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Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag
CEMCON-05036; No of Pages 8 Cement and Concrete Research xxx (2015) xxx–xxx
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Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag Bingqiang Han a,⁎, Peng Wang b, Changming Ke a, Wen Yan a, Yaowu Wei a, Nan Li a a b
The Key State Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan, Hubei, 430081, PR China Sinosteel Equipment & Engineering Co.,Ltd., Design Institute of Wuhan,Wuhan, Hubei, 430081, PR China
a r t i c l e
i n f o
Article history: Received 2 November 2014 Accepted 24 September 2015 Available online xxxx Keywords: Hydration products Pore size distribution Compressive strength Calorimetry SEM
a b s t r a c t Hydration behavior of the spinel containing high alumina cement prepared from high titania blast furnace slag via smelting reduction method is studied. Cooling condition has considerable effect on the phase compositions and hydration behavior of the prepared cements. Hydraulic CA, CA2, inert spinel and gehlenite are the main mineral phases of the naturally cooling cement. Glassy phase, CA and some spinel are the main phases of the splat cooling cement. Both of the prepared cements have controllable setting time, water requirements. Strength of splat cooling cement develops slowly than naturally cooling cement. The naturally cooling cement has satisfactory compressive strength, which is higher than splat cooling cement, but lower than commercial CA80 and Secar71. XRD and SEM observation confirms that CAH10 is the main hydrate of splat cooling cement. Metastable CAH10, C2AH8, are the main hydrates of naturally cooling cement, which will convert to stable C3AH6 with continuing hydration. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Calcium aluminate cement (CAC), particularly high alumina cements(HAC) is a type of hydraulic binder that has been in commercial existence since 1908 manufactured from bauxite and limestone via a fusion method in a reverberatory furnace [1–3]. CAC is classified into pure CAC and high alumina cement based on the alumina and impurities content, such as SiO2, Fe2O3. CAC with high alumina content and low impurities can be used in castable refractory subjects to more severe environments, such as Secar71 and Secar80(Kerneos, France) or CA70 and CA80(PR China). Today, such CACs are manufactured by either a fusion or sintering process. Typical starting raw materials are a source of Al2O3 and CaO and can include limestone, lime, bauxite and calcined alumina. The principal reactive phase in all CACs is monocalcium aluminate (CaAl2O4, CA); other phases present are calcium dialuminate (CaAl4O7,CA2), mayenite(Ca12Al14O33,C12A7). Iron oxides are detrimental and limit the use of these CACs to certain types of refractory applications except for ciment Fondu, which is commonly used in construction engineering. The hydration process, hydration products and the conversion of CAC have been studied by many researchers via different methods [4–8]. Furthermore, the hydration and applications in various environments of CAC with and without inert filler were also studied [9–13].
⁎ Corresponding author. E-mail address: [email protected] (B. Han).
CAC has unique properties: (1)high-temperature resistance and refractory performance;(2)rapid strength development; (3)resistance to chemically aggressive environments [2,11,12]. For lower purity, lower alumina content bauxite based CACs (HACs),such as CA50 and CA60, the phase composition is complicated by the presence of silica and iron oxide. These CACs contain additional phases of the CaO–Al2O3–SiO2 and CaO–Al2O3–Fe2O3 threecomponent systems. Nowadays, HAC is generally produced in rotary kiln. Generally, spinel is widely used in refractory castables together with CACs, especially in steel ladle lining due to its unique high temperature properties [12–21]. Spinel can absorb FeO, MnO, TiO2 and V2O5 and others in slag, making the slag more viscous and hindering the penetration of the slag. Thus, the castable containing spinel expresses excellent durability. Over the last few years, the use of spinel containing castable refractory has increased significantly. However, the high cost of eletrofused or sintered spinel limits their applications. Recently, new spinel-containing aluminate cements have been studied based on CaO–Al2O3–MgO ternary diagram via the reaction sintering of dolomite and alumina by A.H. De. Aza, Xiao Guoqing and other researchers [22–27]. A.H. De. Aza et al. indicated that a spinel-containing aluminate cement with 43 ± 5 wt.% CA, 15 ± 3 wt.% CA2 and 42 ± 2 wt.% MgAl2O4, showed a hydration and dehydration behavior similar to that of commercial calcium aluminate cements [24]. J. M. Auvray compared the hydration and applicative properties of calcium magnesium aluminate cement and calcium aluminate regular cement. Spinel was present as micro crystals embedded into matrix of calcium aluminates, which was favorable to improve the corrosion resistance by
http://dx.doi.org/10.1016/j.cemconres.2015.09.019 0008-8846/© 2015 Elsevier Ltd. All rights reserved.
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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B. Han et al. / Cement and Concrete Research xxx (2015) xxx–xxx
molten slags and metals. They also pointed out that the hydration mechanism of the calcium magnesium aluminate cement followed the same path as a pure calcium aluminate cement [27]. E. Dourdounis et al. has reported a new routine to prepare high alumina cement(HAC) from ferronickel electro-reduction furnace slag, limestone and low-grade diasporic bauxite via smelting reduction method [28]. The prepared HAC has satisfactory compressive strength relative to those of commercial HAC. During the process, FeOx and NiO will be reduced to iron and nickel and the metals can be recovered. Common blast furnace slag(BF slag) has been utilized as cement mixture, or fine aggregates in concrete. However, those blast furnace slags with high titania content(HTBF slag) in PR China and other countries are still reused ineffectively [29–33]. The main problem in the utilization of the slag is the low hydration activity. High titania content(N20 mass%) results in high content titanaugite and low silicates with potential hydration activity. So more than 60 million slag was piled up in PR China and more than 3 million tons of new additional slag is produced per year. The pileup of the BF slag not only pollutes the environment, but also wastes the titanium resource. The comprehensive utilization of the BF slag is becoming the focus. When the titanium and silicon are extracted from the HTBF slag, the residue is classified into CaO–Al2O3–MgO ternary system. The recovered alloy has potential applications in alloy steel production. The residue can be used as spinel containing CAC cement. Thus the HTBF slag can be effectively and completely reused without new second pollution. Based on the above train of thought, in recent years, Ke Changming, Han Bingqiang have tried the smelting reduction of HTBF slag by electric-arc furnace and studied the hydration performance of the residue [34–36]. In this paper, the hydration behavior of the residue after extracting titanium and silicon from HTBF slag were studied. 2. Experimental procedure The main raw materials were the residues extracted from HTBF slag. The HTBF slag and other reducing agents were treated in a 1.5 t arc furnace. After the reaction finishment, the formed melt was kept at 1600– 1700 °C for a few minutes. Then the melts were cooled by two methods. One is splat cooling by compressed air and the other is naturally cooling. For former, the melt was blowed into sphere by compressed air. For the latter, the melt was tapped in a ladle and cooled down up to room temperature. After a 24 h or 48 h cooling period, the solidified melt was crushed and the residue and alloy were separated for following analysis. Two different types of residue were used in this study named HAC1 and HAC2. The chemical compositions of HTBF slag and the prepared HACs analyzed by chemical analysis (Thermo Elemental IRIS Advantage Radial ICP-AES) are given in Table 1. By comparison, chemical compositions of commercial CA50, CA80 and Secar71 cement are also listed in Table 2. CA50 cement is a kind of calcium-based cement prepared by the lower grade bauxite and CA80 and Secar71 are high purity cement. The chemical compositions of the HACs fall into the range of the CA50 cement except higher MgO content. It can be seen that the as-prepared HACs are mainly composed of Al2O3, CaO, MgO, SiO2. The content of TiO2 in the prepared HACs decreases from N20 wt.% to b2 wt.%. Table 1 Chemical compositions of HTBF slag and the prepared HACs after extracting titanium and silicon (wt.%).
