Hydration of Reactive MgO as Partial Cement Replacement and Its Influence on the Macroperformance of Cementitious Mortars

Gonçalves, T.; Silva, R. V.; de Brito, J.; Fernández, J. M.; Esquinas, A. R.

doi:https://doi.org/10.1155/2019/9271507

Advances in Materials Science and Engineering

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Abstract Introduction Materials and Methods Results and Discussion Conclusions Data Availability Conflicts of Interest Acknowledgments References Copyright Related Articles

Research Article | Open Access

Volume 2019 | Article ID 9271507 | https://doi.org/10.1155/2019/9271507

Hydration of Reactive MgO as Partial Cement Replacement and Its Influence on the Macroperformance of Cementitious Mortars

T. Gonçalves,¹R. V. Silva,²J. de Brito,²J. M. Fernández,³and A. R. Esquinas⁴

Academic Editor: Jose M. Monzo

Received05 Nov 2018

Revised17 Jan 2019

Accepted04 Feb 2019

Published26 Feb 2019

Abstract

A recent quest for more sustainable cement-based construction materials has triggered the pursuit of technically viable alternatives of cement, making reactive magnesium oxide (MgO) one of the least known top contenders to reduce this sector’s environmental impact since it participates in the cement’s hydration reactions and presents enhanced carbon capture ability during its life cycle. In this study, two different commercially available reactive MgO samples were evaluated as partial cement replacements (at 10%, 15%, and 20%, by weight) in the production of mortars. Thermogravimetric analysis (TGA), energy-dispersive X-ray (EDX) analysis, differential thermal analysis (DTA), and powder X-ray diffraction (XRD) analysis of cement, MgO samples, and resulting mortars were carried out. All specimens were evaluated in terms of their mechanical and durability-related performance (i.e., flexural and compressive strength, carbonation, water absorption by capillary action, and shrinkage). The main results suggest that, in spite of the decreased, albeit acceptable, performance with increasing incorporation of MgO as partial cement replacement, a significant decrease was observed in the shrinkage strain of cementitious materials.

1. Introduction

Cement industry is one of the most polluting industries in the world, contributing around 5% to 7% of the total global CO₂ emission [1], thereby making it of fundamental importance to seek new solutions that can further decrease its environmental impact. Some of the solutions involve improving the calcination process energy efficiency, using alternative fuels or raw materials, and exploring the role of cement systems in carbon capture and storage technologies [2]. Among the aforementioned, the one considered the most effective from an environmental impact reduction point of view is direct replacement of cement (i.e., supplementary cementitious materials) with constituents that partake in the hydration and strength development processes, especially if sourced from industrial wastes (e.g., fly ash, ground granulated blast-furnace slag, palm oil fuel ash, and electric-arc furnace dust) [3–6]. This makes the use of reactive magnesium oxide (MgO) a unique opportunity to reduce the resulting composite’s carbon footprint, as well as to improve some of its features [7].

The use of MgO-based constituents in the clinkerization process has been restricted to no more than 5% of MgO [8], as the high temperatures involved in this process would yield the production of dead-burned MgO with a very slow hydration rate [9]. Consequently, the late expansion of MgO in a hardened and stable cementitious microstructure would jeopardize its soundness. However, this is not the case for MgO subjected to lower calcination temperatures (i.e., light burned MgO calcined at 700–1000°C). Not only does it typically present a great reactivity and early age expansion [10–12], but its overall CO₂ emissions were reported to be up to 73% lower than those of cement [13], considering the higher carbon capture capacity of MgO-based cements throughout their life cycle [7, 14, 15].

