Combination of SHS and Mechanochemical Synthesis for Nanopowder Technologies

The combination of mechanochemical activation and self-propagating high-temperature synthesis (SHS) have widened the possibilities for both methods. For metallic systems, the investigation showed that a short-term mechanochemical activation of heterophase SHS products leads to singlephase and ultrafine intermetallides obtained from the elements by mechanical alloying. It was demonstrated that metastable phases, usually obtained by mechanical alloying, can be obtained from the equilibrium intermetallic compounds synthesized by SHS. Besides this, the investigations showed that preliminary mechanical activation, during which layered composites are formed from the initial elements, allows one to extend the concentration limits of SHS processes up to a solid solution


Introduction
Mechanochemical synthesis and mechanical alloying are in wide use as experimental methods for the production of highly dispersed powders and nanocomposites. Using mechanical activation we have managed to prepare nanocomposites in immiscible systems such as Cu-Bi, Fe-Bi, Fe-In. It is possible to get supersaturated solid solutions in many other intermetallic systems (Cu-Ga, Cu-In, Cu-Sn, Ni-In, Ni-Bi, Ni-Sn, Ni-Ge, Ni-Al) [1,2]. However, from a technological point of view, mechanical activation poses a lot of problems which cannot be easily solved. The biggest problem is the low productivity of the techniques currently available, as well as the contamination of the end products caused by abrasion of the grinding media. The energy consumption should also be taken into consideration in some large-scale processes.
One of the alternatives to mechanochemistr y may be self-propagating high-temperature synthesis (SHS), which is energy-saving and can be applied to the large-scale production of many intermetallic compounds or complex oxides (nitrides, carbides) [3]. Nevertheless, this method includes a combustion stage which requires very high temperatures. As a result, final products can be obtained only in the form of dense sintered or solidified products when SHS exceeds the melting point of reagents and/or products. To transform such products into commercially interesting powders, one needs to use milling as an unavoidable step. The grinding leads to further contamination and to additional energy consumption, which may be comparable with that required for the mechanochemical synthesis itself.
Another disadvantage of SHS is a relatively high temperature gradient near the combustion surface which leads to a non-uniform phase composition of the obtained products.
The shortcomings of these two methods can be overcome, and their advantages can be extended. This allows one to obtain the final products of required phase composition and disperse state, thus enabling the development of a new energy-saving and waste-free technology.

Experimental
The materials used in the investigation were: carbonyl nickel and iron, powdered aluminum, silicon, germanium, tungsten, molybdenum, titanium, and zirconium, pure gallium, barium peroxide, tungsten and molybdenum oxides.
A ball planetary mill AGO-2 [4] was used for the investigations. The volume of the mill drums was 250 cm 3 , the ball diameter was 5 mm, the ball load was 200 g and the weighed portion of powder treated was 10 g. In order to avoid oxidation of the metals, all the experiments on mechanical alloying were carried out in an argon atmosphere. The X-ray phase analysis was performed with a DRON-3M diffractometer with CuK α radiation.
IR absorption spectra were recorded with a SPECORD 75 IR spectrometer.
Electron microscopic studies were carried out using the JSM-T20 electron microscope and the highresolution electron microscopes JEM-2010 and JEM-400.
The extent of aluminum recovery from intermetallic compounds was calculated as the ratio of the amount of hydrogen evolved in the reaction of samples with a 20% κOH solution to the amount of hydrogen that should be evolved in case of complete dissolution of the aluminum present in the sample. The amount of the evolved hydrogen was recorded automatically with a DAGV-70-2M volumeter.

Homogenizing ef fect of MA
Ultrafine single-phase intermetallic compounds are widely used in powder metallurgy for making heatresistant materials, materials with a high corrosion stability and special magnetic properties. Singlephase intermetallic compounds are usually obtained by fusing or sintering followed by homogenizing and annealing for a long time [5]. The major disadvantages of these processes are deviations in the composition of the mixture in the case of high melting points of the initial components because of evaporation of the lower melting point ingredients. As a result, a long time interval is necessary for the process, and high energy consumption at the stage of homogenizing and annealing. Besides this, most of the powder intermetallides are obtained by grinding for a rather long time, which usually leads to substantial contamination caused by abrasion of the grinding media. In such cases, additional purification is required.
