- Email: [email protected]
Effect of the liquid fragility on flash sintering behavior of oxide nanoparticles
Scripta Materialia 178 (2020) 261–263
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Effect of the liquid fragility on flash sintering behavior of oxide nanoparticles Rachman Chaim Department of Materials Science and Engineering, Technion – Israel Institute of Technology, Haifa 32000, Israel
a r t i c l e
i n f o
Article history: Received 10 November 2019 Accepted 18 November 2019
Keywords: Flash sintering Liquid-film Fragility Oxides Capillarity Nanocrystalline
a b s t r a c t Transient liquid-film formed at the nanoparticle contacts due to the local Joule heating, assists the ultrafast densification of oxide nanopowders by capillarity during flash sintering. The retention of this liquidfilm as a residual intergranular amorphous layer or its crystallization depends on the fragility of the liquid-film, hence on its composition. Impurities concentrated at the particle surfaces and grain boundaries modify the flash temperature by affecting the electrical conductivity and viscosity of this liquid-film. These aspects were discussed for flash sintering of pure and doped aluminas as examples from the literature, where doping changed the flash onset temperatures and densification behavior.
Flash sintering (FS) as a novel technique for ultrafast densification of ceramic powders has been widely investigated. A wide range of ceramics, especially oxides, were fabricated at different electric field - flash temperature conditions, for investigating the various aspects of this technology. The columnar microstructures resulted from melting of 99.8% pure alumina during flash sintering [1], as well as residual grain boundary amorphous films in Yb:(LaY)2 O3 [2] and monoclinic ZrO2 [3] validated the liquid-film assisted sintering mechanism [4–7] as one of the active densification mechanisms. The liquid-film assisted sintering originates from the very high local temperatures developed at the particle contacts, due to the locally high electric resistance, hence local Joule heating that leads to local contact melting and thermal runaway. The liquid-film at the softened/melted contact fully wets the adjacent particles due to full solubility of the solid in its melt, followed by immediate particle rearrangement due to the liquid capillary forces [8,9]. This liquid-film has a transient nature and it may spread over the particle at the velocity of ~1 m s−1 , due to the high capillary forces [9]. Consequently, once the particle rearrangement takes place, the local electric resistance drastically decreases, which leads to the solidification of the liquid-film hence its transient nature. In this respect, while some oxides may retain solid amorphous films remnant of the transient liquid-film, other oxides may reveal no traces for the liquid-film involved process. The thermal instability onset, followed by fast densification at the flash event, was investigated by microwave sintering of Mg-doped
E-mail address: [email protected] https://doi.org/10.1016/j.scriptamat.2019.11.046 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Al2 O3 and related to the change from the solid state to liquidphase assisted processes [10]. Herein, we discuss and relate these differences in the liquid-film behavior to the molecular structure of the supercooled liquid and its transport-type properties, such as electric conductivity, viscosity and its change with temperature which is termed as liquid’s ’fragility’. Recently we showed that spontaneous crystallization of fragile liquid-film immediate to flash sintering leads to temperature difference between the lower bound flash onset temperature and Debye temperature [11]. In our previous paper, we referred the flash temperature of the oxide to its fundamental properties, such as crystal structure, average ionic potential, and the fusion entropy [12]. By normalizing the flash sintering conditions of different oxides we showed that the relative flash onset temperature for a given crystal-type increases with the increase in the average ionic potential. The ionic potential is expressed by the ratio of the ion charge to its radius (i.e. charge density on the ion), and represents the strength of the interionic bond. Therefore, the value of the flash onset temperature for a given crystal-type (i.e. cation coordination number), is relative to the melting point of the oxide [Fig. 2 in ref. 12]. However, we showed that the onset flash temperature strongly depends on the electric conducting nature of the oxides; non-conducting oxides exhibit higher temperatures than semi-conducting oxides, and in turn, than ionic conductors [12]. This arises from the need of some amount of electric energy density (~10–50 W cm−3 ) to locally melt the particle contacts and form the liquid-film needed for current percolation. However, once the percolating liquid is formed and the particles rearranged by the liquid capillary forces, one expects the immediate solidification of the liquid. Solidification of the liquid-
262
R. Chaim / Scripta Materialia 178 (2020) 261–263