Hydration reaction of MgO, which leads to the formation of Mg(OH)₂, is as follows [16]:

Since the volume of hydration products is higher than that of the reagents, the hydration of MgO leads to an expansion, which, if properly controlled, may be used to compensate the shrinkage typically observed in cementitious composites [10, 17–21] and possibly their deformation due to creep as well. The compatible chemical nature of MgO with that of cement may produce beneficial expansive components capable of reducing the number of microcracks and contribute to strength gain under certain circumstances [22]. Similar to that observed in the hydration of cement, after the formation of Mg(OH)₂, ensuing pozzolanic reactions can lead to the formation of magnesium silicate hydrates (i.e., MgO-SiO₂-H₂O or M-S-H) [23] capable of presenting considerable strength gain [24], if in the presence of amorphous silica-bearing mineral additions, such as fly ash [25, 26] or silica fume [27, 28]. Another practical feature of using reactive MgO-based cementitious composites is their ability to capture a considerable amount of CO₂ thereby further decreasing the material’s environmental impact throughout its life cycle. Furthermore, the carbonation of brucite may result in phases (i.e., magnesium carbonate hydrates such as nesquehonite, hydromagnesite, dypingite, and artinite) capable of significantly increasing the strength of the matrix [29, 30].

There have been quite a few studies on the microstructure of cementitious composites containing reactive MgO, by means of X-ray diffraction (XRD) and thermogravimetric analyses [11, 23, 31–36], specifically on the formation of Mg(OH)₂ and M-S-H phases [16, 37]. Various reports based on SEM analyses have suggested that the hydration products of MgO exhibit a tendency to migrate towards the interfacial transition zone (ITZ) area thereby reducing the porosity of the resulting material [38–41].

In the light of the recent positive findings regarding this component and its practical impact in the construction industry, this paper reports the interaction of the hydration products of commercially available MgO, from two different sources, with those of cement, from a cementitious microstructure viewpoint and its implications on the composite’s macroperformance. This study is the first of upcoming state-funded research projects, involving various companies within the construction sector. Contrasting with the mix design of other research studies, which typically involve lower replacement levels of cement with MgO, relatively high contents of 10%, 15%, and 20%, by weight, of MgO samples were used as cement replacement thus constituting one of the novelties of this paper. These relatively higher contents were chosen with the purpose of further reducing the environmental impacts of concrete, given the considerable carbon capture of MgO-containing cementitious systems, and also its costs since the manufacture of cement is an energy-intensive process unlike that of light burned MgO. Thermogravimetric analysis (TGA), energy-dispersive X-ray (EDX) analysis, differential thermal analysis (DTA), and powder XRD analysis were carried out on cement, the two MgO samples, and selected mortars. The specimens’ mechanical- and durability-related performances were evaluated by means of their flexural and compressive strength, carbonation, water absorption by capillary action, and shrinkage.

2. Materials and Methods

2.1. Cement and MgO

The cement used was CEM I 42.5R, and its chemical composition is presented in Table 1. The cement presents initial and final setting times of 161 min and 232 min, respectively. The water used in the production of all specimens was potable tap water.

Two types of MgO were used: MgO A was supplied by Styromag, an Austrian company, and MgO G from Grecian Magnesite S.A., a Greek company. Both are light burned MgO due to their relatively low calcination temperature of about 800°C [11]. Both MgO samples’ chemical compositions are presented in Table 1. The morphology of both MgO samples can be observed by means of the scanning electron microscopy (SEM) micrographs presented in Figure 1.

(a)

(b)

The particle-size distributions of cement and both MgO samples are presented in Figure 2(a). MgO A and cement present similar distributions, as most of their particles have a diameter between 10 μm and 60 μm. MgO G, which is similar to MgO A, contains slightly smaller particles (mostly between 10 μm and 30 μm) and presents a less extensive particle-size distribution when compared to the other two. About 52.1% and 65.3% of the particles of MgO A and MgO G, respectively, were found in the range of 3 µm–32 µm. Above 32 µm (20.2% of MgO A and 3.5% of MgO G), the particles may be too large to hydrate rapidly, and below 3 µm (27.7% of MgO A and 31.1% of MgO G), they may have a smaller contribution to the mechanical strength and simultaneously demand more water [42–44]. The bulk densities of MgO A, MgO B, and cement were 653, 618, and 1040 kg/m³, respectively.