HTBF HAC1 HAC2
TiO2
SiO2
Al2O3
MnO
V2O5
MgO
CaO
Na2O
S
TFe
21.40 1.66 1.45
24.72 4.30 5.86
13.64 59.66 56.35
0.54 0.07 0.12
0.29 0.002 0.002
7.05 6.71 7.65
28.39 28.99 26.60
0.26 0.06 0.07
0.39 0.22 0.32
2.40 0.25 0.27
The mineralogical phases of the prepared HACs were detected by an X-ray diffractometer (Philips, X'pert Pro) with a Ni-filtered Cu Kα radiation (see Fig. 1). The machine settings were: scanning range = 5°–90°, and scanning rate = 1.5°/min. There are no evident differences in chemical compositions between the two HACs, however, the mineralogical phases are quite different. From Fig. 1, it can be seen that the HAC1 has much glassy phase and only a little CA and spinel can be detected. For HAC2, hydraulic CA and CA2, inert MgAl2O4 spinel and gehlenite are the main phases. For quantitative determination of the different phases in the prepared cement, XRD based Rietveld analyses were carried out and the method is briefly discussed below [37–41]. The weight fraction (Wα) of each phase α is related with its scale factor (Sα) which is defined as Eq. (1)[37,38]: Sα ¼
W α ρm K
ð1Þ
V 2α ρα μ m
where ρm and μm are the density and the linear X-ray absorption coefficient of the mixture. K is a constant which depend on the diffractometer operation conditions, Vα is the unit cell volume of α-phase. Once the crystal structure is known, the Vα is known. Finally, the weight fraction for each phase can be calculated by Eq. (2): Wα ¼
Sα ðZMV Þα n X Si ðZMV Þi
ð2Þ
i¼1
where Z is the number of chemical units per unit cell of α-phase, M is the molecular mass of the chemical formular for α-phase, and V is the unit cell volume for α-phase. The use of Eq. (2) eliminates the need to measure the instrument calibration constant K and the mass absorption coefficient μ. Here,it is assumed that all phases present in the sample are included in the analysis and there is not an amorphous phase and/or unaccounted crystalline phases. In this case, a known weight (Ws) of crystalline internal standard must be added to the sample and the relative fraction for each phase is calculated from Eq. (3). It is must be noted that the addition of Si standard does not affect the estimated weight percent of phases, Wα ¼
Ws Sα ðZMV Þα : Ss ðZMV Þs
ð3Þ
In this case, the glassy phases(Wg) is calculated by difference from Eq. (4): W g ¼ 1−W s −
n X
Wα:
ð4Þ
α¼1
Since the cement is produced from recycled raw materials with variable composition, the clinker produced results in a complex variable multiphase–multicomponent system. So it is difficult to take quantitative analysis on the cement, especially for the rapid cooling cement with a great amount of glassy content. Thus only HAC2 is analyzed by quantitative X-ray diffraction analysis and the calculated weight percent values of phases from Eq. (2) is 43 ± 5% CA,11 ± 3%, CA2 23 ± 2% spinel, 19 ± 2% gehlenite and about 3 ± 1% perovskite (CaTiO3). Relative weight percent corrected by calculating the glassy phase with Si as internal standard(Eq. (4)) is 41 ± 5% CA, 10 ± 3% CA2 22 ± 2% spinel, 18 ± 2% gehlenite and about 3 ± 1% perovskite (CaTiO3). The glassy phase of HAC2 is about 4% ± 2%. The prepared high alumina cement was ground to 4000 cm2/g by means of a 100 kg vibration mill and measured by Blaine method. Particles size distribution of the prepared HACs was analyzed with a Malvern Company's laser particle size analyzer(see Fig. 2). The setting time and hydration properties of prepared HACs were detected based on
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
B. Han et al. / Cement and Concrete Research xxx (2015) xxx–xxx
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Table 2 Typical chemical compositions of commercial CAC cement (wt.%).
CA50cement CA80 cement Secar71 cement
Al2O3
CaO
SiO2
MgO
TiO2
Fe2O3
Na2O + K2O
S
49–55 79–82 68.5–70.5
34–39 17–20 28.5–30.5
4–6 ≤0.5 0.2–0.6
b1 b0.2 b0.5
/ b0.1 b0.4
≤2.5 ≤0.5 0.1–0.3
b0.4 b0.7 b0.5
≤0.1 ≤0.1 b0.3(SO3)
GB201-2000. Paste compositions are: standard sand: 1350 g, cement: 450 g and water: 225 g, W/C is = 0.5. Bending strength and compressive strength tests were conducted at each age of the specimens. For the investigations of the hydration behavior of the prepared HACs, the hydration heat liberation was continuous recorded by a microcalorimeter (eight channels, TAM Air C80, SETARAM Company, France) in isothermal mode. During the measurements, the temperature was kept constant at value of 23.0 ± 0.3 °C. The fragments obtained after failure were crushed and immersed in acetone and alcohol to interrupt the hydration process of mineral phases for microstructure observation and pore analysis. Pore size distribution was evaluated using the mercury intrusion porosimetry technique with a Micromeritics Autopore mercury intrusion porosimeter on prismatic samples taken from the specimens, which automatically registers pressure, pore diameter, intrusion volume and pore surface area. Microstructure analysis was studied by a Nova400 Nano SEM (FEI Company, USA) coupled with an Oxford IE350 Penta FETX-3 energy dispersive spectroscopy (EDS) and a Denmark HKL Channel 5 EBSD was for semi-quantitative analysis. The samples were sectioned perpendicularly by diamond saw and the sample with corroded areas were cut and embedded in a low viscosity resin, then ground with diamond and B4C abrasive, then polished with diamond suspensions, and carbon-coated using standard ceramographic techniques. Second electron imaging (SEI) and back-scattered electron imaging (BSE) were used during the microstructure observation.