film as an exothermic reaction is thermodynamically preferred [6]. Despite the transient nature of the liquid-film, the densification and the resultant sintered microstructure of the ceramic particles depend on the liquid-film properties such as its molecular structure, its electric conductivity, and its temperature-dependent viscosity. Below we will further discuss these properties. One of the common features of different oxides subjected to flash sintering is the similar dissipated power at the flash onset, irrespective of the flash temperature value [13]. This is in accordance with the formation of the liquid film at the flash temperature, since all oxide liquids exhibit relatively high electronic conductivity through the percolating liquid-film. However, the oxide melts exhibit different properties, defined by their chemical composition, hence dictate the resultant microstructure of the flash sintered ceramic. It is necessary to refer to the oxide liquid phase structure and properties in order to understand the final microstructure evolved by the flash sintering. The internal structure of the oxide glasses, liquids, and supercooled liquids were the matter of continuous investigation, due to their complex nature. However, at present, several aspects of the glass structure are well-accepted paradigms in this field. Traditionally, selected cations with low coordination number (SiO2 , GeO2 , P2 O5 , B2 O3 and As2 O3 ) are classified as the network-formers, to say they form thermodynamically stable glass. These class of glasses reveal three-dimensional network of interconnected rings, the typical size of which is also present in their crystalline counterparts [14]. The melting points of the glass-forming oxides are significantly different from each other (SiO2 - 1883 K, GeO2 - 1388 K, P2 O5 - 853 K, B2 O3 - 723 K, and As2 O3 - 313 K). These differences in the melting points, together with differences in their liquid viscosity, were related to the density of the crosslinked bonds (defined as the number of the bridging bonds per cation less 2) in the glass network [14]. Therefore, assuming formation of a liquid-layer during the flash sintering, the melting point, and the elastic modulus which is associated with glass softening temperature [15], can act as good measures for the onset flash sintering temperature of these oxides. The viscosity of the network-former glasses near the glass transition temperature (Tg) follows an Arrhenian behavior, hence they were named as ’strong glasses’. The largest group of cations are conditional network-formers and cannot form continuous networks by themselves from their oxide melts. Therefore, their retention as a solid glass may need extremely high cooling rates, if possible at all. Accordingly, the viscosity of their supercooled liquids is nonArrhenian and drastically increases below Tg, hence their liquid is termed ’fragile’. The fragile behavior is related to the liquid’s internal structure that composes polyhedra clusters similar to those existing in their crystalline counterpart. Wang et al. [16] explored the cluster structure within the liquid in metallic glasses, using molecular dynamic simulations supported by the experimental alloy fabrication. Although the cluster type was important for the crystallization of the melt, they found strong correlation between the variance of the cluster distribution and the glass forming ability (GFA) of the melt. Large variance of the cluster distribution leads to difficulties in the crystallization of the melt due to a large difference between the cluster types, hence an increase in the glass forming ability of the liquid. A similar trend may be applicable to ceramic systems, albite the cluster types in the ionic melts are better defined due to the electrostatic bonding, and extremely high cooling rates are impractical due to the low thermal conductivities. Based on the fact that any liquid can be formed as solid glass when a critical cooling rate is applied, Jiusti et al. [17] analyzed the GFA of glass forming oxides. They related the GFA to the cooling rate in general, and defined it to be proportional to η (TL )/(TL ) 2 , i.e. the ratio between the melt viscosity at the liquidus temperature divided by the square of the liquidus
temperature. In this respect, viscous flow flash sintering results of silica glass by Prado et al. [fig. 4 in 18] show that the flash onset in glass under different applied fields is associated with reaching a certain electric conductivity (hence viscosity), rather than a certain dissipated power. In order to highlight the effect of the liquid film nanostructure on the flash temperature, we will use alumina as an example, the liquid of which is extremely poor glass former, and exhibits non-Arrhenian viscosity. Several flash sintering studies consistently exhibited lower flash onset temperatures for MgO-doped Al2 O3 compared to pure Al2 O3 powders [19–21]. MgO-doped (0.25 wt.