Journal of the European Ceramic Society 41 (2021) 5634–5643

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Journal of the European Ceramic Society

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Sintering and grain growth behaviour of magnesium aluminate spinel: Effect of lithium hydroxide addition

Ali Talimian a,*, H.F. El-Maghraby b,c, Monika Michalkov´ a´ b, Duˇsan Galusek a,b a Centre for Functional and Surface Functionalised Glass, Alexander Dubcek University of Trencin, Trencin, Slovakia b Joint Glass Centre of the IIC SAS, TnUAD and FChPT STU, Trencin, Slovakia c Refractories, Ceramics, and Building Materials Department, National Research Centre, 33 El-Bohous St., 12622, Cairo, Egypt

ARTICLE INFO ABSTRACT

Keywords: Lithium hydroxide, LiOH, in the amounts ranging from 0.1 to 1.2 wt% has been used as a sintering aid to improve Magnesium aluminate spinel the densificationof MgAl2O4. The addition of 0.3 wt% LiOH promotes densificationand limits grain growth. The Sintering activation energy of sintering, calculated using master sintering curve approach, decreases from 790 ± 20 kJ. Master Sintering Curve 1 1 mol to 510 ± 20 kJ.mol with the addition of 0.3 wt% of LiOH. In addition, MgAl2O4 was also mixed with 10 High-temperature XRD wt% of LiOH to amplify the formation of reaction products. High-temperature XRD results showed that sec­ ◦ ondary phases (MgO and LiAlO2) are produced above 1040 C. The secondary phases start to disappear at T > ◦ 1200 C, and MgAl2O4 is produced. While adding small amounts of LiOH, up to ca. 0.3 wt%, is beneficial for + densification and suppressing grain growth, there exists a critical concentration of Li that is accounted for by the preferential incorporation of lithium ions into MgAl2O4 crystal lattice.

1. Introduction sintering: the incorporation of lithium ions into MgAl2O4 changes cat­ ions’ stoichiometry, introducing oxygen vacancies and improving Polycrystalline magnesium aluminate spinel is an interesting candi­ diffusion [21]. However, other fluorides,such as MgF2 and MnF2, have date for various engineering applications; it exhibits a favourable com­ also been reported to accomplish the task of LiF during sintering [17, bination of chemical and physical properties, such as high melting point, 22]. There are, therefore, some reservations about whether lithium has chemical inertness, low coefficient of thermal expansion, high thermal any significant influence on sintering. Although liquid LiF, acting as a shock resistance and excellent mechanical properties, i.e. hardness, and lubricant, facilitates particles’ rearrangement during sintering, it has fracture toughness [1–3]. Moreover, owing to its isotropic reflection detrimental effects on the properties of finalmaterial; grain growth and index and wide bandgap, polycrystalline magnesium aluminate spinel is cracking at grain boundaries are inevitable consequences of using LiF a cost-effective alternative to sapphire single crystals for optical appli­ that result in the decrease of the mechanical properties or impair the cations [4–6]. Producing dense magnesium aluminate spinel is a pre­ transparency [23–27]. requisite for obtaining excellent properties; however, sintering of While the effects of using LiF as sintering aid on the sintering MgAl2O4 is difficult due to the slow diffusion of constituent elements behaviour of MgAl2O4 have been studied extensively, there are few re­ and oxygen in particular [2,7–9]. Therefore, careful sintering processes ports on the influenceof other sources of lithium ions incorporation on at high temperatures with the application of pressure are required to the densification of magnesium aluminate spinel. Mordekovitz et al. densify MgAl2O4 [6,10–12]; even then, highly dense bodies are usually have studied the sintering of lithium doped magnesium aluminate spinel produced with the help of sintering additives, such as CaO, B2O3, AlF3, [28]. Although the instability of Li-doped MgAl2O4 during sintering transition element fluorides, or LiF [13–20]. Among these additives, results in the formation of secondary phases, such as MgO and γ-LiAlO2, only LiF has been consistently used to fabricate highly dense MgAl2O4. that can suppress grain growth by Zener pinning [28], no meaningful LiF promotes sintering during early-stage densification by producing change in the densificationof samples was observed. Therefore, the role transient liquid. Also, there had been a general agreement on how of lithium ions as solid-state sintering aids remains equivocal. Moreover, lithium facilitates sintering, particularly during the final stages of the possible reactions between the lithium source and MgAl2O4 during

* Corresponding author. E-mail address: [email protected] (A. Talimian). https://doi.org/10.1016/j.jeurceramsoc.2021.05.003 Received 24 October 2020; Received in revised form 26 April 2021; Accepted 1 May 2021 Available online 6 May 2021 0955-2219/© 2021 The Author(s). Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license