3. Results and discussion 3.1. Hydration properties of the spinel containing HAC Table 3 lists the setting time and water requirement for normal consistency of cement pastes of the commercial Secar71, CA50 and the as-prepared spinel containing HACs. From the results, it can be seen
Fig. 1. Mineralogical phases of the prepared HACs.
that the prepared cements have controllable setting time and water requirement. Fig. 3 shows the heat flow curves of the prepared HACs. HAC2 is characterized by a relatively shorter dormant period and an intensive main hydration peak with a maximum at about 11 h, which is similar to regular CAC. HAC1 system exhibits a relatively longer dormant period with a very low heat flow. The maximum heat liberation appears at about 18 h and 40 h, respectively. HAC1 hydrated slowly than HAC2, which is attributed to the difference of relative content of CA, CA2, glassy phase and specific surface. Microstructure and compositions of the glassy phase are also important. HAC1 shows many heaps, which is due to different grain sizes. Strength development strongly depends on the kinds and amount of the formed hydrates, which varied with the w/c ratio, curing temperature and aging time [42,43]. Furthermore, space also can affect the reaction between HACs with H2O and the conversion of CAH10 to more stable C3AH6. For HAC1, CAH10 is the main hydrate and the amount increases with the aging time. Some hydrogarnet (C3AH6) can be detected for 28 days paste. For HAC2, C2AH8, C3AH6 and gibbsite (AH3) are the main hydrates instead of CAH10 for all pastes. At the same time, unreacted CA and CA2 are present in all pastes. Fig. 4 gives the physical properties of the prepared HACs and three commercial cements. The bending strength and strength development of the prepared HACs increase with the aging time, which is similar to the three commercial cements. For the bending strength, it is observed that the HAC1 shows the early strength of 3.2 MPa in 1 day, compared to 14.5 MPa, 12.2 MPa, 7.7 MPa, 6.9 MPa of Secar71, CA80, HAC2 and CA50 cement, respectively. However, the 3 days and 7 days bending strength of the HAC1 and HAC2 varies a little. In 28 days, HAC2 develops strength of 8.8 MPa, higher than the 6.7 MPa of HAC1 and 8.1 MPa of CA50, but lower than the 15.7 MPa of Secar 71, 13.7 MPa of CA80. HAC1 develops more slowly than HAC2. However, HAC1 has a greater increase than HAC2 for 28 days curing. For compressive strength, the similar trend is also found in 28 day strengths of 32.8 MPa, 74.5 MPa, 60.9 MPa, 111.1 MPa and 115.1 MPa for the prepared HAC1 and HAC2, commercial CA50, CA80 and Secar71. From the above discussions, it can be concluded that the strength development of the prepared HAC2 is higher than commercial CA50 and HAC1, but lower than Secar71 and CA80.
Fig. 2. Particle size distribution of the prepared HACs.
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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B. Han et al. / Cement and Concrete Research xxx (2015) xxx–xxx
Table 3 Setting time and water requirement for normal consistency of cement paste. Types
Secar71cement CA50 cement CA80 cement HAC1 HAC2
Water requirement for normal consistency of cement paste (%)
Setting time Initial setting time
Final setting time
38.0 30.7 18 29 27
104 min 87 min 155 min 165 min 75 min
2 h21 min 3 h42 min 3 h50 min 4 h20 min 4 h41 min
Long-term strength development of the paste depends on the kinds and amount of hydration products, microstructure, curing temperature and admixture. HAC2 has little glassy phase; however, HAC1 has a great amount of glassy phase, which is belonged to CaO–Al2O3–MgO–SiO2 system, similar to granulated blast furnace slag formed under markedly non-equilibrium conditions. The hydraulic activities of HAC1 are not determined solely by bulk composition and glass content. Reactivity may be presumed to depend on glass composition and structure, crystalline size and morphology. Furthermore, the separation degree of crystal phase with glass maybe has considerable influence on the hydration. It must be noted that inert spinel and gehlenite perhaps have some influence on the hydration process, similar to inert filler. The metastable phases of CAH10 and C2AH8 are of low density but occupies a large volume due to high amount of combined water. It is helpful to occupy much of the space filled by water during early hydration and account for rapid early strength gain. During conversion the overall porosity of the system increases and the strength decreases. The slow strength development of HAC2 paste may be related to the conversion of C3AH6. The rapid strength increase of HAC1 is due to continuing hydration and the formation of CAH10. Cumulative and differential pore volume with pore diameter of hydrated pastes for 3 days and 28 days are shown in Fig. 5. Properties of pores in hydrated pastes are listed in Table 4. The pore size distribution curve of HAC2 paste has a single peak for 3 days hydration, but has double peak for 28 days. However, HAC1 paste has double peak pore for 3 days and a single peak for 28 days hydration. It also can be seen from Fig. 5 that the general pore size distribution shifts to the left with an increasing hydration period, i.e. the pores become smaller pores, which is similar to common Portland cement paste. Median pore diameter and average pore diameter of HAC1 are far smaller than HAC2.
Fig. 4. Bending strength and compressive strength of mortar paste with curing time.
3.2. Phase analysis of hydrates
Fig. 3. Heat flow curves of two HACs.
1 day, 7 days and 28 days pastes were also analyzed with XRD and the results are shown in Fig. 6. XRD observation reveals that two inert components, spinel and gehlenite are present in all pastes. For pastes of HAC1, CAH10 with intensive peaks between 6 and 8° (2θ), 10 and 15° are present and increase with the curing time. C2AH8, C3AH6 and AH3 are detected in the paste with the proceeding of hydration. It must be noted that a great amount of amorphous phases are detected in all pastes. For pastes of HAC2, XRD observation shows increasing CAH10, C3AH6 and decreasing CA. CAH10, C2AH8, C3AH6, AH3 and unreacted CA, CA2, which coexist in HAC2 paste after 1 day hydration. For 7 days paste of HAC2, CAH 10 , C 2 AH8 C3 AH6 and AH3 were the dominant hydrates. For 28 days paste of HAC2, more stable C 3AH6 and AH3 are the dominant hydrates in replacement of metastable CAH10 and C2AH8. The conversion is temperature and time dependent. With the conversion of unstable CAH 10 and C 2 AH8 to stable C3AH6, some water was released in hardened matrix, causing an increase in porosity and thereby a reduction in strength. It was also mentioned by other researchers that the bonds of C3 AH6 and AH3 are weaker than those of CAH10 and C2AH8, even if at equal porosity [1,44–46].
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
B. Han et al. / Cement and Concrete Research xxx (2015) xxx–xxx
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Fig. 6. XRD analysis of the hydrated pastes for different aging time. Fig. 5. Cumulative and differential pore volume vs. pore diameter of hydrated paste for 3 days and 28 days.
3.3. Microstructure observation The hydration process and microstructure of CAC has been discussed by many researchers [1–11,44–48]. Naturally, the reactions are highly dependent upon composition, temperature, moisture state, admixture and W/C ratio. When the temperature is less than 10 °C, the formation of CAH10 is predominant. While between 10 °C and 27 °C, CAH10 and C2AH8 will be both formed. At higher temperatures CAH10 is no longer formed and the stable phase of C3AH6 occurs early in the process of hydration. It is claimed that the formation of C3AH6 is always preceded by some formation of C2AH8, The direct formation of C3AH6 from CA can take place after the nucleation of some C3AH6. Once nucleation of C3AH6 has occurred, further hydration will lead to direct formation of stable hydrates since there is no nucleation barrier to overcome.