%) 99.9% alumina with 200 nm average particle size (diameter) was flash sintered at 1268 °C under the field of 10 0 0 V cm−1 , while pure Al2 O3 powder failed to flash sinter at this field, resulting in arcing at 1400 °C [19]. Alumina powder with larger particle size (d50 = 600 nm) and higher impurity content (i.e. 99.8%) was flash sintered around 1050 °C at the same applied field [20]. For comparison, 99.5% pure MgAl2 O4 spinel powder with 250 nm particle size was successfully flash sintered at 1410 °C at the same electric field [22]. Therefore, we can evaluate the effect of Mg dopant on the liquid film properties based on the above discussions. Alumina melts were intensively investigated for their melt structure, their electric conductivity and viscosity in the supercooled condition. Ansell et al. [23] used Synchrotron x-ray diffraction for levitating alumina melts, and determined the (AlO4 )5− tetrahedra to be the main structural units, where the Al coordination (number) changes from octahedral (6) in the solid to tetrahedral (4) in the liquid. Similar experiments by Skinner et al. [24] with molecular dynamic and reverse Monte Carlo simulations showed the predominant structural units to be AlO4 and AlO5 in the 2:1 ratio, and minor fractions of the AlO3 and AlO6 molecules in the liquid. They showed that 83% of the Al-O-Al connections were involved in corner sharing polyhedra. Therefore, alumina melt can form glass in certain extreme conditions [23]. Levi et al. [25] used extremely high cooling rates of ~105 Ks − 1 to form amorphous alumina. Decrease in the cooling rate resulted in metastable transition (Gamma) alumina, where Al ions are both tetrahedral and octahedral coordinated. Cooling rates of alumina melt below 102 K s−1 led to the crystallization of the stable alphaalumina [26]. This melt behavior may partially explain the lack of a residual grain boundary amorphous layer in flash sintered pure aluminas. On the other hand, the formation of (AlO4 )5− clusters in liquid Al2 O3 is believed to be responsible for the ionic conductivity of 6 ± 1 [ cm]−1 at the melting point [27]. The electric conductivity of the molten oxides depends primarily on the cation’s charge, size, and mass [28,29]. In this respect, generally, the increase in the cation charge, cation radius and cation mass lead to decease in the electric conductivity. Cations with multiple oxidation states exhibit the lowest electric conductivity for their molten oxide with the highest valence state. Thus, the electric conductivity of molten oxide follows approximately the location of their cations in the periodic table of the elements. Consequently, molten MgO exhibits higher electric conductivity (i.e. 35 [cm]−1 [28,29]) than molten Al2 O3 (i.e. 15 [ cm]−1 [28]; 6 [ cm]−1 [27]; 0.2 [ cm]−1 [30]). It is believed that the higher electrical conductivity values were due to the impurity content in molten alumina [30]. Electric conductivity in molten oxides is also inversely proportional to the melt viscosity according to Walden’s rule [4]. This arises from the size of the ionic clusters present within the melt and the degree of the covalent bonds (i.e. degree of ionicity) between the clusters (molecules). The viscosity of molten MgO (at 2852 °C) was calculated as 1.4 [mPas] together with its specific electrical conductivity as 20.9 [ cm]−1 [31]. For comparison, viscosity of pure molten Al2 O3 was reported as 40 [mPas] at the melting point (2072 °C), and decreased to ~15 [mPa s] at 2800 °C
R. Chaim / Scripta Materialia 178 (2020) 261–263
[32]. Assuming the viscosity of molten mixed oxides to depend linearly on the viscosities of its oxide constituents in small composition ranges [32,33], addition of MgO is expected to increase the electrical conductivity of molten MgO-doped alumina and decrease its viscosity, whereas SiO2 dopant will act in the opposite way. Since the flash event is expected to take place by percolation of the electric current through the liquid-film layer, one expects higher current through the MgO-doped liquid-film than the pure alumina liquid-film. Consequently, the dissipated power needed to initiate and proceed the flash should be reached at relatively lower temperatures than for pure alumina. Furthermore, the lower viscosity of the liquid-film in the MgO-doped alumina should increase its flow kinetics by rapid wetting of the surrounding particles and their liquid-aided rearrangement and densification. On the other hand, SiO2 exhibits extremely higher viscosities and lower electric conductivities than Al2 O3 at temperatures around the alumina melting point [32]. Therefore, one expects a reduced effect of densification during the flash sintering in SiO2 -doped alumina than in pure alumina or MgO-doped alumina. Indeed, flash sintering of alumina containing 8 wt.% SiO2 and 2 wt.% MgO exhibited flash temperature of ~700 °C under the field of 1500 V cm−1 , but with very limited densification; linear shrinkage of 0.3% was measured compared to 5.6% in almost pure alumina under the same electric field [34]. The theoretical observations and expectations mentioned above are in full agreement with the experimental flash sintering results from the literature [1,19–21,34] and explain the possible effects of the impurities on the liquidlayer properties, hence on the flash sintering temperature and behavior. Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.11. 046.
