residue still cannot meet the needs of society and environment. In order to expand the application scope and further reduce the pressure on the environment, a new kind of binding material system based on soda residue-slag-cement was developed through a series of experiments. The microstructure and composition were analyzed by the scanning electron microscope (SEM), energy dispersive X-ray spectroscopy (EDS), thermogravimetricdifferential scanning calorimetry (DSC-TG), and X-ray diffraction (XRD).

## EXPERIMENTAL

## Raw Materials

The soda residue used in the experiment came from San You group in Tangshan, Hebei, which was dried and crushed into fine powder particles with the maximum diameter of 0.16 mm. The XRD pattern was shown in **Figure 1**. It was apparent that the peak of CaCO3, CaSO4, CaCl2, and NaCl appeared in **Figure 1**, which indicated that CaCO3, CaSO4, CaCl2, and NaCl were contained in the identified phases in the XRD diffractogram of soda residue. The slag in the experiment was S95 grade granulated blast furnace slag powder. The cement used in the soda residue-slagcement binding material system was Portland cement (P. II 42.5), cement used in preparing cement mortar specimen for strength comparison was ordinary Portland cement (P.O 42.5R). Both slag and cement came from Qinhuangdao municipal construction group co., LTD in Hebei. The chemical components of soda residue, slag and cement were shown in **Tables 1**, **2** respectively. The particle size of nature quartz river sand ranged from 0.16 to 2.36 mm.

## Preparation of Mortar Specimens

In the mortar specimen mix design, the water-binder ratio was fixed at 0.5, the sand-binder ratio was fixed as 3:1 and the mix proportions of the binding materials were shown in **Table 3**. The cement mortar was stirred evenly and the specimens of 40 mm × 40 mm × 160 mm were prepared. All mortar specimens were cured in the standard curing box of a controlled environment of (20 ± 2)◦C and relative humidity ≥ 95% for 28 days and then used to test the flexural strength and compressive strength. Each group was molded into six mortar specimens, which were then taken for average calculation.

## Methods

## SEM and EDS Test of Paste Specimens

According to the test method described in the literature (Aughenbaugh et al., 2016), the paste specimens of 20 mm × 20 mm × 80 mm were prepared according to the mix proportion in **Table 3**, and then removed into containers of a controlled environment of (20 ± 2)◦C and relative humidity ≥ 95% to be cured to the ages of 28 days. The microstructure of the hardened pastes curing for 28 days was observed by a Hitachi-3400N SEM after the vacuum and gold spray treatment, and the content of the elements of Ca, Si, and Cl in hydration products was analyzed by the energy dispersive EDS, then the ratio of Ca/Si in the C-S-H gel and the ratio of Si/Cl in the hydration products were calculated.

#### TABLE 1 | Chemical composition of soda residue w/%.


TABLE 2 | Chemical compositions of cement and slag w/%.




### Mineralogical Composition

The component of hydration product was determined by XRD. The paste specimens were taken from the non-carbonated zone. During the testing period, the specimens were processed as following steps. Firstly, the non-carbonated specimens were split into small pieces and put in ethanol and were ground into fine powder particles till they all passed the sieve with the diameter of 0.12 mm. Secondly, the powder was filtered in a vacuum by the suction filter bottle and diluted five times with ethanol and twice with ether. Then, the powder particles were dried in a vacuum drying oven with desiccant at 105◦C for 24 h. Finally, the powder particles were placed in the small glass bottles to be used for XRD. The 2θ values ranged from 10◦ to 80◦ and the scanning rate applied was 1◦ /min for all specimens.

## DSC-TG Test of Hydration Product Component

Thermogravimetry-differential scanning calorimetries (DSC-TG) of the paste specimens curing for 28 days were carried out by the STA 449F5 integrated thermal analyzer (Zhao et al., 2016). Firstly, the soda residue, slag and cement were dried in an oven at 105◦C for 12 h, and the coarse particles over 0.9 mm were removed by a standard sieve. Secondly, the flesh pastes were prepared based on mix proportions in **Table 3,** and the waterbinder ratio was 0.5 for all the pastes. Then the flesh pastes were sealed in plastic bags and then cured under the environment of 20 ± 2 ◦C and relative humidity ≥ 95%.