Table 4 Properties of pore in hydrated paste. Sample HAC1 HAC2
3d 28d 3d 28d
Porosity/%
Total intrusion volume (mg/L)
Median pore diameter (nm)
Average pore diameter (nm)
16.2 21.6 17.2 18.5
0.1003 0.1224 0.1022 0.1123
12.5 24.1 426.8 56.9
8.4 18.9 121.9 26.4
The microstructures of 1 day, 7 days and 28 days pastes were also studied with second electron imaging (SEI). The shape and texture of samples can be deduced from Fig. 7. For 1 day paste of HAC1, fine meshed CAH10 was observed (Fig. 7a). Clavate AH3 gel, fine meshed CAH10 with anhydrous, i.e., unreacted clinker grains formed a porous structure in HAC2 paste (Fig. 7b). Most of inert spinel and gehlenite have been masked and can't be seen in SEM observation. By 7 days hydration, layered crystal C2AH8 and fine meshed CAH10 crystal are observed in HAC1 paste (Fig. 7c). More CAH10, AH3 gel are present in HAC2 paste (Fig. 7d). By 28 days, the hydration reaction had progressed considerably, with a decline in the amount of anhydrous cement grains. Most of metastable CAH 10 and C 2 AH8 have converted to very small C3AH6 in HAC2 paste, which is the only stable ternary phase in CaO–Al2O3–H2O system at ordinary temperature. So C3AH6 combined with some CAH 10 and gibbsite are the main hydrates (Fig. 7f). However, CAH10 and C 2AH8 are still the main hydrate in HAC1 paste, unreacted CA grains are present (Fig. 7e). Furthermore, it can be seen that the higher the curing time, the more compact paste was formed. The hydration process of the prepared HACs is complicated. Further experiments investigated by other methods at exactly the same hydration conditions have to be done, and to get more
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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Fig. 7. SEM second electron image of prepared HACs hydrated for 1 days, 7 days and 28 days with w/c = 0.5 at room temperature, showing the different morphology of hydration products. a: HAC1-1d. Clinker grains, frozen water at the background, fine meshed CAH10. c:HAC1-7d. Layered crystal C2AH8 and fine meshed CAH10 and unreacted clinker grains. e: HAC1-28d. Wellrecrystallized CAH10 and layered crystal C2AH8, unreacted clinker grains covered with some kind of amorphous layers. b: HAC2-1d. Clavate AH3 gel, fine meshed CAH10 with unreacted clinker grains. d: HAC2-7d; Clavate AH3 gel, fine meshed CAH10. f: HAC2-28d. Intermixed C3AH6 and AH3 gel constitute the hydrated phases, non reacted gehlenite (C2AS) is present. The converted microstructure has significant porosity.
detailed and comparable results about the differences in HACs hydration reactions. The effects of the composition and microstructure of glass phase on the hydration mechanism still need to be further studied.
4. Conclusions Hydration behavior of spinel containing high alumina cement are studied, which is prepared in an electric arc furnace with high titania
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
B. Han et al. / Cement and Concrete Research xxx (2015) xxx–xxx
BF slag as raw material after extracted Fe–Si–Ti alloy . Physical properties of the prepared cement and the evolution of hydrates are discussed. The following conclusions can be drawn. Cooling condition has considerable effects on the mineral phases of the prepared HAC cement. Hydraulic CA and CA2, inert spinel and gehlenite are the main mineral phases of HAC2. However, glassy phase is the main component of the splat cooling HAC1, which is classified into CaO–Al2O3–MgO–SiO2 system, similar to granulated blast furnace slag formed under markedly non-equilibrium conditions. The hydraulic activities of HAC1 are not determined solely by bulk composition and glass content. Reactivity may be presumed to depend on glass composition and structure, crystallite size and morphology, degree of crystal phase separation with glass. The hydration reactivity and the kinds and amount of hydration products are quite different. The prepared cement has controllable setting time and water requirement. The hydration of the HACs is similar to commercial aluminate cement. The strength development of the prepared HAC2 is between CA50 cement and Secar71, however, strength of HAC1 develops more slowly than HAC2. CAH10 is the main hydrate for all pastes of HAC1 at different aging time. CAH10, and C2AH8 are the dominant phases in HAC2 paste except for spinel and gehlenite and most of them will convert to more stable C3AH6 after 28 days curing. With the conversion of unstable CAH10 and C2AH8 to stable C3AH6, some water was released in hardened matrix, causing an increase in porosity and thereby a reduction in strength. Acknowledgments The authors wish to thank the Education Bureau of Hubei Province(D20141103) and Ministry of Science and Technology, PR China for funding the research under 973 plan of PR China (2014CB660802). References [1] H.F.W. Taylor, Cement Chemistry, Academic Press, London, 1990 316–361. [2] Lea's Chemistry of Cement and Concrete, in: Peter Hewlett (Ed.) fourth ed.Elsevier, London 2003, pp. 715–782. [3] D. Sorrentino, F. Sorrentino, C.M. George, Mechanisms of hydration of calcium aluminate cements, in: J.P. Scalny (Ed.), Materials science of concrete, American Ceramics Society, vol. IV 1995, pp. 41–90. [4] S. Ng, T. HjellastrÖm, Quantitative and qualitative analysis of CAC and its hydrate products, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 17–30. [5] M. Schmidt, H. Pöllmann, Hydration behaviour of CAC at various temperatures by isoperibolic heat flow calorimetry using acetates, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2008, pp. 93–108. [6] C. Gosselin, K.L. Scrivener, Microstructure development of calcium aluminate cements accelerated by lithium sulphate, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2008, pp. 109–122. [7] F. Goetz-Neunhoeffer, S.R. Klaus, J. Neubauer, Kinetics of CA and CA2 dissolution determined by QXRD and corresponding enthalpies of reactions, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 54–64. [8] S.R. Klaus, J. Neubauer, F. Goetz-Neunhoeffer, A. Buhr, D. Schmidtmeier, Application of heat flow calculation to synthetic calcium aluminate cement mixes, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 65–74. [9] B. Lothenbach, Thermodynamic modelling of effects of time and silica on the conversion process, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 153–164. [10] J. Bizzozero, K.L. Scrivener, Hydration of calcium aluminate cement based systems with calcium sulfate and limestone, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 189–196. [11] S. Zhang, W.E. Lee, Spinel-containing refractories, in: C.A. Schacht (Ed.), Refractories Handbook, Marcel Dekker, Monticello 2004, pp. 81–139. [12] A. Buhr, D. Schmidtmeier, G. Wams, S. Kuiper, S. Klaus, Testing of calcium aluminate cement bonded castables and influence of curing conditions on the strength development, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 437–449. [13] J. Kasper, O. Krause, Mixing optimization of an alumina based LC-castable by applying variable power inputs, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 450–456. [14] M.A.L. Braulio, M. Rigaud, A. Buhr, C. Parr, V.C. Pandolfelli, Spinel-containing alumina-based refractory castables, Ceram. Int. 37 (2011) 1705–1724.