263
References [1] M. Biesuz, V.M. Sglavo, Scr. Mater. 150 (2018) 82–86. [2] Y.V. Bykov, S.V. Egorov, A.G. Eremeev, V.V. Kholoptsev, K.I. Rybakov, A.A. Sorokin, J. Am. Ceram. Soc. 98 (2015) 3518–3524. [3] N. Morisaki, H. Yoshida, T. Tokunaga, K. Sasaki, T. Yamamoto, J. Am. Ceram. Soc. 100 (2017) 3851–3857. [4] R. Chaim, Materials (Basel) 9 (2016) 280. [5] R. Chaim, Materials (Basel) 10 (2017) 179. [6] R. Chaim, C. Estournès, J. Mater. Sci. 53 (2018) 6378–6389. [7] J. Narayan, Scr. Mater. 176 (2020) 117–121. [8] R. Chaim, G. Chevallier, A. Weibel, C. Estournes, J. Appl. Phys. 121 (2017) 145103. [9] R. Chaim, Scr. Mater. 158 (2019) 88–90. [10] S.V. Egorov, A.G. Eremeev, V.V. Kholoptsev, I.V. Rybakov, A.A. Sorokin, Y. Bykov, Scr. Mater. 174 (2020) 68–71. [11] R. Chaim, Scr. Mater. 169 (2019) 6–8. [12] R. Chaim, C. Estournès, Scr. Mater. 163 (2019) 130–132. [13] R. Raj, J. Am. Ceram. Soc. 99 (2016) 3226–3232. [14] B. Bridge, N.D. Path, D.N. Waters, Phys. Stat. Sol. (a) 77 (1983) 655–668. [15] B.D. Sanditov, S. Sh. Sangadiev, D.S. Sanditov, J. Exper. Theor. Phys. 115 (2012) 445–454. [16] J. Wang, A. Agrawal, K. Flores, Acta Mater. 171 (2019) 163–169. [17] J. Jiusti, E.D. Zanotto, D.R. Cassar, M.R.B. Andreeta, J. Am. Ceram. Soc. (2019), doi:10.1111/jace.16732. [18] M.O. Prado, M. Biesuz, M. Frasnelli, F.E. Benedetto, V.M. Sglavo, J. Non-Cryst. Solids 476 (2017) 60–66. [19] M. Cologna, J.S.C. Francis, R. Raj, J. Eur. Ceram. Soc. 31 (2011) 2827–2837. [20] M. Biesuz, V.M. Sglavo, J. Eur. Ceram. Soc. 36 (2016) 2535–2542. [21] M. Biesuz, P. Luchi, A. Quaranta, V.M. Sglavo, J. Appl. Phys. 120 (2016) 145107. [22] H. Yoshida, P. Biswas, R. Johnson, M.K. Mohan, J. Am. Ceram. Soc. 100 (2017) 554–562. [23] S. Ansell, S. Krishnan, J.K. Richard Weber, J.J. Felton, P.C. Nordine, M.A. Beno, D.L. Price, M.-.L. Saboungi, Phys. Rev. Let. 78 (1997) 464–466. [24] L.B. Skinner, A.C. Barnes, P.S. Salmon, L. Hennet, H.E. Fischer, C.J. Benmore, S. Kohara, J.K.R. Weber, A. Bytchkov, M.C. Wilding, J.B. Parise, T.O. Farmer, I. Pozdnyakova, S.K. Tumber, K. Ohara, Phys. Rev. B 87 (2013) 024201. [25] C.G. Levi, V. Jayaram, J.J. Valencia, R. Mehrabian, J. Mater. Res. 5 (1988) 969–983. [26] J.K.R. Weber, C.D. Anderson, D.R. Merkley, P.C. Nordine, J. Am. Ceram. Soc. 78 (1995) 577–582. [27] J.E. Enderby, S. Ansell, S. Krishnan, D.L. Price, M.-L. Saboungi, Appl. Phys. Lett. 71 (1997) 116–118. [28] A.E. van Arkel, E.A. Flood, N.F.H. Bright, Canadian J. Chemistry 31 (1953) 1009–1019. [29] E. Thibodeau, I.-.H. Jung, Metall. Mater. Trans. B 47 (2016) 355–383. [30] A.N. Shatunov, A.I. Maksimov, A.Y. Pechenkov, I.V. Poznyak, Inorg. Mater. 47 (2011) 1579–1583. [31] A.-.L. Leu, S.-.M. Ma, H. Eyring, Proc. Nat. Acad. Sci. USA 72 (1975) 1026–1030. [32] G. Wu, E. Yazhenskikh, K. Hack, E. Wosch, M. Müller, Fuel Processing Tech. 137 (2015) 93–103. [33] G.-.H. Zhang, K.-.C. Chou, Metall. Mater. Trans. B 43B (2012) 849–855. [34] M. Biesuz, V.M. Sglavo, J. Eur. Ceram. Soc. 37 (2017) 705–713.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Effect of the liquid fragility on flash sintering behavior of oxide nanoparticles Rachman Chaim Department of Materials Science and Engineering, Technion – Israel Institute of Technology, Haifa 32000, Israel
a r t i c l e
i n f o
Article history: Received 10 November 2019 Accepted 18 November 2019
Keywords: Flash sintering Liquid-film Fragility Oxides Capillarity Nanocrystalline
a b s t r a c t Transient liquid-film formed at the nanoparticle contacts due to the local Joule heating, assists the ultrafast densification of oxide nanopowders by capillarity during flash sintering. The retention of this liquidfilm as a residual intergranular amorphous layer or its crystallization depends on the fragility of the liquid-film, hence on its composition. Impurities concentrated at the particle surfaces and grain boundaries modify the flash temperature by affecting the electrical conductivity and viscosity of this liquid-film. These aspects were discussed for flash sintering of pure and doped aluminas as examples from the literature, where doping changed the flash onset temperatures and densification behavior.