(a)

(b)

The BET specific surface areas of MgO A and MgO G were 20.12 m²/g and 18.01 m²/g, respectively. The pore size distributions, obtained from the adsorption-desorption isotherms of N₂, are shown Figure 2(b). Both additions show a range of pore diameter of 4–200 nm, according to the BJH method. Although the distribution of the pore volume is comparable in both materials, for MgO A, the highest quantity of the pore volume is centred at a slightly lower pore diameter than MgO G (23 nm vs. 33 nm, respectively).

The results of the XRD analysis are presented in Figure 3. The pattern of MgO A shows that the main phase was periclase (MgO) (International Centre for Diffraction Data-ICDD #45-0946), suggesting its crystallinity due to the narrow peak. For MgO A, peaks corresponding to the presence of calcite (CaCO₃) (ICDD #05-0586), dolomite (CaMg(CO₃)₂) (ICDD #36-0426), and talc (ICDD #13-0558) are also visible. The presence of the first two could also be observed in the TGA in Figure 4(a). Unlike the XRD pattern of MgO A, the one of MgO G shows just a periclase phase. The absence of other smaller phases indicates a lower degree of impurity when compared to MgO A.

(a)

(b)

Figure 4 presents the results of TGA and DTA of MgO A and MgO G. The initial weight loss up to 150°C, slightly higher in MgO G when compared to that in MgO A, is possibly due to the OH group residual in the surface of the material or to the presence of the hydromagnesite compound. At temperatures between 300°C and 400°C, a weight loss was observed in MgO G (4.67%), attributable to the dehydroxylation of Mg(OH)₂. This suggests greater partial hydration of MgO G in comparison with MgO A. The increase in weight loss and decrease in heat flow up to around 750°C may have been due to the decarbonation of calcite (CaCO₃) and dolomite (CaMg(CO₃)₂) as suggested by the XRD patterns of MgO A.

2.2. Characterization of Aggregates

The sand’s particle-size distribution was made in accordance with EN 1015-1 [45]. The bulk density was evaluated according to EN 1097-3 [46], whereas water absorption and apparent density as per EN 1097-6 [47]. The fine particles content classification was determined in accordance to EN 13139 [48]. The physical properties and particle-size distributions of both fine and coarse sand (0/2 and 0/4 mm, respectively) are presented in Table 2.

An XRD analysis was also carried out on the sand used in this study (Figure 5). Both fine and coarse sand showed that the main phase was quartz (crystalline SiO₂) (ICDD #33-1161) and were thus expected to have marginal reactivity with the cement’s products of hydration and Mg(OH)₂. The presence of orthoclase (ICDD #31-0966), albite (ICDD #41-1480), and microcline (ICDD #19-0926) was observed.

2.3. Mix Design

The mortar’s composition was based on the method proposed by Nepomuceno et al. [49], and the contents are presented in Table 3. The method adopted for the mix design allows the production of mortars exhibiting characteristics and performance equivalent to those of a concrete counterpart and not mortars typically used in rendering and masonry. All mortars were produced with a volumetric ratio of 1/3 (binder/aggregate). The mixes’ spread was fixed at 200 ± 15 mm, which was achieved by varying the water/binder () ratio, starting from 0.50 for the control mixes (CMs). The cement was partially replaced by each type of MgO with the ratios of 10%, 15%, and 20%, by weight. Mortars incorporating MgO A and MgO G, henceforth, are MA and MG, respectively.