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Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag Bingqiang Han a,⁎, Peng Wang b, Changming Ke a, Wen Yan a, Yaowu Wei a, Nan Li a a b
The Key State Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan, Hubei, 430081, PR China Sinosteel Equipment & Engineering Co.,Ltd., Design Institute of Wuhan,Wuhan, Hubei, 430081, PR China
a r t i c l e
i n f o
Article history: Received 2 November 2014 Accepted 24 September 2015 Available online xxxx Keywords: Hydration products Pore size distribution Compressive strength Calorimetry SEM
a b s t r a c t Hydration behavior of the spinel containing high alumina cement prepared from high titania blast furnace slag via smelting reduction method is studied. Cooling condition has considerable effect on the phase compositions and hydration behavior of the prepared cements. Hydraulic CA, CA2, inert spinel and gehlenite are the main mineral phases of the naturally cooling cement. Glassy phase, CA and some spinel are the main phases of the splat cooling cement. Both of the prepared cements have controllable setting time, water requirements. Strength of splat cooling cement develops slowly than naturally cooling cement. The naturally cooling cement has satisfactory compressive strength, which is higher than splat cooling cement, but lower than commercial CA80 and Secar71. XRD and SEM observation confirms that CAH10 is the main hydrate of splat cooling cement. Metastable CAH10, C2AH8, are the main hydrates of naturally cooling cement, which will convert to stable C3AH6 with continuing hydration. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Calcium aluminate cement (CAC), particularly high alumina cements(HAC) is a type of hydraulic binder that has been in commercial existence since 1908 manufactured from bauxite and limestone via a fusion method in a reverberatory furnace [1–3]. CAC is classified into pure CAC and high alumina cement based on the alumina and impurities content, such as SiO2, Fe2O3. CAC with high alumina content and low impurities can be used in castable refractory subjects to more severe environments, such as Secar71 and Secar80(Kerneos, France) or CA70 and CA80(PR China). Today, such CACs are manufactured by either a fusion or sintering process. Typical starting raw materials are a source of Al2O3 and CaO and can include limestone, lime, bauxite and calcined alumina. The principal reactive phase in all CACs is monocalcium aluminate (CaAl2O4, CA); other phases present are calcium dialuminate (CaAl4O7,CA2), mayenite(Ca12Al14O33,C12A7). Iron oxides are detrimental and limit the use of these CACs to certain types of refractory applications except for ciment Fondu, which is commonly used in construction engineering. The hydration process, hydration products and the conversion of CAC have been studied by many researchers via different methods [4–8]. Furthermore, the hydration and applications in various environments of CAC with and without inert filler were also studied [9–13].
⁎ Corresponding author. E-mail address: [email protected] (B. Han).
CAC has unique properties: (1)high-temperature resistance and refractory performance;(2)rapid strength development; (3)resistance to chemically aggressive environments [2,11,12]. For lower purity, lower alumina content bauxite based CACs (HACs),such as CA50 and CA60, the phase composition is complicated by the presence of silica and iron oxide. These CACs contain additional phases of the CaO–Al2O3–SiO2 and CaO–Al2O3–Fe2O3 threecomponent systems. Nowadays, HAC is generally produced in rotary kiln. Generally, spinel is widely used in refractory castables together with CACs, especially in steel ladle lining due to its unique high temperature properties [12–21]. Spinel can absorb FeO, MnO, TiO2 and V2O5 and others in slag, making the slag more viscous and hindering the penetration of the slag. Thus, the castable containing spinel expresses excellent durability. Over the last few years, the use of spinel containing castable refractory has increased significantly. However, the high cost of eletrofused or sintered spinel limits their applications. Recently, new spinel-containing aluminate cements have been studied based on CaO–Al2O3–MgO ternary diagram via the reaction sintering of dolomite and alumina by A.H. De. Aza, Xiao Guoqing and other researchers [22–27]. A.H. De. Aza et al. indicated that a spinel-containing aluminate cement with 43 ± 5 wt.% CA, 15 ± 3 wt.% CA2 and 42 ± 2 wt.% MgAl2O4, showed a hydration and dehydration behavior similar to that of commercial calcium aluminate cements [24]. J. M. Auvray compared the hydration and applicative properties of calcium magnesium aluminate cement and calcium aluminate regular cement. Spinel was present as micro crystals embedded into matrix of calcium aluminates, which was favorable to improve the corrosion resistance by
http://dx.doi.org/10.1016/j.cemconres.2015.09.019 0008-8846/© 2015 Elsevier Ltd. All rights reserved.
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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molten slags and metals. They also pointed out that the hydration mechanism of the calcium magnesium aluminate cement followed the same path as a pure calcium aluminate cement [27]. E. Dourdounis et al. has reported a new routine to prepare high alumina cement(HAC) from ferronickel electro-reduction furnace slag, limestone and low-grade diasporic bauxite via smelting reduction method [28]. The prepared HAC has satisfactory compressive strength relative to those of commercial HAC. During the process, FeOx and NiO will be reduced to iron and nickel and the metals can be recovered. Common blast furnace slag(BF slag) has been utilized as cement mixture, or fine aggregates in concrete. However, those blast furnace slags with high titania content(HTBF slag) in PR China and other countries are still reused ineffectively [29–33]. The main problem in the utilization of the slag is the low hydration activity. High titania content(N20 mass%) results in high content titanaugite and low silicates with potential hydration activity. So more than 60 million slag was piled up in PR China and more than 3 million tons of new additional slag is produced per year. The pileup of the BF slag not only pollutes the environment, but also wastes the titanium resource. The comprehensive utilization of the BF slag is becoming the focus. When the titanium and silicon are extracted from the HTBF slag, the residue is classified into CaO–Al2O3–MgO ternary system. The recovered alloy has potential applications in alloy steel production. The residue can be used as spinel containing CAC cement. Thus the HTBF slag can be effectively and completely reused without new second pollution. Based on the above train of thought, in recent years, Ke Changming, Han Bingqiang have tried the smelting reduction of HTBF slag by electric-arc furnace and studied the hydration performance of the residue [34–36]. In this paper, the hydration behavior of the residue after extracting titanium and silicon from HTBF slag were studied. 2. Experimental procedure The main raw materials were the residues extracted from HTBF slag. The HTBF slag and other reducing agents were treated in a 1.5 t arc furnace. After the reaction finishment, the formed melt was kept at 1600– 1700 °C for a few minutes. Then the melts were cooled by two methods. One is splat cooling by compressed air and the other is naturally cooling. For former, the melt was blowed into sphere by compressed air. For the latter, the melt was tapped in a ladle and cooled down up to room temperature. After a 24 h or 48 h cooling period, the solidified melt was crushed and the residue and alloy were separated for following analysis. Two different types of residue were used in this study named HAC1 and HAC2. The chemical compositions of HTBF slag and the prepared HACs analyzed by chemical analysis (Thermo Elemental IRIS Advantage Radial ICP-AES) are given in Table 1. By comparison, chemical compositions of commercial CA50, CA80 and Secar71 cement are also listed in Table 2. CA50 cement is a kind of calcium-based cement prepared by the lower grade bauxite and CA80 and Secar71 are high purity cement. The chemical compositions of the HACs fall into the range of the CA50 cement except higher MgO content. It can be seen that the as-prepared HACs are mainly composed of Al2O3, CaO, MgO, SiO2. The content of TiO2 in the prepared HACs decreases from N20 wt.% to b2 wt.%. Table 1 Chemical compositions of HTBF slag and the prepared HACs after extracting titanium and silicon (wt.%).