Flash sintering (FS) as a novel technique for ultrafast densification of ceramic powders has been widely investigated. A wide range of ceramics, especially oxides, were fabricated at different electric field - flash temperature conditions, for investigating the various aspects of this technology. The columnar microstructures resulted from melting of 99.8% pure alumina during flash sintering [1], as well as residual grain boundary amorphous films in Yb:(LaY)2 O3 [2] and monoclinic ZrO2 [3] validated the liquid-film assisted sintering mechanism [4–7] as one of the active densification mechanisms. The liquid-film assisted sintering originates from the very high local temperatures developed at the particle contacts, due to the locally high electric resistance, hence local Joule heating that leads to local contact melting and thermal runaway. The liquid-film at the softened/melted contact fully wets the adjacent particles due to full solubility of the solid in its melt, followed by immediate particle rearrangement due to the liquid capillary forces [8,9]. This liquid-film has a transient nature and it may spread over the particle at the velocity of ~1 m s−1 , due to the high capillary forces [9]. Consequently, once the particle rearrangement takes place, the local electric resistance drastically decreases, which leads to the solidification of the liquid-film hence its transient nature. In this respect, while some oxides may retain solid amorphous films remnant of the transient liquid-film, other oxides may reveal no traces for the liquid-film involved process. The thermal instability onset, followed by fast densification at the flash event, was investigated by microwave sintering of Mg-doped
E-mail address: [email protected] https://doi.org/10.1016/j.scriptamat.2019.11.046 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Al2 O3 and related to the change from the solid state to liquidphase assisted processes [10]. Herein, we discuss and relate these differences in the liquid-film behavior to the molecular structure of the supercooled liquid and its transport-type properties, such as electric conductivity, viscosity and its change with temperature which is termed as liquid’s ’fragility’. Recently we showed that spontaneous crystallization of fragile liquid-film immediate to flash sintering leads to temperature difference between the lower bound flash onset temperature and Debye temperature [11]. In our previous paper, we referred the flash temperature of the oxide to its fundamental properties, such as crystal structure, average ionic potential, and the fusion entropy [12]. By normalizing the flash sintering conditions of different oxides we showed that the relative flash onset temperature for a given crystal-type increases with the increase in the average ionic potential. The ionic potential is expressed by the ratio of the ion charge to its radius (i.e. charge density on the ion), and represents the strength of the interionic bond. Therefore, the value of the flash onset temperature for a given crystal-type (i.e. cation coordination number), is relative to the melting point of the oxide [Fig. 2 in ref. 12]. However, we showed that the onset flash temperature strongly depends on the electric conducting nature of the oxides; non-conducting oxides exhibit higher temperatures than semi-conducting oxides, and in turn, than ionic conductors [12]. This arises from the need of some amount of electric energy density (~10–50 W cm−3 ) to locally melt the particle contacts and form the liquid-film needed for current percolation. However, once the percolating liquid is formed and the particles rearranged by the liquid capillary forces, one expects the immediate solidification of the liquid. Solidification of the liquid-
262
R. Chaim / Scripta Materialia 178 (2020) 261–263
film as an exothermic reaction is thermodynamically preferred [6]. Despite the transient nature of the liquid-film, the densification and the resultant sintered microstructure of the ceramic particles depend on the liquid-film properties such as its molecular structure, its electric conductivity, and its temperature-dependent viscosity. Below we will further discuss these properties. One of the common features of different oxides subjected to flash sintering is the similar dissipated power at the flash onset, irrespective of the flash temperature value [13]. This is in accordance with the formation of the liquid film at the flash temperature, since all oxide liquids exhibit relatively high electronic conductivity through the percolating liquid-film. However, the oxide melts exhibit different properties, defined by their chemical composition, hence dictate the resultant microstructure of the flash sintered