At the age of 28 days, the soda residue-slag-cement paste specimens were processed as per the following steps. Firstly, the specimens were split into small pieces and put in ethanol, and were then crushed into fine powder particles until all passed the sieve with the diameter of 0.12 mm. Secondly, the powder was filtered in a vacuum by the suction filter bottle and diluted five times with ethanol and twice with ether. Then the powder samples were dried in a vacuum drying oven with desiccant at 105◦C for 24 h. Finally, DSC-TG of the paste specimens curing for 28 days was carried out by the STA 449F5 integrated thermal analyzer. To prevent the carbonization of the specimens during the heating process, the Ar was used as the protective gas and the temperature rising rate was 10◦C/min.

## RESULTS AND DISCUSSION

## Mechanical Properties of Mortar Specimens

**Table 3** shows the flexural strength and compressive strength of mortar specimens after curing for 7 days and 28 days, with different percentages of binding materials. It can be seen from **Table 3** that the mechanical property of D2 group curing for 28 days was optimum when the soda residue, slag and cement were prepared as the mass percentage of 2:6:2. **Figure 2** shows the compressive strength of specimens at the age of 7, 28, 60, 90, and 120 days under the ratio of D2 group and the compressive strength of P.O 42.5R cement mortar specimen at the same age. It can be seen from **Figure 2** that the compressive strength of mortar specimen of D2 group curing after 28 days was quite close to the strength of mortar specimens of P.O 42.5R cement.

It was apparent that compressive strength of D2 group curing for 7 days was 15.3 MPa, and the value was much lower than that of P.O 42.5R cement mortar specimens at age of 7 days. Slag is a potentially active cementitious material. The alkaline environment of soda residue makes the slag hydrate slowly. At the beginning, less hydration products are generated, and the system is less dense. With increasing age, the number of hydration products increases and the density increases gradually, and the compressive strength at 28 days increases significantly, reaching 45.1 MPa.

## Micro-Morphology Analysis of Hydration Products

**Figure 3** shows the SEM photographs of paste specimens of D2 group curing for 28 days. **Figures 3A–D** indicated that the hydration products of D2 group cured to the ages of 28 days were rich, and the microstructure of hardened paste was formed basically. The hydration products were closely bonded with pores and cracks. The crystals were rod-shaped and irregularly shaped, and were distributed on the surface of hydration products. At the same time, a small amount of unreacted binding materials particles were distributed around the hydration products.

## Energy Dispersive Spectrum Analysis of Hydration Products

Slices of paste of D2 group curing for 28 days were impregnated and polished to prepare BSE samples as shown in **Figure 3E**. In the SEM-BSE photographs, the detection dots were marked, respectively, in the position of different brightness in the image, and EDS analysis was carried out on every detection dot, respectively. The marked position was shown in **Figure 3F** and the detection results were shown in **Figure 4** and **Table 4**.

It can be seen from **Table 4** that the content of Cl elements at each detection position was far higher than the content of the S elements in the hydration products of D2 group. Therefore, it was inferred that there was a compound containing Cl elements in the hydration products. Meanwhile, it can be seen that the content of Si elements in the detection range was relatively large, because the diameter of the spectrum analysis range was larger than the size of the compound containing Cl elements, thus causing the existence of C-S-H gel in the various measuring regions. In the analysis of elements, it was assumed that all Si elements came from C-S-H gel and the hydration product containing Cl elements was 3CaO·Al2O3·CaCl2·10H2O crystal (Shao et al., 2013). Then the numbers of atoms of the Ca and Cl elements were subtracted from the numbers of atoms of the element in the measured region according to the elemental composition of 3CaO·Al2O3·CaCl2·10H2O, subsequently, the numbers of atoms of Ca and Si elements in C-S-H gel were calculated. The result of the atomic weight of Ca and Si elements in C-S-H gel were shown in **Table 5**. It can be seen from **Table 5** that the ratio of Ca/Si in C-S-H gel on the spot 1 was 0.7, the ratio of Ca/Si in C-S-H gel on the spot 2 was 1.2, the ratio of Ca/Si in C-S-H gel on the spot 3 was 1.1, and the ratio of Ca/Si in C-S-H gel on the spot 4 was 1.4 when the hydration product containing Cl element was 3CaO·Al2O3·CaCl2·10H2O crystal. The ratio of Ca/Si of the remaining portion in detection region was within the range of Ca/Si ratio of C-S-H gel, thus, it was proved that the hydration product containing Cl elements was 3CaO·Al2O3·CaCl2·10H2O crystal.