HTBF HAC1 HAC2
TiO2
SiO2
Al2O3
MnO
V2O5
MgO
CaO
Na2O
S
TFe
21.40 1.66 1.45
24.72 4.30 5.86
13.64 59.66 56.35
0.54 0.07 0.12
0.29 0.002 0.002
7.05 6.71 7.65
28.39 28.99 26.60
0.26 0.06 0.07
0.39 0.22 0.32
2.40 0.25 0.27
The mineralogical phases of the prepared HACs were detected by an X-ray diffractometer (Philips, X'pert Pro) with a Ni-filtered Cu Kα radiation (see Fig. 1). The machine settings were: scanning range = 5°–90°, and scanning rate = 1.5°/min. There are no evident differences in chemical compositions between the two HACs, however, the mineralogical phases are quite different. From Fig. 1, it can be seen that the HAC1 has much glassy phase and only a little CA and spinel can be detected. For HAC2, hydraulic CA and CA2, inert MgAl2O4 spinel and gehlenite are the main phases. For quantitative determination of the different phases in the prepared cement, XRD based Rietveld analyses were carried out and the method is briefly discussed below [37–41]. The weight fraction (Wα) of each phase α is related with its scale factor (Sα) which is defined as Eq. (1)[37,38]: Sα ¼
W α ρm K
ð1Þ
V 2α ρα μ m
where ρm and μm are the density and the linear X-ray absorption coefficient of the mixture. K is a constant which depend on the diffractometer operation conditions, Vα is the unit cell volume of α-phase. Once the crystal structure is known, the Vα is known. Finally, the weight fraction for each phase can be calculated by Eq. (2): Wα ¼
Sα ðZMV Þα n X Si ðZMV Þi
ð2Þ
i¼1
where Z is the number of chemical units per unit cell of α-phase, M is the molecular mass of the chemical formular for α-phase, and V is the unit cell volume for α-phase. The use of Eq. (2) eliminates the need to measure the instrument calibration constant K and the mass absorption coefficient μ. Here,it is assumed that all phases present in the sample are included in the analysis and there is not an amorphous phase and/or unaccounted crystalline phases. In this case, a known weight (Ws) of crystalline internal standard must be added to the sample and the relative fraction for each phase is calculated from Eq. (3). It is must be noted that the addition of Si standard does not affect the estimated weight percent of phases, Wα ¼
Ws Sα ðZMV Þα : Ss ðZMV Þs
ð3Þ
In this case, the glassy phases(Wg) is calculated by difference from Eq. (4): W g ¼ 1−W s −
n X
Wα:
ð4Þ
α¼1
Since the cement is produced from recycled raw materials with variable composition, the clinker produced results in a complex variable multiphase–multicomponent system. So it is difficult to take quantitative analysis on the cement, especially for the rapid cooling cement with a great amount of glassy content. Thus only HAC2 is analyzed by quantitative X-ray diffraction analysis and the calculated weight percent values of phases from Eq. (2) is 43 ± 5% CA,11 ± 3%, CA2 23 ± 2% spinel, 19 ± 2% gehlenite and about 3 ± 1% perovskite (CaTiO3). Relative weight percent corrected by calculating the glassy phase with Si as internal standard(Eq. (4)) is 41 ± 5% CA, 10 ± 3% CA2 22 ± 2% spinel, 18 ± 2% gehlenite and about 3 ± 1% perovskite (CaTiO3). The glassy phase of HAC2 is about 4% ± 2%. The prepared high alumina cement was ground to 4000 cm2/g by means of a 100 kg vibration mill and measured by Blaine method. Particles size distribution of the prepared HACs was analyzed with a Malvern Company's laser particle size analyzer(see Fig. 2). The setting time and hydration properties of prepared HACs were detected based on
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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Table 2 Typical chemical compositions of commercial CAC cement (wt.%).
CA50cement CA80 cement Secar71 cement
Al2O3
CaO
SiO2
MgO
TiO2
Fe2O3
Na2O + K2O
S
49–55 79–82 68.5–70.5
34–39 17–20 28.5–30.5
4–6 ≤0.5 0.2–0.6
b1 b0.2 b0.5
/ b0.1 b0.4
≤2.5 ≤0.5 0.1–0.3
b0.4 b0.7 b0.5
≤0.1 ≤0.1 b0.3(SO3)
GB201-2000. Paste compositions are: standard sand: 1350 g, cement: 450 g and water: 225 g, W/C is = 0.5. Bending strength and compressive strength tests were conducted at each age of the specimens. For the investigations of the hydration behavior of the prepared HACs, the hydration heat liberation was continuous recorded by a microcalorimeter (eight channels, TAM Air C80, SETARAM Company, France) in isothermal mode. During the measurements, the temperature was kept constant at value of 23.0 ± 0.3 °C. The fragments obtained after failure were crushed and immersed in acetone and alcohol to interrupt the hydration process of mineral phases for microstructure observation and pore analysis. Pore size distribution was evaluated using the mercury intrusion porosimetry technique with a Micromeritics Autopore mercury intrusion porosimeter on prismatic samples taken from the specimens, which automatically registers pressure, pore diameter, intrusion volume and pore surface area. Microstructure analysis was studied by a Nova400 Nano SEM (FEI Company, USA) coupled with an Oxford IE350 Penta FETX-3 energy dispersive spectroscopy (EDS) and a Denmark HKL Channel 5 EBSD was for semi-quantitative analysis. The samples were sectioned perpendicularly by diamond saw and the sample with corroded areas were cut and embedded in a low viscosity resin, then ground with diamond and B4C abrasive, then polished with diamond suspensions, and carbon-coated using standard ceramographic techniques. Second electron imaging (SEI) and back-scattered electron imaging (BSE) were used during the microstructure observation.