For metal systems with high enthalpies of the intermetallide formation, the SHS method can be applied; heterophase SHS products can be homogenized during mechanochemical activation at a simultaneous substantial decrease of the particle size, which should finally lead to finely dispersed material with a high structural defect content. This material is similar in its properties to that obtained by mechanical alloying of the elements; however, the time for mechanical treatment in the former case is several decades shorter.
Phase formation and morphological changes were investigated for the mechanochemical synthesis of intermetallides from elements and for the mechanochemical activation of heterophase products of the SHS process. For comparison purposes, systems were considered for which it is known that the SHS process can be performed, but the products of this synthesis are substantially different in phase composition: -Ni-Al within the concentration region of NiAl phase existence. It is known that this phase can be obtained by SHS [3]; -Ni-Si, in which the SHS process is possible [6], but it is difficult to achieve single-phase composition during SHS synthesis; -Ni-Ge in the NiGe intermetallide homogeneity region. Due to the rather high melting point of Ni (1455°C) and the rather low temperature of NiGe intermetallide decomposition (850°C), the homogeneity region is very narrow for this compound. The X-ray investigation of the mechanical alloying of an NiAl intermetallide from a mixture of metal powders of the composition Ni 32 wt.% Al demonstrated that the synthesis started after mechanical activation for about 2.5Ҁ3 minutes. First, the Ni 2 Al 3 intermetallide forms. Some amounts of unreacted nickel and aluminum still remain (Fig. 1a). After activation for 5Ҁ7 min., the amount of the formed Ni 2 Al 3 phase starts to decrease and the ref lections of the NiAl intermetallide increase (Fig. 1b). Then, the Ni 2 Al 3 intermetallide disappears, the nickel and aluminum contents decrease gradually and the intensities of NiAl ref lections increase. After mechanochemical activation for 25Ҁ30 min., very broad ref lections only of the NiAl phase are present in the X-ray diffraction patterns (Fig. 1c). The coherent scattering domains for this product are 8Ҁ10 nm. The electron microscopic investigation of the products of the mechanical alloying of NiAl-based solid solutions in a ball mill demonstrated that after activation for 30 s, the initial nickel and aluminum particles could not be detected in the sample. The products consist of agglomerates of various shapes and sizes. An increase of the activation time to 1 min. leads to an increase of the fraction of coarse agglomerates which attain a plate-like shape as their density increases. Figure 2a shows a microphotograph of the cross section of the agglomerate, clearly exhibiting its layered structure. A further increase to the time of mechanical alloying causes a drastic change to the morphology of the product particles. The destruction of large layered agglomerates starts after mechanical alloying for 3 min.; whereas they disappear completely after 5 min. of alloying. The major part of the samples by this time is composed of rather dense agglomerates of irregular shape with a size of 2Ȗ10 µm. Larger particles (up to 50 µm) are also present. These agglomerates are composed of the layered particles 1Ҁ5 µm in size. It should be noted that according to XRD data, it is the mechanical alloying for 3 min. that causes the appearance of the Ni 2 Al 3 intermetallide in the products, which is likely to be the reason of the significant change of the particle morphology. An increase in the time of mechanical alloying brings insignificant changes to the morphology of the formed products. The fine fraction (1Ҁ5 µm) content increases, the density of the formed particles increases (Fig. 2b).
The SHS of a metal powder mixture comprising Ni 32 wt.% Al results in the NiAl intermetallide with admixture of the Ni 2 Al 3 phase; judging from the width of diffraction peaks, the formed products are well crystallized. Electron microscopic studies showed that the intermetallides were most likely crystallized from the melt (Fig. 3a).