2.4. Test Methods

The unbound constituents and the resulting hardened specimens were analyzed by means of XRD patterns using a Bruker D8 Discover A25 instrument with Cu-Kα radiation. All diffraction patterns were obtained by scanning using a goniometer from 10° to 80° (2θ) at a rate of 0.05°·s⁻¹. TGA was performed using a Setaram SETSYS Evolution 16/18 apparatus at a heating rate of 5°C/min. The electron microprobe technique was conducted using an electron microscope JEOL JSM-6300 at an acceleration voltage of 20 kV and a working distance of 15 mm. The X-ray detector model is Oxford Instruments ATW2-6699. The particle sizes were measured using a Mastersizer S laser diffraction particle size analyzer (Malvern Instruments) with ethanol as dispersant.

Mortars were produced and cast in accordance with EN 1015-2 [50]. Fresh mortars were tested for their consistence [51] and bulk density [52]. Hardened mortars were evaluated in terms of their flexural and compressive strength (3 and 6 specimens, respectively, per mix and age; total of 105 and 210 specimens, respectively) [53], carbonation (3 specimens per mix and age; total of 42 specimens, carbonation chamber with 5 ± 0.1% CO₂, 60 ± 5% RH, and at a temperature of 23 ± 3°C) [54], water absorption by capillary action (3 specimens per mix; total of 21 specimens) [55], and shrinkage (2 specimens per mix; total of 14 specimens) [56].

3. Results and Discussion

3.1. Physicochemical Characterization of Hardened Mortars

Figure 4 presents the TGA and DTA of control mortars and those incorporating 20% MgO A as partial cement replacement with a curing age of 28 days. The initial mass loss for temperatures up to 200°C can be attributed to the dehydration of C-S-H and AFt phases. The thermal decomposition of M-S-H usually occurs in the same range of temperatures [24] and thus cannot be clearly observed, but it is likely that this phase did not form due to lack in the presence of amorphous SiO₂ capable of reacting with Mg(OH)₂ [35]. Two drops are readily noticeable for MA20 between 300°C and 500°C. The first one, between 300°C and 400°C, can be attributed to the dehydroxylation of Mg(OH)₂. The second drop, observed in both specimens, between 400°C and 500°C, is due to the dehydroxylation of Ca(OH)₂, as a result of the cement’s hydration. The final endothermic peak at around 700°C may be attributed to the decarbonation of CaCO₃ for control mixes and both CaCO₃ and MgCO₃ for MA20 specimens.

Figure 6 presents the results of the XRD analysis of the CM and MA20 specimens. It is clear from both patterns that the main mineral phase was quartz (ICDD #33-1161), being a result of the incorporation of sand. The presence of portlandite (Ca(OH)₂) (ICDD #04-0733), from the hydration of cement, was also noted; and, to a lesser extent, calcite (ICDD #05-0586), albite (ICDD #41-1480), orthoclase (ICDD #31-0966), Ca₃SiO₅ (ICDD #42-0551), and microcline (ICDD #19-0926) were noted. For MA20, quartz (ICDD #33-1161) was also the main mineral phase. Other phases include portlandite (ICDD #04-0733), calcite (ICDD #05-0586), albite (ICDD #41-1480), orthoclase (ICDD #31-0966), Ca₃SiO₅ (ICDD #42-0551), microcline (ICDD #19-0926), dolomite (ICDD #36-0426), periclase (ICDD #45-0946), and talc (ICDD #13-0558); the last three were derived from the incorporation of MgO A.

3.2. Flexural and Compressive Strength

Figures 7(a) and 7(b) present the flexural and compressive strength, respectively, of all mixes at 3, 7, 14, 28, and 56 days. A general decrease in strength can be observed with increasing MgO content, which was also observed by other researchers assessing lower replacement levels [57]. The magnitude of this decrease in strength is more noticeable for MG specimens. This decline is a result of two factors, i.e., the need to increase the ratio and the cement’s dilution with increasing MgO content. In the former, the additional water led to a more porous microstructure for specimens containing 15% and 20% MgO. In the latter, the addition of MgO as cement replacement, in spite of the formation of Mg(OH)₂, leads to a decreasing quantity of available C₂S and C₃S and consequently of C-S-H phases [35, 58, 59]. Nevertheless, if in the presence of an addition with amorphous SiO₂, a greater rate of strength development could be observed as a result of pozzolanic reactions with Mg(OH)₂ [25]. Although the formation of Mg(OH)₂ is vital for strength gain at early curing ages, the ensuing formation of M-S-H will be the main contributing factor of the improvement of cementitious materials incorporating MgO [23]. Moreover, the existence of alumina from aluminosilicate mineral additions may also offer an advantage concerning the material’s mechanical behaviour. It may interact with Mg(OH)₂ and form a hydrotalcite-like phase [60] that is capable of enhancing the overall mechanical performance through improved bond with the cementitious matrix [23] and due to the hydrate’s ability of decreasing the porosity of the resulting composite [61].