3. Results and discussion 3.1. Hydration properties of the spinel containing HAC Table 3 lists the setting time and water requirement for normal consistency of cement pastes of the commercial Secar71, CA50 and the as-prepared spinel containing HACs. From the results, it can be seen
Fig. 1. Mineralogical phases of the prepared HACs.
that the prepared cements have controllable setting time and water requirement. Fig. 3 shows the heat flow curves of the prepared HACs. HAC2 is characterized by a relatively shorter dormant period and an intensive main hydration peak with a maximum at about 11 h, which is similar to regular CAC. HAC1 system exhibits a relatively longer dormant period with a very low heat flow. The maximum heat liberation appears at about 18 h and 40 h, respectively. HAC1 hydrated slowly than HAC2, which is attributed to the difference of relative content of CA, CA2, glassy phase and specific surface. Microstructure and compositions of the glassy phase are also important. HAC1 shows many heaps, which is due to different grain sizes. Strength development strongly depends on the kinds and amount of the formed hydrates, which varied with the w/c ratio, curing temperature and aging time [42,43]. Furthermore, space also can affect the reaction between HACs with H2O and the conversion of CAH10 to more stable C3AH6. For HAC1, CAH10 is the main hydrate and the amount increases with the aging time. Some hydrogarnet (C3AH6) can be detected for 28 days paste. For HAC2, C2AH8, C3AH6 and gibbsite (AH3) are the main hydrates instead of CAH10 for all pastes. At the same time, unreacted CA and CA2 are present in all pastes. Fig. 4 gives the physical properties of the prepared HACs and three commercial cements. The bending strength and strength development of the prepared HACs increase with the aging time, which is similar to the three commercial cements. For the bending strength, it is observed that the HAC1 shows the early strength of 3.2 MPa in 1 day, compared to 14.5 MPa, 12.2 MPa, 7.7 MPa, 6.9 MPa of Secar71, CA80, HAC2 and CA50 cement, respectively. However, the 3 days and 7 days bending strength of the HAC1 and HAC2 varies a little. In 28 days, HAC2 develops strength of 8.8 MPa, higher than the 6.7 MPa of HAC1 and 8.1 MPa of CA50, but lower than the 15.7 MPa of Secar 71, 13.7 MPa of CA80. HAC1 develops more slowly than HAC2. However, HAC1 has a greater increase than HAC2 for 28 days curing. For compressive strength, the similar trend is also found in 28 day strengths of 32.8 MPa, 74.5 MPa, 60.9 MPa, 111.1 MPa and 115.1 MPa for the prepared HAC1 and HAC2, commercial CA50, CA80 and Secar71. From the above discussions, it can be concluded that the strength development of the prepared HAC2 is higher than commercial CA50 and HAC1, but lower than Secar71 and CA80.
Fig. 2. Particle size distribution of the prepared HACs.
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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Table 3 Setting time and water requirement for normal consistency of cement paste. Types
Secar71cement CA50 cement CA80 cement HAC1 HAC2
Water requirement for normal consistency of cement paste (%)
Setting time Initial setting time
Final setting time
38.0 30.7 18 29 27
104 min 87 min 155 min 165 min 75 min
2 h21 min 3 h42 min 3 h50 min 4 h20 min 4 h41 min
Long-term strength development of the paste depends on the kinds and amount of hydration products, microstructure, curing temperature and admixture. HAC2 has little glassy phase; however, HAC1 has a great amount of glassy phase, which is belonged to CaO–Al2O3–MgO–SiO2 system, similar to granulated blast furnace slag formed under markedly non-equilibrium conditions. The hydraulic activities of HAC1 are not determined solely by bulk composition and glass content. Reactivity may be presumed to depend on glass composition and structure, crystalline size and morphology. Furthermore, the separation degree of crystal phase with glass maybe has considerable influence on the hydration. It must be noted that inert spinel and gehlenite perhaps have some influence on the hydration process, similar to inert filler. The metastable phases of CAH10 and C2AH8 are of low density but occupies a large volume due to high amount of combined water. It is helpful to occupy much of the space filled by water during early hydration and account for rapid early strength gain. During conversion the overall porosity of the system increases and the strength decreases. The slow strength development of HAC2 paste may be related to the conversion of C3AH6. The rapid strength increase of HAC1 is due to continuing hydration and the formation of CAH10. Cumulative and differential pore volume with pore diameter of hydrated pastes for 3 days and 28 days are shown in Fig. 5. Properties of pores in hydrated pastes are listed in Table 4. The pore size distribution curve of HAC2 paste has a single peak for 3 days hydration, but has double peak for 28 days. However, HAC1 paste has double peak pore for 3 days and a single peak for 28 days hydration. It also can be seen from Fig. 5 that the general pore size distribution shifts to the left with an increasing hydration period, i.e. the pores become smaller pores, which is similar to common Portland cement paste. Median pore diameter and average pore diameter of HAC1 are far smaller than HAC2.
Fig. 4. Bending strength and compressive strength of mortar paste with curing time.
3.2. Phase analysis of hydrates
Fig. 3. Heat flow curves of two HACs.
1 day, 7 days and 28 days pastes were also analyzed with XRD and the results are shown in Fig. 6. XRD observation reveals that two inert components, spinel and gehlenite are present in all pastes. For pastes of HAC1, CAH10 with intensive peaks between 6 and 8° (2θ), 10 and 15° are present and increase with the curing time. C2AH8, C3AH6 and AH3 are detected in the paste with the proceeding of hydration. It must be noted that a great amount of amorphous phases are detected in all pastes. For pastes of HAC2, XRD observation shows increasing CAH10, C3AH6 and decreasing CA. CAH10, C2AH8, C3AH6, AH3 and unreacted CA, CA2, which coexist in HAC2 paste after 1 day hydration. For 7 days paste of HAC2, CAH 10 , C 2 AH8 C3 AH6 and AH3 were the dominant hydrates. For 28 days paste of HAC2, more stable C 3AH6 and AH3 are the dominant hydrates in replacement of metastable CAH10 and C2AH8. The conversion is temperature and time dependent. With the conversion of unstable CAH 10 and C 2 AH8 to stable C3AH6, some water was released in hardened matrix, causing an increase in porosity and thereby a reduction in strength. It was also mentioned by other researchers that the bonds of C3 AH6 and AH3 are weaker than those of CAH10 and C2AH8, even if at equal porosity [1,44–46].
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
B. Han et al. / Cement and Concrete Research xxx (2015) xxx–xxx
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Fig. 6. XRD analysis of the hydrated pastes for different aging time. Fig. 5. Cumulative and differential pore volume vs. pore diameter of hydrated paste for 3 days and 28 days.
3.3. Microstructure observation The hydration process and microstructure of CAC has been discussed by many researchers [1–11,44–48]. Naturally, the reactions are highly dependent upon composition, temperature, moisture state, admixture and W/C ratio. When the temperature is less than 10 °C, the formation of CAH10 is predominant. While between 10 °C and 27 °C, CAH10 and C2AH8 will be both formed. At higher temperatures CAH10 is no longer formed and the stable phase of C3AH6 occurs early in the process of hydration. It is claimed that the formation of C3AH6 is always preceded by some formation of C2AH8, The direct formation of C3AH6 from CA can take place after the nucleation of some C3AH6. Once nucleation of C3AH6 has occurred, further hydration will lead to direct formation of stable hydrates since there is no nucleation barrier to overcome.