Mechanochemical activation of SHS products for 1.5 min. results in a material similar in morphology, particle size and coherent scattering domain to the material obtained by the mechanical alloying of nickel and aluminum powders for 25 min. This means that a short-time mechanochemical activation of a mixture of phases obtained by the SHS process in this system allows preparation of an ultra-fine single-phase intermetallic product (Fig. 3b). Similar products were observed for the Ni-Si system (Fig. 4).
X-ray analysis showed that a significant amount of the Ni 3 Si 2 phase can be found in the products of the mechanical alloying of an NiѿSi mixture corresponding to an Ni 5 Si 2 intermetallic compound.
The X-ray diffraction investigation of the mechanical alloying of nickel and germanium at the ratio of Ni 55 wt.% Ge demonstrated that formation of the NiGe intermetallide starts within the first minute (Fig. 5a); a drastic increase in the intensity of the ref lections from this phase starts after mechanical activation for 6Ҁ7 minutes (Fig. 5b); the maximum is achieved after 25 minutes of mechanical treatment (Fig. 5c). However, an admixture of the second phase (Ni 5 Ge 2 ) appears during mechanochemical synthesis of the NiGe intermetallide; at first, the content of the second phase increases, then decreases and disappears completely after 12Ҁ13 minutes of mechanical alloying. By the end of the process (25Ҁ30 minutes) there is only one phase, which is the NiGe intermetallide, with a coherent scattering domain of 5Ҁ8 nm.
According to the data of the electron microscopic investigation, mechanical alloying in the mixture of a stoichiometric composition NiGe for 40 s results in the formation of rounded three-dimensional agglomerates with a maximal size of up to 400 µm, originating from particles of 0.2Ҁ1 µm in diameter. Besides this, separate particles of 0.2Ҁ1 µm in diameter are present; their number is rather large. An increase of the mechanical alloying time for this system to 5Ҁ7 minutes leads to the formation of very dense agglomerates; traces of a strong plastic deformation are sometimes observed on their surface. However, in this case the amount of separate fine particles is also rather large. A further increase of the mechanical alloying time to 30 min. has practically no effect on the morphology and size of the resulting particles; it is only the fraction of coarse agglomerates which increases. The final product of mechanochemical synthesis NiGe is composed of rather dense particles with a size of 0.5Ҁ2 µm; they sometimes form larger agglomerates.
An SHS process was carried out in a mixture of this composition. According to XRD data, it is mainly the equilibrium well-crystallized NiGe which forms, with an admixture of a second intermetallide (Ni 5 Ge 2 ) (Fig. 6a). One can see from microphotographs that the final products are formed from the melt (Fig. 7a), which means that the SHS proceeds via the liquid phase. The X-ray diffraction investigations of the mechanochemical treatment of the SHS products demonstrate that a broadening of the diffraction patterns starts after mechanical activation for 30 s. Phase homogenization of the system (NiGe) is achieved after 1.5Ҁ2 minutes (Fig. 6b). After mechanical activation for 3 min., the material exhibits a coherent scattering domain of 8Ҁ10 nm. Further mechanical activation leads to some narrowing of the peaks. According to the electron microscopic data, after mechanical activation of the SHS products for 30 s, the major part of the sample is composed of fine particles of irregular shapes, their size being 0.2Ҁ3 µm (Fig. 7b). Larger particles are still present, though they rarely occur in this case. A sintered agglomerate of fine particles starts to form. An increase of the treatment time causes an intensive formation of such agglomerates. Their density increases gradually. Traces of substantial plastic deformation are observed on the surface of some of these agglomerates. After mechanical activation for 3 min., the sample consists of agglomerates of different density and shape, their size being 1 to 400 µm.
Similar results were obtained for the Ni-Si system (Fig. 4a, 4b).
The investigation of the metallic systems showed that a short-term mechanochemical activation of heterophase SHS products leads to single-phase and ultrafine intermetallides obtained from the elements by mechanical alloying.

Formation of non-equilibrium phases by activation of SHS products
It is known that it is practically impossible to obtain non-equilibrium intermetallic compounds by SHS; however, highly reactive intermetallic compounds (which is exactly what metastable phases usually are) are used in hydrogen energetics, in preparation of Raney catalysts, metal cements, diffusion-hardening solder, metal dental materials, etc. Mechanochemical synthesis and mechanical activation are among the most efficient methods to obtain metastable phases in metal systems; it is known that a series of equilibrium intermetallic compounds can be transferred into nonequilibrium concentration regions by means of mechanical activation.