(http://creativecommons.org/licenses/by-nc-nd/4.0/). A. Talimian et al. Journal of the European Ceramic Society 41 (2021) 5634–5643 the initial stages of sintering have never been examined. LiOH) spread over a Pt strip acting both as heater and as sample holder, In the present work, the sintering behaviour of a commercial mag­ using the hight temperature diffraction chamber Anton Paar HTK 16. nesium aluminate spinel powder doped with LiOH has been studied. The The measurements were performed by heating the samples at the con­ ◦ ◦ effects of the addition of LiOH, up to 1.2 wt%, on the densificationand stant heating rate of 5 C up to 1200 C and recording the diffraction ◦ ◦ grain growth of MgAl2O4 were investigated. The reactions between Li2O pattern every 10 C, over the 2θ range between 20 and 38 . Additional and MgAl2O4 during the sintering and the stability of transient phases HT-XRD isothermal experiments were also carried out at the tempera­ ◦ have been studied by adding a relatively large amount of LiOH, 10 wt%, tures 1050, 1070, and 1120 C. The measured data were fitted and to make any reaction detectable. analysed by Rietveld refinement methods using the MAUD software package [31]; the total occupancies of tetrahedral and octahedral sites 2. Experimental procedures were constrained to stoichiometry values describing the Mg-Al distri­ + bution and assuming both sites are able to host Li . The background was Commercial magnesium aluminate spinel powder, S30CR (Bai­ modelled using a fourth-order polynomial, and the crystallographic kowski, Paris, France), and lithium hydroxide monohydrate, ACS grade variables were the lattice constants and occupancies of the tetrahedral > 99.0 (Sigma-Aldrich, MO, USA) were used as raw materials in this and octahedral sites. study. The main impurities of the spinel powder, reported by the sup­ plier, are (in wt ppm): Na: 70, K: 60, Ca: 60, Si:30, Fe:15, and S:600. The 2.1. Master Sintering Curve spinel powder was dispersed in isopropanol using an ultrasonic mixer (Sonopuls HD 3400, BANDELIN, Berlin, Germany). Then, an aqueous The densification kinetics of samples was studied by constructing solution of LiOH (25 mg/mL) was added to the suspension in order to Master Sintering Curves (MSC) and following the method developed by prepare a mixture with up to 1.2 wt % of LiOH, with respect to spinel. Su and Johnson [32]. The instantaneous densification rate of samples Afterwards, the mixture was transferred to a rotary evaporator, during sintering, (dρ/dt), can be described using the combined-stage ◦ concentrated under vacuum, and then dried at 120 C overnight. The sintering model proposed by Hansen et al. [33]: ( ) obtained powder was granulated by passing through a sieve of 0.5 mm 1 dρ 3γΩ δD Γ D Γ = b b + V V mesh. 4 3 (1) ρ dt kBT G G Cylindrical pellets with a diameter of 8 mm were produced using uniaxial pressing at 80 MPa. A portion of samples was sintered in a where γ and Ω are the surface energy and atomic volume, respectively; D thermo-mechanical analyser (TMA 402 Hyperion, Netzsch, Germany); being the diffusion coefficient. Γ is a scaling parameter related to the the temperature was increased at constant heating rates of 5, 10 and 20 ◦ ◦ driving force, mean diffusion distance and the microstructural features C min 1 up to 1550 C and a dwell time of 1 min. Some samples were ◦ of sintered samples. The indices are pointing out the diffusion mecha­ prepared by sintering at 1400 and 1500 C with a constant heating rate nism: b, grain boundary and V, lattice diffusion. K is Boltzmann con­ ◦ – .B. of 10 C and various dwelling time between 1 240 min in an electric stant, and T represents absolute temperature. G represents the grain size. furnace under ambient condition; the samples were removed from the It is considered that the diffusion is a thermally activated process furnace and quenched to room temperature to freeze the microstructure. following the Arrhenius relation: The density of samples was measured using Archimedes’ principle, with the theoretical density of MgAl O assumed to be 3.58 g cm 3. Q 2 4 D = D0exp( ) (2) The chemical reactions between the spinel powder and LiOH were RT investigated using a Simultaneous Thermal Analyser (STA 449 F1 where D , Q and R are a pre-exponential factor of diffusion, the apparent Jupiter®, Netzsch, Germany) in DTA-TG configurationupon heating to 0 ◦ ◦ diffusion activation energy and the gas constant, respectively. The sin­ 1350 C using a constant heating rate of 10 C.min 1. The measurements tering mechanism is controlled by the diffusion of the slowest species in were carried out by heating a powder mixture comprising MgAl2O4 and the fastest diffusion path [2]; hence, one can assume that the densifi­ 10 wt % LiOH; a relatively large amount of LiOH was added to the spinel cation occurs through only one mechanism [32,34], and Eq. (1) can be powder to make any reaction detectable, adopting the approach applied simplified, rearranged and integrated: by Rozenburg et al. [29]. The powder mixture was prepared by ∫ ρ m ∫ t dispersing the spinel powder in an aqueous solution of lithium hy­ KB 1 G 1 Q ◦ ∙ dρ = exp( ) dt (3) droxide, with subsequent drying at 60 C for 24 h. γΩD0 ρ0 3ρ Γ 0 T RT A scanning electron microscope, SEM, (JEOL 7600 F, JEOL, Tokyo, Japan) was used to examine the microstructure of sintered samples. where m is a constant depending on the diffusion mechanism. Moreover, Ceramographic cross-sections were prepared by cutting small samples the grain size of samples is assumed to be independent of thermal history from the centre of sintered bodies, which were then ground and polished and, hence, G and Γ depend only on the density of samples. Considering with diamond slurries down to 1 μm. The cross-sections were then this, Eq. (3) can be simplified further into two sets of the equations as: cleaned in an ultrasonic bath with acetone. Afterwards, the samples ∫ ρ m ◦ KB 1 G were subjected to thermal etching at the temperature 50 C below the Φ(ρ) = ∙ dρ (4a) γΩD0 3ρ Γ sintering temperature for 5 min in an electric furnace. The samples were ρ0 lightly coated with carbon. The grain size was estimated by measuring ∫ t 1 Q the size of at least 400 grains over five different regions of the sample Θ(t, T(t) ) = exp( ) dt T RT (4b) using line-cut method [30]. 0 Phase analyses were performed using an X-ray diffractometer where Eq. (4a) is independent of the thermal history and quantifies ef­ (Empyrean, Malvern Panalytical, Almelo, Netherlands) using Cu-Kα fects of the microstructure and material properties on densification, radiation (45 kV and 40 mA) and equipped with a high-temperature cell while Eq. (4b) depends only on sintering history. The relation between ρ (Anton Paar, HTK 16). First, the reference pellets consisting of the spinel and Θ (t, T(t)) is defined as the master sintering curve; here Q is an powder and 10 wt% LiOH were heat-treated at various temperatures ◦ average of the activation energies of sintering and hereafter is called between 1200 and 1500 C for 1 h in an electric furnace. After heat apparent activation energy. treatment, the pellets were removed from the furnace, pulverised and measured by XRD at room temperature. High-temperature XRD, HT-XRD, experiments were carried out by collecting the diffraction patterns from the powder mixture (10 wt%