Table 4 Properties of pore in hydrated paste. Sample HAC1 HAC2
3d 28d 3d 28d
Porosity/%
Total intrusion volume (mg/L)
Median pore diameter (nm)
Average pore diameter (nm)
16.2 21.6 17.2 18.5
0.1003 0.1224 0.1022 0.1123
12.5 24.1 426.8 56.9
8.4 18.9 121.9 26.4
The microstructures of 1 day, 7 days and 28 days pastes were also studied with second electron imaging (SEI). The shape and texture of samples can be deduced from Fig. 7. For 1 day paste of HAC1, fine meshed CAH10 was observed (Fig. 7a). Clavate AH3 gel, fine meshed CAH10 with anhydrous, i.e., unreacted clinker grains formed a porous structure in HAC2 paste (Fig. 7b). Most of inert spinel and gehlenite have been masked and can't be seen in SEM observation. By 7 days hydration, layered crystal C2AH8 and fine meshed CAH10 crystal are observed in HAC1 paste (Fig. 7c). More CAH10, AH3 gel are present in HAC2 paste (Fig. 7d). By 28 days, the hydration reaction had progressed considerably, with a decline in the amount of anhydrous cement grains. Most of metastable CAH 10 and C 2 AH8 have converted to very small C3AH6 in HAC2 paste, which is the only stable ternary phase in CaO–Al2O3–H2O system at ordinary temperature. So C3AH6 combined with some CAH 10 and gibbsite are the main hydrates (Fig. 7f). However, CAH10 and C 2AH8 are still the main hydrate in HAC1 paste, unreacted CA grains are present (Fig. 7e). Furthermore, it can be seen that the higher the curing time, the more compact paste was formed. The hydration process of the prepared HACs is complicated. Further experiments investigated by other methods at exactly the same hydration conditions have to be done, and to get more
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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Fig. 7. SEM second electron image of prepared HACs hydrated for 1 days, 7 days and 28 days with w/c = 0.5 at room temperature, showing the different morphology of hydration products. a: HAC1-1d. Clinker grains, frozen water at the background, fine meshed CAH10. c:HAC1-7d. Layered crystal C2AH8 and fine meshed CAH10 and unreacted clinker grains. e: HAC1-28d. Wellrecrystallized CAH10 and layered crystal C2AH8, unreacted clinker grains covered with some kind of amorphous layers. b: HAC2-1d. Clavate AH3 gel, fine meshed CAH10 with unreacted clinker grains. d: HAC2-7d; Clavate AH3 gel, fine meshed CAH10. f: HAC2-28d. Intermixed C3AH6 and AH3 gel constitute the hydrated phases, non reacted gehlenite (C2AS) is present. The converted microstructure has significant porosity.
detailed and comparable results about the differences in HACs hydration reactions. The effects of the composition and microstructure of glass phase on the hydration mechanism still need to be further studied.
4. Conclusions Hydration behavior of spinel containing high alumina cement are studied, which is prepared in an electric arc furnace with high titania
Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019
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BF slag as raw material after extracted Fe–Si–Ti alloy . Physical properties of the prepared cement and the evolution of hydrates are discussed. The following conclusions can be drawn. Cooling condition has considerable effects on the mineral phases of the prepared HAC cement. Hydraulic CA and CA2, inert spinel and gehlenite are the main mineral phases of HAC2. However, glassy phase is the main component of the splat cooling HAC1, which is classified into CaO–Al2O3–MgO–SiO2 system, similar to granulated blast furnace slag formed under markedly non-equilibrium conditions. The hydraulic activities of HAC1 are not determined solely by bulk composition and glass content. Reactivity may be presumed to depend on glass composition and structure, crystallite size and morphology, degree of crystal phase separation with glass. The hydration reactivity and the kinds and amount of hydration products are quite different. The prepared cement has controllable setting time and water requirement. The hydration of the HACs is similar to commercial aluminate cement. The strength development of the prepared HAC2 is between CA50 cement and Secar71, however, strength of HAC1 develops more slowly than HAC2. CAH10 is the main hydrate for all pastes of HAC1 at different aging time. CAH10, and C2AH8 are the dominant phases in HAC2 paste except for spinel and gehlenite and most of them will convert to more stable C3AH6 after 28 days curing. With the conversion of unstable CAH10 and C2AH8 to stable C3AH6, some water was released in hardened matrix, causing an increase in porosity and thereby a reduction in strength. Acknowledgments The authors wish to thank the Education Bureau of Hubei Province(D20141103) and Ministry of Science and Technology, PR China for funding the research under 973 plan of PR China (2014CB660802). References [1] H.F.W. Taylor, Cement Chemistry, Academic Press, London, 1990 316–361. [2] Lea's Chemistry of Cement and Concrete, in: Peter Hewlett (Ed.) fourth ed.Elsevier, London 2003, pp. 715–782. [3] D. Sorrentino, F. Sorrentino, C.M. George, Mechanisms of hydration of calcium aluminate cements, in: J.P. Scalny (Ed.), Materials science of concrete, American Ceramics Society, vol. IV 1995, pp. 41–90. [4] S. Ng, T. HjellastrÖm, Quantitative and qualitative analysis of CAC and its hydrate products, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 17–30. [5] M. Schmidt, H. Pöllmann, Hydration behaviour of CAC at various temperatures by isoperibolic heat flow calorimetry using acetates, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2008, pp. 93–108. [6] C. Gosselin, K.L. Scrivener, Microstructure development of calcium aluminate cements accelerated by lithium sulphate, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2008, pp. 109–122. [7] F. Goetz-Neunhoeffer, S.R. Klaus, J. Neubauer, Kinetics of CA and CA2 dissolution determined by QXRD and corresponding enthalpies of reactions, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 54–64. [8] S.R. Klaus, J. Neubauer, F. Goetz-Neunhoeffer, A. Buhr, D. Schmidtmeier, Application of heat flow calculation to synthetic calcium aluminate cement mixes, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 65–74. [9] B. Lothenbach, Thermodynamic modelling of effects of time and silica on the conversion process, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 153–164. [10] J. Bizzozero, K.L. Scrivener, Hydration of calcium aluminate cement based systems with calcium sulfate and limestone, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 189–196. [11] S. Zhang, W.E. Lee, Spinel-containing refractories, in: C.A. Schacht (Ed.), Refractories Handbook, Marcel Dekker, Monticello 2004, pp. 81–139. [12] A. Buhr, D. Schmidtmeier, G. Wams, S. Kuiper, S. Klaus, Testing of calcium aluminate cement bonded castables and influence of curing conditions on the strength development, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 437–449. [13] J. Kasper, O. Krause, Mixing optimization of an alumina based LC-castable by applying variable power inputs, in: C. Fentiman, R. Mangabhai, K. Scrivener (Eds.),Calcium aluminate cements, Avignon, France 2014, pp. 450–456. [14] M.A.L. Braulio, M. Rigaud, A. Buhr, C. Parr, V.C. Pandolfelli, Spinel-containing alumina-based refractory castables, Ceram. Int. 37 (2011) 1705–1724.
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Please cite this article as: B. Han, et al., Hydration behavior of spinel containing high alumina cement from high titania blast furnace slag, Cem. Concr. Res. (2015), http://dx.doi.org/10.1016/j.cemconres.2015.09.019