ceramic. It is necessary to refer to the oxide liquid phase structure and properties in order to understand the final microstructure evolved by the flash sintering. The internal structure of the oxide glasses, liquids, and supercooled liquids were the matter of continuous investigation, due to their complex nature. However, at present, several aspects of the glass structure are well-accepted paradigms in this field. Traditionally, selected cations with low coordination number (SiO2 , GeO2 , P2 O5 , B2 O3 and As2 O3 ) are classified as the network-formers, to say they form thermodynamically stable glass. These class of glasses reveal three-dimensional network of interconnected rings, the typical size of which is also present in their crystalline counterparts [14]. The melting points of the glass-forming oxides are significantly different from each other (SiO2 - 1883 K, GeO2 - 1388 K, P2 O5 - 853 K, B2 O3 - 723 K, and As2 O3 - 313 K). These differences in the melting points, together with differences in their liquid viscosity, were related to the density of the crosslinked bonds (defined as the number of the bridging bonds per cation less 2) in the glass network [14]. Therefore, assuming formation of a liquid-layer during the flash sintering, the melting point, and the elastic modulus which is associated with glass softening temperature [15], can act as good measures for the onset flash sintering temperature of these oxides. The viscosity of the network-former glasses near the glass transition temperature (Tg) follows an Arrhenian behavior, hence they were named as ’strong glasses’. The largest group of cations are conditional network-formers and cannot form continuous networks by themselves from their oxide melts. Therefore, their retention as a solid glass may need extremely high cooling rates, if possible at all. Accordingly, the viscosity of their supercooled liquids is nonArrhenian and drastically increases below Tg, hence their liquid is termed ’fragile’. The fragile behavior is related to the liquid’s internal structure that composes polyhedra clusters similar to those existing in their crystalline counterpart. Wang et al. [16] explored the cluster structure within the liquid in metallic glasses, using molecular dynamic simulations supported by the experimental alloy fabrication. Although the cluster type was important for the crystallization of the melt, they found strong correlation between the variance of the cluster distribution and the glass forming ability (GFA) of the melt. Large variance of the cluster distribution leads to difficulties in the crystallization of the melt due to a large difference between the cluster types, hence an increase in the glass forming ability of the liquid. A similar trend may be applicable to ceramic systems, albite the cluster types in the ionic melts are better defined due to the electrostatic bonding, and extremely high cooling rates are impractical due to the low thermal conductivities. Based on the fact that any liquid can be formed as solid glass when a critical cooling rate is applied, Jiusti et al. [17] analyzed the GFA of glass forming oxides. They related the GFA to the cooling rate in general, and defined it to be proportional to η (TL )/(TL ) 2 , i.e. the ratio between the melt viscosity at the liquidus temperature divided by the square of the liquidus
temperature. In this respect, viscous flow flash sintering results of silica glass by Prado et al. [fig. 4 in 18] show that the flash onset in glass under different applied fields is associated with reaching a certain electric conductivity (hence viscosity), rather than a certain dissipated power. In order to highlight the effect of the liquid film nanostructure on the flash temperature, we will use alumina as an example, the liquid of which is extremely poor glass former, and exhibits non-Arrhenian viscosity. Several flash sintering studies consistently exhibited lower flash onset temperatures for MgO-doped Al2 O3 compared to pure Al2 O3 powders [19–21]. MgO-doped (0.25 wt.