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3. Results

3.1. Effects of LiOH on MgAl2O4 densification

◦ The relative density and porosity of samples, sintered at 1550 C with a dwell time of 1 min and containing various concentrations of lithium hydroxide, are summarised in Table 1; the corresponding atomic con­ centration of lithium ions is also presented. The addition of lithium hydroxide has improved the densification: the relative density has increased from 95.0% to 98.5% with the addition of 0.3 wt% LiOH; however, the density has remained virtually unchanged for higher LiOH additions. Fig. 1a shows the shrinkage of LiOH-free sample and the samples ◦ containing various amounts of LiOH upon heating at 10 C min 1 up to ◦ 1550 C. The LiOH-free sample exhibits an onset of shrinkage at around ◦ ◦ 1100 C; afterwards, the shrinkage continues until 1550 C. Although the addition of 0.1 and 0.3 wt% of LiOH has no significantinfluence on the onset shrinkage temperature, the samples containing 0.6 or 1.2 wt% of lithium hydroxide exhibit almost similar shrinkage with a lower ◦ sintering threshold at around 950 C. Moreover, the shrinkage curves of samples containing 0.3, 0.6 and 1.2 wt% of LiOH reached a plateau in the measured temperature range. The sintering behaviour of samples containing 0.3 and 1.2 wt % LiOH was investigated in order to study the influence of excessive lithium amount on densification. The shrinkage rate was estimated by calculating the slope of shrinkage curves using the Eq. (5):