In order to characterize the relative content of C-S-H gel and 3CaO·Al2O3·CaCl2·10H2O crystal in the hydration product of D2 group, the ratio of Si/Cl was introduced, the ratio of which was shown in **Table 5**. It could be seen from **Table 5** that the ratio of Si/Cl was 3.0 on the spot 1, the ratio of Si/Cl was 1.8 on the spot 2, the ratio of Si/Cl was 0.9 on the spot 3, and the ratio of Si/Cl was 3.2 on the spot 4. It was assumed that all the Si elements were

derived from the C-S-H gel and all the Cl element were derived from the 3CaO·Al2O3·CaCl2·10H2O crystal. Meanwhile, the content of Si element in the C-S-H gel was the benchmark, it was known that the ratio of n(C-S-H)/n(CaO·Al2O3·CaCl2·10H2O) on spot 1, spot 2, spot 3, and spot 4 was 6.0, 3.6, 1.8, and 6.4, respectively. Each detection position was representative when the detection marker was made on the position of different brightness in the EDS spectrum analysis, so the data measured at four positions also represent the overall situation of the hydration products. According to the analysis above, the variation range of n(Cx-S-Hy)/n(3CaO·Al2O3·CaCl2·10H2O) was 1.8–6.4 in the hydration products. That is, the ratio of molar content of Cx-S-H<sup>y</sup>

gel to the molar content of CaO·Al2O3·CaCl2·10H2O crystal was fluctuated in the range of 1.8–6.4. Therefore, C-S-H gel and CaO·Al2O3·CaCl2·10H2O crystals that densely arranged and closely bonded with pores and cracks were the main contributors to the strength of the whole soda residue-slag-cement binding material system.

To verify the accuracy of the analysis above, the detection points were marked and the energy dispersive spectrum analysis was carried out on different position of brightness. The marked position was shown in **Figure 3F** and the detection results were shown in **Table 4** and **Figures 4E,F**. According to the analysis method above, it can be calculated that the ratio of Ca/Si in C-S-H

gel on the spot 5 was 0.75 and the ratio of Ca/Si in C-S-H gel on the spot 6 was 1.3. The ratio of Ca/Si of the remaining portion in detection region was within the range of Ca/Si ratio of C-S-H gel, thus, it supported the conclusion above that the hydration product containing Cl elements was 3CaO·Al2O3·CaCl2·10H2O crystal. It could be seen from **Table 5** that the ratio of Si/Cl on spot 5 and spot 6 was 2.05 and 2.9, respectively. In other words, the ratio of n(C-S-H)/n(CaO·Al2O3·CaCl2·10H2O) on spot 5 and spot 6 was 4.1 and 5.8, respectively. The ratio of n(C-S-H)/n(CaO·Al2O3·CaCl2·10H2O) on the two measuring surfaces were within the range of 1.8–6.4, which support the conclusion that the ratio of molar content of Cx-S-H<sup>y</sup> gel to the molar content of CaO·Al2O3·CaCl2·10H2O crystal was fluctuated in the range of 1.8–6.4.

TABLE 4 | Elementary compositions of detection spot.


TABLE 5 | Calculation results of Si, Ca, and Cl (by atomic) for spot.