%) 99.9% alumina with 200 nm average particle size (diameter) was flash sintered at 1268 °C under the field of 10 0 0 V cm−1 , while pure Al2 O3 powder failed to flash sinter at this field, resulting in arcing at 1400 °C [19]. Alumina powder with larger particle size (d50 = 600 nm) and higher impurity content (i.e. 99.8%) was flash sintered around 1050 °C at the same applied field [20]. For comparison, 99.5% pure MgAl2 O4 spinel powder with 250 nm particle size was successfully flash sintered at 1410 °C at the same electric field [22]. Therefore, we can evaluate the effect of Mg dopant on the liquid film properties based on the above discussions. Alumina melts were intensively investigated for their melt structure, their electric conductivity and viscosity in the supercooled condition. Ansell et al. [23] used Synchrotron x-ray diffraction for levitating alumina melts, and determined the (AlO4 )5− tetrahedra to be the main structural units, where the Al coordination (number) changes from octahedral (6) in the solid to tetrahedral (4) in the liquid. Similar experiments by Skinner et al. [24] with molecular dynamic and reverse Monte Carlo simulations showed the predominant structural units to be AlO4 and AlO5 in the 2:1 ratio, and minor fractions of the AlO3 and AlO6 molecules in the liquid. They showed that 83% of the Al-O-Al connections were involved in corner sharing polyhedra. Therefore, alumina melt can form glass in certain extreme conditions [23]. Levi et al. [25] used extremely high cooling rates of ~105 Ks − 1 to form amorphous alumina. Decrease in the cooling rate resulted in metastable transition (Gamma) alumina, where Al ions are both tetrahedral and octahedral coordinated. Cooling rates of alumina melt below 102 K s−1 led to the crystallization of the stable alphaalumina [26]. This melt behavior may partially explain the lack of a residual grain boundary amorphous layer in flash sintered pure aluminas. On the other hand, the formation of (AlO4 )5− clusters in liquid Al2 O3 is believed to be responsible for the ionic conductivity of 6 ± 1 [ cm]−1 at the melting point [27]. The electric conductivity of the molten oxides depends primarily on the cation’s charge, size, and mass [28,29]. In this respect, generally, the increase in the cation charge, cation radius and cation mass lead to decease in the electric conductivity. Cations with multiple oxidation states exhibit the lowest electric conductivity for their molten oxide with the highest valence state. Thus, the electric conductivity of molten oxide follows approximately the location of their cations in the periodic table of the elements. Consequently, molten MgO exhibits higher electric conductivity (i.e. 35 [cm]−1 [28,29]) than molten Al2 O3 (i.e. 15 [ cm]−1 [28]; 6 [ cm]−1 [27]; 0.2 [ cm]−1 [30]). It is believed that the higher electrical conductivity values were due to the impurity content in molten alumina [30]. Electric conductivity in molten oxides is also inversely proportional to the melt viscosity according to Walden’s rule [4]. This arises from the size of the ionic clusters present within the melt and the degree of the covalent bonds (i.e. degree of ionicity) between the clusters (molecules). The viscosity of molten MgO (at 2852 °C) was calculated as 1.4 [mPas] together with its specific electrical conductivity as 20.9 [ cm]−1 [31]. For comparison, viscosity of pure molten Al2 O3 was reported as 40 [mPas] at the melting point (2072 °C), and decreased to ~15 [mPa s] at 2800 °C
R. Chaim / Scripta Materialia 178 (2020) 261–263
[32]. Assuming the viscosity of molten mixed oxides to depend linearly on the viscosities of its oxide constituents in small composition ranges [32,33], addition of MgO is expected to increase the electrical conductivity of molten MgO-doped alumina and decrease its viscosity, whereas SiO2 dopant will act in the opposite way. Since the flash event is expected to take place by percolation of the electric current through the liquid-film layer, one expects higher current through the MgO-doped liquid-film than the pure alumina liquid-film. Consequently, the dissipated power needed to initiate and proceed the flash should be reached at relatively lower temperatures than for pure alumina. Furthermore, the lower viscosity of the liquid-film in the MgO-doped alumina should increase its flow kinetics by rapid wetting of the surrounding particles and their liquid-aided rearrangement and densification. On the other hand, SiO2 exhibits extremely higher viscosities and lower electric conductivities than Al2 O3 at temperatures around the alumina melting point [32]. Therefore, one expects a reduced effect of densification during the flash sintering in SiO2 -doped alumina than in pure alumina or MgO-doped alumina. Indeed, flash sintering of alumina containing 8 wt.% SiO2 and 2 wt.% MgO exhibited flash temperature of ~700 °C under the field of 1500 V cm−1 , but with very limited densification; linear shrinkage of 0.3% was measured compared to 5.6% in almost pure alumina under the same electric field [34]. The theoretical observations and expectations mentioned above are in full agreement with the experimental flash sintering results from the literature [1,19–21,34] and explain the possible effects of the impurities on the liquidlayer properties, hence on the flash sintering temperature and behavior. Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.11. 046.