1 d(Δl) 1 ΔlT+δT ΔlT δT = . T˙ (5) l dt l 2δT + 0 0 Fig. 1. (a) Shrinkage curves and (b) shrinkage rate curves of free- and Li doped-MgAl2O4 samples against temperature. Due to large difference of the where l and l0 are the instantaneous and the initial lengths of the sample, shrinkage rates of the 0.3-LiOH the data are presented separately, on the right Y ˙ respectively; T is the temperature and T being the heating rate. The axis (highlighted in red) (For interpretation of the references to colour in this instantaneous densification rates are smoothed by averaging ten adja­ figure legend, the reader is referred to the web version of this article). cent points. Fig. 1b shows the shrinkage rates of samples as a function of ◦ ◦ temperature between 800 C and 1550 C; samples containing 0.3 wt% The influence of lithium hydroxide addition on the sintering of LiOH show a significantly larger shrinkage rate (almost 6 times) in magnesium aluminate spinel was further evaluated using the master comparison with other samples. The corresponding values are therefore sintering curve approach (MSC); a series of dilatometry measurements ◦ presented separately on the secondary axis, on the right side of the graph were carried out using heating rates of 5, 10 and 20 C min 1. The effect (highlighted in red). The addition of limited amounts of LiOH promotes of lithium hydroxide decomposition on the mass of samples during the densification of magnesium aluminate spinel significantly. The + sintering was neglected due to the limited change in weight (less than shrinkage rate curve of Li -free is characterised by a broad peak at 0.5 %). The instantaneous relative density ρ(T) of the samples was ◦ around 1400 C. The temperature corresponding to the maximum calculated using the Eq. (6) [35]: shrinkage rate decreased with the addition of LiOH. The samples con­ ⎛ ⎞ ◦ 3 taining 0.3 wt% of LiOH show the maximum shrinkage rate at 1350 C, ⎜ ⎟ ◦ ⎜ 1 ⎟ while the 1.2 wt% samples exhibit a maximum shrinkage rate at 1175 C ρ(T) = ρg∙ (6) ◦ ⎝1 + Δl⎠ with a shoulder at 1350 C. l0 Fig. 2 shows representative microstructures of samples containing 0.0, 0.3 and 1.2 wt% LiOH and sintered using the sintering regime where ρ is the relative density of green bodies measured using the shown in Fig. 1. The sample containing 0.3 wt% LiOH exhibits a smaller g weight and dimensions of the samples. grain size in comparison to LiOH-free sample, while the addition of 1.2 Fig. 3 shows the grain size as a function of sintering regimes and wt% LiOH results in a microstructure with larger grains (1.0 μm vs 1.5 sample composition. The boxes represent the 25th and 75th percentile, μm). Moreover, residual pores located between the grains are visible in whereas the whiskers indicate 5th and 95th percentile; the median grain the microstructure of LiOH-free sample, marked by arrows in Fig. 2a; + size is also marked. The mean grain size of samples slightly increases this is related to incomplete densification of the Li -free samples. ◦ with the decreasing heating rate: especially for the heating rate of 5 C min 1. This is mainly related to the higher density of samples. Statistical Table 1 differences in the grain sizes were investigated using T-test. Although Relative density, as a function of theoretical density (3.58 g cm 3), of samples there is a change in the mean grain size of samples, the differences ◦ subjected to conventional sintering at 1550 C for 3 min (dilatometry mea­ among the samples related to various LiOH concentrations are statisti­ surements); the numbers in parentheses represent the standard deviation of cally insignificant. The microstructure can be thus considered as inde­ measurements. pendent of the heating rate, allowing the use of the MSC theory for + LiOH concentration Li (atom %) Relative density (%) evaluation of apparent activation energy of sintering.

LiOH-free 0 95.1(0.4) The MSCs were constructed from the dilatometry measurements 0.1 0.34 96.5(0.4) using the method described in the experimental part and following the 0.3 1.08 98.3(0.3) approach proposed by Maca et al. using the green densities of samples 0.6 2.14 97.5 (0.6) (ρgreen ≈ 43.0 ± 2.0 %) [35]. Fig. 4 shows MSC corresponding to the 1.2 4.25 97.7(0.5) LiOH-free sample and the samples with 0.3, and 1.2 wt% LiOH; the

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activation energy of 510 ± 15 kJ mol 1; the activation energy is significantlylarger than the reported values for liquid phase sintering of magnesium aluminate spinel doped with LiF [16]. The trajectories of the samples doped with 1.2 wt% LiOH begin to diverge at densities >60 % (θ = 10-20.3 K s-1); the poor overlap of the curves resulted in large error of the calculated activation energy of sintering. This is related to the oversimplifying of the process of sintering by assuming that it is controlled only by one mechanism. Nonetheless, the divergence points out that the governing mechanism changed during sintering; this be­ ◦ comes more significantin the case of samples sintered at 5 C min-1 [38].