(a)

(b)

Concerning the strength development of MA and MG, the former presented improved performance when compared to the latter. One of the reasons for this may be due to the fact that MgO A presented a more extensive particle-size distribution similar to that of cement (Figure 2), which may have improved the packing of the resulting mortar specimens. Furthermore, as a result of the smaller particle size and greater surface area of MgO G, more water was required (Table 3) to obtain mixes with comparable workability levels. Naturally, this led to a more porous and less resistant microstructure. Still, in cases wherein an additional amount of water was not required (i.e., MA10 and MG10), specimens containing MgO A exhibited lower loss in performance. The TGA on both MgO samples showed that MgO G presented a greater mass loss probably due to the dihydroxylation of Mg(OH)₂ (mass losses of 1.76% and 4.67% within the 300–400°C range for MgO A and MgO G, respectively), which suggests that it was already partially hydrated and did not contribute to the mortars’ strength gain as effectively as MgO A.

3.3. Carbonation

Table 4 presents the carbonation depth of the MgO-containing specimens after being exposed for 28 and 91 days in a CO₂ chamber. The results show greater carbonation depths with increasing MgO content, and there is no significant difference between MA and MG specimens. Even though less porous microstructures incorporating MgO have been observed elsewhere [25, 62] suggesting reduced carbonation, the opposite trend was observed in the present study, and in others [62, 63], they were substantiated by the results in Figure 8. The higher porosity in mixes containing increasing MgO content may have led to a greater surface area within the cementitious microstructure available to carbonate in the presence of CO₂. Nevertheless, this higher porosity may exist only at an initial stage as ensuing carbonation yields the formation of further products of hydration reducing the volume of the porous microstructure [62].

(a)

(b)

Another factor must also be taken into account, which is the suitability of the phenolphthalein test to ascertain the amount of CO₂ uptake in the specific case of mortars containing MgO. Since cement was partially replaced with MgO, there was a lower quantity of the former available to produce Ca(OH)₂, which exhibits a higher pH level than that of its counterpart from the hydration of MgO (i.e., Mg(OH)₂) (∼12.5 and ∼10.5, respectively) thereby reducing the specimens’ overall pH level. Nevertheless, the slightly greater mass loss of MA20 specimens in comparison with that of CM specimens between 500°C and 850°C in Figure 4 (3.71% vs. 3.34%, respectively) suggests increased CO₂ uptake and thus increased rate of carbonation.

3.4. Water Absorption by Capillary Action

Figure 8 and Table 4 present the results of the water absorption by capillary action of MA and MG. As expected, there was a slight increase in water absorption with increasing MgO content, though not as obvious for MG10 and MG15. Even though the hydration of MgO may result in the formation of a phase that exhibits higher volume than that of the initial constituents, replacing a part of the cement will directly decrease the amount of C-S-H that is capable of producing a more tortuous and less interconnected porous network [37], especially after carbonation of the specimen [62, 64]. However, lower total pore volume and water absorption would likely be observed over time with ensuing carbonation reactions and formation of additional magnesium carbonate hydrates [30, 37, 65].