## XRD Analysis of Hydration Products

**Figure 5** shows the XRD patterns of paste specimens of D2 group cured to the age of 28 days. It was apparent that the peak of paste specimens (3CaO·Al2O3·CaCl2·10H2O) appeared in **Figure 5**. As shown in the reference (Baquerizo et al., 2015), 3CaO·Al2O3·CaCl2·10H2O, also known as Friedel's salt, which belongs to a compound of the AFm family. A variety of compounds contained in the AFm family, in which 3CaO·Al2O3·CaSO4·nH2O, 3CaO·A2O3·CaCO3·nH2O, and 3CaO·Al2O3·CaCl2·nH2O were the typical compounds. A variety of ions will replace each other to produce complex compounds when a large number of mineral admixtures existed. AFm crystal was a six square sheet crystal, but it tended to show tiny needle rod and irregular sheet shape when it grew incompletely. As shown in **Figure 3C**, the crystal in the shape of needle rod and irregular sheet distributed on the surface of hydration products was the 3CaO·Al2O3·CaCl2·10H2O crystal.

It was apparent that the peak of paste specimens (CaCO3) appeared in **Figure 5**. The presence of CaCO<sup>3</sup> may be caused by carbonation of the Ca(OH)<sup>2</sup> during the sampling and specimens preparation. Although the C-S-H gel belongs to the amorphous state and the spectral line was mainly diffuse peak, but there was a prominent "convex hull" phenomenon at 25–35◦ in the XRD patterns, indicating that there was the C-S-H gel in the hydration products of the soda residue-slag-cement binding material system curing for 28 days (Wu et al., 2014).

## DSC-TG Analysis of Hydration Products

The DSC-TG patterns of paste specimens of D2 group were shown in **Figure 6**. It can be seen from **Figure 6** that there were seven main heat absorption peaks and two main exothermic peaks. The endothermic peaks that located, respectively, at 106.9, 186.4, 299.3, 334.2, 467, 609.9, and 672.6◦C were corresponding to the temperature of decomposition of the hydration products. According to the literature (Yang and Xue,

2000), the endothermic peak at 106.9◦C was due to the removal of water molecule in 3CaO·Al2O3·CaCl2·10H2O, and forming 3CaO·Al2O3·CaCl2·6H2O. The endothermic peak at 299.3◦C resulted from 3CaO·Al2O3·CaCl2·6H2O continued taking off 5 crystalline water. The endothermic peak at 186.4, 334.2, and 467◦C were caused by dehydration and heat absorption of C-S-H gel. The crest at 609.9◦C was due to the continuous hydration and heat absorption of calcium aluminum chloride hydrate. The existence of CaCO<sup>3</sup> in the paste specimens was showed in the XRD patterns. Most of the carbonate mineral decomposition temperature are in the range of 600–900◦C. It can be seen that there was an endothermic peak at 859.5◦C which was due to the decomposition reaction of CaCO<sup>3</sup> resulting in releasing CO<sup>2</sup> and generating CaO. The exothermic peaks were at the position of 717.6 and 904.9◦C. The exothermic peak at 717.6◦C was due to the further crystallization of calcium aluminum chloride hydrate. The exothermic peak at 862.2◦C was due to the crystal transformation of the corresponding C-S-H gel and the transformation into the β-wollastonite (Wang and Yan, 2008, 2011). By the analyzation of DSC-TG, the existence of C-S-H gel and 3CaO·Al2O3·CaCl2·10H2O crystal in hydration products was evidenced.

## REFERENCES


## CONCLUSIONS

The icrostructure and composition of hardened paste of the soda residue-slag-cement binding material system were analyzed in this study. Several conclusions could be drawn as follows.


## AUTHOR CONTRIBUTIONS

RS: experimental design, investigation, methodology, and writing. QZ: funding acquisition, investigation, methodology, project administration, and writing. JZ: conceptualization, data analysis, supervision, validation, writing, and revising. JL: data analysis, methodology, and writing.

## FUNDING

Financial support from the National Natural Science Foundation of China under the grants of 51578477 and 51708403, the Innovative Research Group Science Fund of the National Natural Science Foundation of China under the grant of 51621092, the Key R&D Project of Hebei Province under the grant of 19211505D, and the China Postdoctoral Science Foundation under the grants of 2018T110200 and 2018M640236 was gratefully acknowledged.


J. Constr. Build. Mater. 162, 714–723. doi: 10.1016/j.conbuildmat.201 7.12.034


**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Song, Zhao, Zhang and Liu. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.