3.2. Chemical reactions between LiOH and MgAl2O4

The chemical reactions between magnesium aluminate spinel and the sintering aid were studied using a sample with a relatively large LiOH content, 10 wt% vs 1.2 wt%, to amplify the effect of sintering and make any reaction product detectable. Fig. 5 shows the DTA-TG result of the reference MgAl2O4 samples with the 10 wt% of lithium hydroxide; the T.G. curve exhibits three main regions: (i) a rapid weight-loss (ca. ◦ ◦ 3%) around 100 C, (ii) a slow weight decrease between 100 C and 550 ◦ ◦ ◦ C, (iii) and a weight-loss between 550 C–800 C. The T.G. curve rea­ ◦ ches a plateau over 800 C. The sharp endothermic peak in the DTA ◦ curve at 100 C confirmsthat the firstweight loss of samples is related to the dehydration of lithium hydroxide. Further decreases in samples’ weight are related to thermal decompositions of solid lithium hydroxide. ◦ There is a sharp endothermic peak at 420 C that is probably attributed to the eutectic melting of LiOH; melting of LiOH results in an apparent ◦ deviation from the DTA curve base-line between 480 850 C, and the fast decomposition of LiOH [39]. Moreover, there is a small endothermic ◦ peak at 724 C which can be attributed to the melting of lithium car­ bonate, Li2CO3, which is produced by the reactions of the hydroxide with COx species from DI water or the ambient atmosphere during mixing or drying. The decomposition of LiOH produces solid Li2O below ◦ 800 C, which is significantly lower than the shrinkage onset of the samples containing LiOH as sintering aid (Fig. 1). The weight of the sample remains constant with further temperature increase. However, the observed constant rate decrease of DTA signal implies the existence of a solid-state reaction between Li2O and MgAl2O4 that yields a product with different heat capacity. The reaction between Li2O and MgAl2O4 + and incorporation of Li ions into the spinle structure promote the ◦ densifitcaion of samples above 850 C (Fig. 1). Fig. 6 shows the X-ray diffraction patterns of samples containing 10 wt% of lithium hydroxide and subjected to heat treatments at various ◦ ◦ temperatures between 1200 C–1500 C; the samples were removed from the furnace after one-hour dwelling and quenched to room tem­ perature in order to preserve the phase composition. In addition to magnesium aluminate spinel, samples subjected to the heat treatment contain periclase and lithium aluminate as secondary phases. Lithium aluminate is mainly in the form of γ-LiAlO2 (high-temperature phase) and traces of α-LiAlO2 (low-temperature phase); the latter was produced by phase transition during cooling. Table 2 summarises the result of quantitative phase analysis ◦ following the Rietveld refinementmethod and assuming only crystalline Fig. 2. SEM images of MgAl2O4 samples sintered at 1550 C for 1 min (dila­ tometry measurements) (a) LiOH free, (b) 0.3 wt% and (c) 1.2 wt% LiOH. phases were present in the samples. The concentration of MgAl2O4 in­ creases with the temperature of heat-treatment, while the amount of LiAlO decreases. The concentration of MgO is approximately constant, apparent activation energy associated with the lowest Mean Perpendicular 2 and MgAl O concentration increases with the increasing temperature. Curves Distance (MPCD) among the curves is also reported and shown as 2 4 The reaction between Li O and MgAl O can be described by Eq. (7): the insets in Fig. 4 [35]. The LiOH-free sample is characterised by an 2 2 4 1 apparent activation energy of sintering of 790 ± 20 kJ.mol . The MgAl2O4(s) + Li2O(s)→MgO(s) + 2LiAlO2(s) (7) measured value is in good agreement with the reported values in the Fig. 7a shows the in situ high-temperature XRD patterns of the literature for samples shaped by cold isostatic pressing (750 kJ mol 1) reference samples with 10 wt% LiOH, as a function of temperature: the [36] while it is significantlylower than the reported energy of samples ◦ 1 data were recorded upon heating with the constant heating rate of 10 C produced by slip casting (950 kJ.mol ) [37]. The addition of lithium ◦ min 1. γ-LiAlO forms at temperatures over 1020 C. The peaks of hydroxide results in a significant decrease in the apparent activation en­ 2 γ-LiAlO become more intensive at higher temperatures, indicating that ergy of sintering; the samples doped with lithium hydroxide show the 2 the concentration of γ-LiAlO2 increases. Moreover, MgAl2O4 peaks shift

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Fig. 3. Grain size distribution of different MgAl2O4 samples as a function of heating rate during sintering and powder preparation; median value, 25th and 75th percentiles, maximum and minimum values are shown.