3.5. Shrinkage

Figure 9 presents the shrinkage strains of MA and MG. Up to a replacement level of 15%, both MA and MG exhibited lower shrinkage strain when compared to the control. This trend was also observed for MA-20, but not for MG-20, which presented significantly higher shrinkage strain. These findings are in agreement with those of others [19, 66] and can be explained by the aforementioned reason that Mg(OH)₂ presents a greater volume than that of its initial reagents. The specimens’ expansion, prompted by the MgO particles’ rapid reaction (typical of light-burned MgO [67]), occurred 2-3 days after casting. A similar behaviour was observed by Polat et al. [68] for mixes containing up to 7.5% MgO. After that, a distinctive behaviour of drying shrinkage could be observed, nonetheless significantly offset by the initial expansion; MA-15 and MA-20 exhibited about 80% lower shrinkage strain after 90 days. Others observed similar decreases in shrinkage for MgO contents of 7.5% [68, 69]. However, for MgO G, the lower shrinkage strain could be observed only for replacement levels of 10% and 15%, with values of about 40% and 20%, respectively, lower than those of the control. This can be explained by the partial hydration of MgO G before its inclusion and thus did not contribute as effectively in the specimens’ expansion as in MgO A-bearing mortars. Furthermore, for higher replacement levels of MgO G, the water requirement also increased, which increased the specimens’ open porosity and decreased their stiffness. This facilitated the evaporation of free water and decreased the material’s ability to restrain shrinkage, which outweighed the expansion of MgO G.

(a)

(b)

4. Conclusions

An experimental investigation on the role of MgO, as partial cement replacement, in the hydration reactions was performed. The following conclusions were drawn:(i)Commercially available reactive MgO exhibits a particle-size distribution that is comparable to that of cement, which avoids further size-reduction processing stages after the manufacture of light burned MgO before its inclusion in cementitious composites.(ii)Although the formation of Mg(OH)₂ was apparent in MgO-bearing mortars, TGA results did not present visible M-S-H-like patterns or those of hydrotalcite-like compounds, suggesting marginal interaction with the cement’s products of hydration and siliceous aggregates. Nevertheless, additional analysis should be carried out on specimens without the incorporation of sand to further understand the hydration reactions.(iii)The dilution of cement with progressively higher replacement levels of MgO, and thus lower quantity of C-S-H, probably led to the formation of a less resistant microstructure (lower flexural and compressive strength). Additionally, greater water requirement was observed in mixes including MgO since the mortars in this study were produced with equivalent consistence and without the use of water-reducing admixtures. The resulting increased porosity, apart from causing a decline in mechanical performance, also led to increased carbonation and water absorption by capillary action.(iv)Mortars incorporating MgO are likely to present lower shrinkage strain in comparison with corresponding control specimens made solely with cement. Nevertheless, this enhanced performance (∼80% less shrinkage strain for mixes containing 15% and 20% MgO A), prompted by the expansion of Mg(OH)₂, heavily depends on the quantity of hydrated compounds and reactivity of MgO. From a structural point of view, the MgO-containing elements are likely to exhibit reduced shrinkage-induced cracking thereby improving the material’s overall structural and durability-related performances.

Data Availability

The data used to support the results of this study are included in the manuscript.

Conflicts of Interest

The authors declare that they have no conflicts of interest.

Acknowledgments

The authors gratefully acknowledge the support of the CERIS-ICIST Research Institute, IST-University of Lisbon, and FCT-Foundation for Science and Technology. This work was partly supported by the Andalusian Regional Government (Research Groups FQM-391) and MECD, Spain (FPU-13/04030; http://www.mecd.gob.es/educacion-mecd/). The authors also wish to thank Styromag, Grecian Magnesite S.A., Research Plan of the University of Córdoba (2016), and the staff at the Electron Microscopy and Elemental Analysis Units of the Central Research Support Service (SCAI) of University of Córdoba for their technical assistance and the Fine Chemistry Institute of the University of Córdoba for their support.

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Copyright © 2019 T. Gonçalves et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

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