263
References [1] M. Biesuz, V.M. Sglavo, Scr. Mater. 150 (2018) 82–86. [2] Y.V. Bykov, S.V. Egorov, A.G. Eremeev, V.V. Kholoptsev, K.I. Rybakov, A.A. Sorokin, J. Am. Ceram. Soc. 98 (2015) 3518–3524. [3] N. Morisaki, H. Yoshida, T. Tokunaga, K. Sasaki, T. Yamamoto, J. Am. Ceram. Soc. 100 (2017) 3851–3857. [4] R. Chaim, Materials (Basel) 9 (2016) 280. [5] R. Chaim, Materials (Basel) 10 (2017) 179. [6] R. Chaim, C. Estournès, J. Mater. Sci. 53 (2018) 6378–6389. [7] J. Narayan, Scr. Mater. 176 (2020) 117–121. [8] R. Chaim, G. Chevallier, A. Weibel, C. Estournes, J. Appl. Phys. 121 (2017) 145103. [9] R. Chaim, Scr. Mater. 158 (2019) 88–90. [10] S.V. Egorov, A.G. Eremeev, V.V. Kholoptsev, I.V. Rybakov, A.A. Sorokin, Y. Bykov, Scr. Mater. 174 (2020) 68–71. [11] R. Chaim, Scr. Mater. 169 (2019) 6–8. [12] R. Chaim, C. Estournès, Scr. Mater. 163 (2019) 130–132. [13] R. Raj, J. Am. Ceram. Soc. 99 (2016) 3226–3232. [14] B. Bridge, N.D. Path, D.N. Waters, Phys. Stat. Sol. (a) 77 (1983) 655–668. [15] B.D. Sanditov, S. Sh. Sangadiev, D.S. Sanditov, J. Exper. Theor. Phys. 115 (2012) 445–454. [16] J. Wang, A. Agrawal, K. Flores, Acta Mater. 171 (2019) 163–169. [17] J. Jiusti, E.D. Zanotto, D.R. Cassar, M.R.B. Andreeta, J. Am. Ceram. Soc. (2019), doi:10.1111/jace.16732. [18] M.O. Prado, M. Biesuz, M. Frasnelli, F.E. Benedetto, V.M. Sglavo, J. Non-Cryst. Solids 476 (2017) 60–66. [19] M. Cologna, J.S.C. Francis, R. Raj, J. Eur. Ceram. Soc. 31 (2011) 2827–2837. [20] M. Biesuz, V.M. Sglavo, J. Eur. Ceram. Soc. 36 (2016) 2535–2542. [21] M. Biesuz, P. Luchi, A. Quaranta, V.M. Sglavo, J. Appl. Phys. 120 (2016) 145107. [22] H. Yoshida, P. Biswas, R. Johnson, M.K. Mohan, J. Am. Ceram. Soc. 100 (2017) 554–562. [23] S. Ansell, S. Krishnan, J.K. Richard Weber, J.J. Felton, P.C. Nordine, M.A. Beno, D.L. Price, M.-.L. Saboungi, Phys. Rev. Let. 78 (1997) 464–466. [24] L.B. Skinner, A.C. Barnes, P.S. Salmon, L. Hennet, H.E. Fischer, C.J. Benmore, S. Kohara, J.K.R. Weber, A. Bytchkov, M.C. Wilding, J.B. Parise, T.O. Farmer, I. Pozdnyakova, S.K. Tumber, K. Ohara, Phys. Rev. B 87 (2013) 024201. [25] C.G. Levi, V. Jayaram, J.J. Valencia, R. Mehrabian, J. Mater. Res. 5 (1988) 969–983. [26] J.K.R. Weber, C.D. Anderson, D.R. Merkley, P.C. Nordine, J. Am. Ceram. Soc. 78 (1995) 577–582. [27] J.E. Enderby, S. Ansell, S. Krishnan, D.L. Price, M.-L. Saboungi, Appl. Phys. Lett. 71 (1997) 116–118. [28] A.E. van Arkel, E.A. Flood, N.F.H. Bright, Canadian J. Chemistry 31 (1953) 1009–1019. [29] E. Thibodeau, I.-.H. Jung, Metall. Mater. Trans. B 47 (2016) 355–383. [30] A.N. Shatunov, A.I. Maksimov, A.Y. Pechenkov, I.V. Poznyak, Inorg. Mater. 47 (2011) 1579–1583. [31] A.-.L. Leu, S.-.M. Ma, H. Eyring, Proc. Nat. Acad. Sci. USA 72 (1975) 1026–1030. [32] G. Wu, E. Yazhenskikh, K. Hack, E. Wosch, M. Müller, Fuel Processing Tech. 137 (2015) 93–103. [33] G.-.H. Zhang, K.-.C. Chou, Metall. Mater. Trans. B 43B (2012) 849–855. [34] M. Biesuz, V.M. Sglavo, J. Eur. Ceram. Soc. 37 (2017) 705–713.