◦ more than what is expected from the thermal expansion of spinel (the LiOH and sintered at 1400 and 1500 C with a dwell time of 1 240 min. dashed lines); this is accounted for by changes in the lattice constant of The grain size of samples containing different amounts of LiOH remains + spinel as a result of the incorporation of Li in the structure [40–42]. unchanged and is almost identical at all relative densities lower than 90 Fig. 7b shows the relative concentration of phases as a function of %, which is the characteristic behaviour during the intermediate stage of temperature; the quantitative analysis was carried out by fitting the sintering [53,54]. There are, however, two distinct trajectories for collected spectra using the Rietveld method and assuming that MgO is densities larger than 92 %: while samples containing 0.1 and 0.3 wt% of produced according to the stoichiometry of Eq. 7. The concentration of LiOH are characterised by larger densities and smaller grain size, the γ-LiAlO2 and, accordingly, MgO increases continuously with tempera­ addition of larger amounts of LiOH resulted in significantgrain growth. ◦ ◦ ture from 1040 C to 1170 C; then, it decreases. This indicates that the Lithium oxide, produced from LiOH, is responsible for promoting the decomposition of LiAlO2 and the reaction of the products at elevated sintering of samples by facilitating oxygen diffusion. Although the faster temperatures results in the formation of “secondary spinel” [43,44]. oxygen diffusion would be expected to aid both densification and microstructure coarsening, the densificationrates decrease at the larger 4. Discussion concentrations of LiOH. XRD results showed that the reaction between MgAl2O4 and Li2O yields the secondary phases, MgO, γ-LiAlO2. After­ The addition of lithium hydroxide to magnesium aluminate spinel wards, at higher temperatures, the reactions between secondary phases promotes densification.The results of the thermal analysis indicate that reproduces MgAl2O4 (Figs. 6 and 7), in the absence of liquid phase. ◦ lithium hydroxide yields lithium oxide below 800 C. Moreover, despite However, the kinetics of the reaction, the stability of phases and how the the lower sintering onset of samples containing 1.2 wt% LiOH, their reformation of spinel contribute to the sintering require more in­ shrinkage rate is similar to the additive-free samples, which indicates a vestigations. The reactions between MgAl2O4 and Li2O were studied similar densification mechanism is at play. Considering the melting further using in situ HT-XRD under isothermal heat treatment condi­ ◦ temperature of Li2O, 1438 C [45], no liquid phase is present in samples tions. Fig. 9 shows the relative concentration of γ-LiAlO2 and MgAl2O4 ◦ at the initial stages of sintering. Moreover, HT-XRD confirmedthat Li2O as a function of time at 1020, 1070 and 1120 C. While the concentration reacts with MgAl2O4 producing high melting-point phases (Tm γ-LiAlO2 of γ-LiAlO2 continuously increases during the isothermal heat treatment ◦ ◦ ≈ 1785 C) [45]. at 1020 and 1070 C, its concentration increases in the first 1200s and The apparent activation energy of sintering for LiOH doped samples is afterwards decreases. Interestingly, the finalcomposition of all samples significantly larger than the values reported for LiF doped spinel (510 is almost identical, regardless of the treatment temperature. In the early kJ. mol 1 vs 300 kJ.mol 1) [16]. It can be therefore assumed that, unlike stages of heat treatment, the concentration of phases depends on both using LiF, the densification is governed by the contribution of grain time and temperature, while the equilibrium concentration of phases is boundary mobility and pore mobility through solid-state diffusion of independent of temperature. The secondary phases are present in the ◦ cations and oxygen. Due to the significantlysmaller diffusion coefficient system even after several hours at 1500 C (Fig. 6); the inclusion par­ of oxygen in comparison with the cations in MgAl2O4 structure (3.8 × ticles might prevent grain growth by the Zener pinning effect [28]. 10-6 m2s-1 vs ~ 2 × 10-2 m2s-1) [46–48], the diffusion of oxygen ulti­ Although further microstructural examination, e.g. TEM examinations, mately controls the densification of the spinel ceramics [7–9]. The is required to show the segregation of secondary phases at grain decrease of apparent activation energy of sintering implies that the boundaries, it is plausible that the increase of additive concentration introduction of LiOH improves the diffusion of oxygen and, in turn, and, in turn, the number of inclusions would suppress grain growth and promotes the densification. Interestingly, the activation energy of sin­ result in microstructure refinement.However, the results were quite the tering of lithium doped samples is of the same order as the activation opposite: a lower concentration of lithium hydroxide (0.3 wt%) pro­ energy of oxygen diffusion (450 550 kJ.mol 1). [49–52] duces dense, fine-grainedmicrostructures, while higher amounts of the Fig. 8 shows the sintering trajectories (the dependence of grain size additive result in grain growth (Fig. 8). against the relative density) of samples doped with various amounts of Moreover, the shift of spinel diffraction peaks upon heating is larger

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Fig. 4. Master sintering curve and MCDP (inset) calculated for samples produced from (a) LiOH-free MgAl2O4, and samples doped with (b) 0.3 wt% and (c) 1.2 wt % LiOH. than what is expected form thermal expansion (Fig. 7-dashed lines): this using dilatometry measurements. The lattice constants are larger than indicates that the introduction of Li2O changes the crystal structure of the estimated ones from the thermal expansion, particularly at tem­ ◦ MgAl2O4. Magnesium aluminate spinel is capable of producing solid- peratures higher than 950 C. Such differences imply that some cations solution with Li2O [44,55]. Fig. 10 shows spinel lattice constant as a are replaced with larger ones in spinel structure [40]; this becomes more ◦ 1 3+ function of temperature upon heating at a constant rate of 10 C min ; significantif octahedral sites are filledwith magnesium or lithium: (rAl : 2+ + the data were measured by refining the HT-XRD results. 0.59 Å vs rMg: 0.72 Å rLi: 0.76 Å). The XRD results confirmed that the The lattice constant is influenced by both thermal expansion of the reaction between MgAl2O4 and Li2O yields MgO and LiAlO2 which can material and the dissolution of ions; the influenceof thermal expansion be dissolved in MgAl2O4, producing nonstoichiometric spinel according on the lattice constant of MgAl2O4 was therefore calculated using the to the Eq. 9-a and 9-b: empirical formula reported by Kapralik [56]: MgAl O 3MgO ̅̅̅̅̅̅2 ̅4→ Mg× + 2Mg/ + 3O× + V∙∙ 5 8 2 Mg Al O O (9a) a(T) = 8.066 + 5.015 × 10 T + 1.550 × 10 T (8) MgAl O LiAlO ̅̅̅̅̅̅2 ̅4→ Li/ + Al× + V/// + 2O× + 2V∙∙ where T is the absolute temperature. Also, changes of lattice parameter 2 Mg Al Al O O (9b) were calculated using the coefficient of thermal expansion determined + 2+ The incorporation of Li or Mg into MgAl2O4 is accompanied by

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Fig. 5. DTA /TG curves for powder mixture containing magnesium aluminate ◦ spinel and 10 wt% LiOH upon heating at the constant rate of 10 C.min 1.

Fig. 7. (a) In-situ high-temperature XRD patterns (in colour contour) of MgAl2O4-10 wt% LiOH mixture (b) relative concentration of MgAl2O4 (◼), ● γ ◆ Fig. 6. XRD patterns of samples comprising MgAl2O4 and 10 wt% LiOH at MgO ( ) and -LiAlO2 ( ) as a function of temperature. ◦ room temperature and after subjecting to heat-treatment at 1200-1500 C for 1 h. (MgAl2O4 (◼), MgO (●), γ-LiAlO2 (◆), α -LiAlO2 (▾), and LiOH.H2O ( )). Inset: zoom in on the peak shifts of MgO and MgAl2O4.

Table 2 The phase composition, in wt%, of magnesium aluminate spinel doped with 10 ◦ wt % LiOH and subjected to heat treatment at a temperature between 1200 C ◦ and 1500 C for 1 h (Fig. 6). Numbers between parentheses represent the error in quantitative analyses (N.D.: not detected).

o Temperature ( C) MgAl2O4 MgO γ-LiAlO2 α-LiAlO2 LiOH. H2O

Room 94.5 N.D. N.D. N.D. 5.4 (0.6) Temperature (0.5) 1200 62.5 8.5 (0.9) 25.2 (2.5) 3.8 (0.6) N.D. (6.2) 1300 64.3 10.1 23.8 (1.8) 1.8 (0.8) N.D. (4.8) (0.7) 1400 66.3 10.7 21.6 (2.0) 1.5 (1.4) N.D. (6.1) (0.9) 1500 76.1 7.7 (0.6) 14.62 1.5 (0.9) N.D. (6.0) (1.2) Fig. 8. Grain size against relative density (sintering trajectory) of samples doped with various amounts of LiOH (0-1.2 wt%); the dashed lines are to guide the production of oxygen vacancies to preserve charge neutrality which the eye. facilitates the diffusion of oxygen, promoting the densification of MgAl2O4. the second phases inhibit grain growth, and the final grain size is The formation of oxygen vacancies accounts for the decrease of inversely proportional to the volume fraction of precipitates [57]. In our apparent activation energy of sintering. However, the peculiar grain + case, the grain growth was observed mainly in the samples with the level growth of samples doped with high amounts of Li does not meet the of LiOH doping > 0.3 wt% (Fig. 8). Mordekovits et al. have reported on expected behaviour characteristic for Zener pinning mechanism where the thermal stability of Li-doped magnesium aluminate spinel: there is a

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while the substitution of magnesium sites by lithium ions above the critical concentration can be described by the Eq. (11):

2MgAl2O4 / /// × ∙∙ Li2O ̅̅̅̅̅̅̅→ 2LiMg + 4VAl + OO + 7VO (11)

Although such substitutions should be limited, particularly in the + case of Li/Al due to the differences in ionic charge and ionic radii of Li + and Al3 , they increase the concentration vacancies and improve the diffusion of oxygen at the surfaces dramatically. Consequently, grain growth increases significantly. Indeed, more studies are required to determine the critical concentration of lithium, its effect on the inver­ sion of spinel structure and to identify its role in the sintering of mag­ nesium aluminate spinel. Although the results indicated that reactions between LiOH and MgAl2O4 produce transient phases (LiAlO2 and MgO) at grain bound­ aries and, thus, affect the grain growth of magnesium aluminate spinel, one finalcomment should be made about the approach used to identify the transient phases. Large amounts of LiOH have been added to

Fig. 9. Relative concentration of MgAl2O4 and γ-LiAlO2 against time during MgAl2O4 to amplify the reaction products due to the limited sensitivity ◦ isothermal heat-treatment at 1020, 1070 and 1120 C in the reference sample of the used characterisation techniques. However, the reactions and doped with 10 wt% LiOH. densificationmechanism might be altered by the additive concentration. Therefore, further studies, i.e. TEM examination, or electrochemical impedance spectroscopy of densifying bodies, are necessary to identify the reactions at the grain boundaries and shed light on the mechanisms at play during sintering MgAl2O4 doped with small concentrations of LiOH.

5. Conclusions

The addition of LiOH improves the densification of magnesium aluminate spinel; using small amounts of LiOH (ca. 0.3 wt%) decreases the apparent activation energy of sintering of MgAl2O4 significantly (790 ± 20 kJ.mol 1 vs 510 ± 15 kJ.mol 1), and suppresses grain growth. + Incorporation of Li ions into MgAl2O4 changes the stoichiometry by introducing oxygen vacancies; this, in turn, enhances the diffusivity of ions and decreases the apparent activation energy of sintering. However, adding excessive LiOH results in grain growth without a further decrease in the activation energy of sintering. There is a critical amount of lithium above which it preferentially occupies tetrahedral sites of the spinel structure at grain boundaries, promoting grain

Fig. 10. Lattice constant of MgAl2O4-spinel against temperature: the lattice boundary mobility and thus, also grain growth. constant is determined form in-situ high- temperature XRD of (MgAl2O4-10 wt % LiOH) samples (Fig. 7-a); lattice constant estimated using measured thermal Contributions expansion of samples measured by dilatometry (black line) is also shown for comparison. The manuscript was written through the contribution of all authors. The idea was conceived by A.T. A.T, and H.E. produced the samples and γ critical concentration of lithium beyond which Mg(Li,Al)O and -LiAlO2 performed electron microscopy and analysis of Results. D.G. acquired precipitate [58]. Caliman et al. have reported that due to the segregation the in-situ high-temperature XRD. Thermal analyses were carried out by of lithium to surface or grain boundary, adding more than 5 mol% of M.M.; D.G. provided supervision. All authors contributed to the text and lithium to MgAl2O4 results in grain growth [59]. Therefore, there is a gave their approval to the final version of the manuscript. critical concentration of lithium beyond which grain growth is dominant during the sintering process. It is interesting to consider the atomic + concentration of Li ions in samples (Table 1 and Fig. 8): the critical Declaration of Competing Interest concentration of lithium is between 1.08 and 2.14 at%. Bernatdes et al. have reported that the segregation of lithium ions in magnesium + 3+ The authors declare that they have no known competing financial aluminate spinel depends on its concentration: while Li replaces Al interests or personal relationships that could have appeared to influence ions for lithium-ion concentration below 1.8 mol %, for samples con­ + the work reported in this paper. taining higher concentrations of lithium, Mg2 is preferentially substituted, particularly at the surfaces [42]. Considering that the for­ Acknowledgements mation of Schottky type defects is energetically favourable to Frenkel ones (~ 4 eV vs ~7 eV) in magnesium aluminate spinel [60,61], the This paper is a part of dissemination activities of the project Fun­ incorporation of lithium ions into the octahedral sites of MgAl2O4 can be + Glass. This project has received funding from the European Union’s described by Eq. (10) below the critical concentration of Li : Horizon 2020 research and innovation programme under grant agree­ MgAl2O4 // // × ∙∙ ment No 739566. Financial support of this work by the grants SAS- Li2O ̅̅̅̅̅̅̅→ VMg + 2LiAl + OO + 3VO (10) MOST JRP2015/6, and VEGA2/0026/17 is gratefully acknowledged. We appreciate the help of Dr A. Prnova with STA measurements.

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