A comparative investigation of phase and microstructural transformations during the mechanochemical synthesis of non-equilibrium solid solutions from the elements and during mechanical activation of NiGe equilibrium intermetallic compounds obtained by SHS was carried out. The Ni-Al system was selected for investigation, because it is known that the SHS process in the Ni-Al system can be carried out within a broad concentration range [3]. The limit solubility of aluminum in nickel ranges from 3.85 wt.% at 500°C to 10.3 wt.% at 1385°C. The structural similarity of the Ni 3 Al intermetallide (cubic lattice of the Cu 3 Au type) and βnickel makes it possible to mechanochemically pass from the intermetallide to the non-equilibrium solid solution [7]. The structure of the Ni 2 Al 3 intermetallide (40.8 wt.% Al) is close to that of NiAl. Aluminum atoms form distorted cubes in the rhombohedral lattice of Ni 2 Al 3 . Two-thirds of the positions in the centres of cubes are occupied by nickel atoms, the other positions are vacant. Disordering of these vacancies should lead to the formation of a supersaturated solid solution with a non-equilibrium concentration of vacancies.
The X-ray and electron microscopic studies of the mechanical alloying in a mixture of nickel and aluminum powders at a ratio of Ni 40.8 wt.% Al (calculated for the formation of the Ni 2 Al 3 intermetallide) showed that the process is largely similar to the mechanical alloying of the Ni 32 wt.% Al mixture. At first, the Ni 2 Al 3 intermetallide is formed; during further activation it is transformed into the supersaturated solid solution with a non-equilibrium concentration of vacancies based on the intermetallic compound NiAl. After mechanical activation for 30 minutes, the size of its coherent scattering domains is around 8 to10 nm.
The X-ray diffraction patterns of the product of the SHS process in the Ni 40.8 wt.% Al mixture exhibit the formation of a well-crystallized Ni 2 Al 3 intermetallide (Fig. 8a). Mechanochemical activation of this product for ȁ2 min., under the same conditions as those under which mechanical alloying was carried out, gives the diffraction patterns corresponding to the material obtained by mechanical alloying after 25Ҁ30 minutes (Fig. 8b), which has been mentioned above to be a supersaturated solid solution with a non-equilibrium concentration of vacancies based on the NiAl intermetallide. Further mechanical activation does not bring any substantial changes.
Electron microscopic investigations showed that the formed porous fused SHS product (Fig. 9a) can be transformed mechanochemically within 2 min. into a product which is morphologically very similar to that formed at the final stage of mechanical alloying (Fig. 9b).
Metastable phases obtained by mechanical alloying and by mechanical activation of the SHS products have approximately equal reactivity.
For the example of a well-studied aluminum leach- ing reaction, it was stated that non-equilibrium solid solutions based on NiAl, which were synthesized by mechanical alloying from metal powders within the concentration range of the Ni 2 Al 3 intermetallide and activated for 1 min., exhibit a very high reactivity (Fig. 10). These samples exhibit similar dynamics to hydrogen evolution in leaching. This is one more confirmation of the fact that the formed metastable phases are identical. This means that the non-equilibrium solid solutions based on the NiAl intermetallide, with largely similar structure, particle morphology, coherent scattering domains, and reactivity can be obtained both by the mechanical alloying of the initial powders of the composition Ni 40.8 wt.% Al, and also by short-time mechanical activation of the Ni 2 Al 3 intermetallide obtained by SHS. An important fact! The X-ray diffraction studies of the mechanical alloying of the Ni 13.5 wt.% Al composition (the concentration region of equilibrium Ni 3 Al intermetallide) showed that a broadening and a slight decrease of the intensity of diffraction peaks related to metals occur within the first two minutes of activation. Diffraction ref lections of the Ni 2 Al 3 intermetallide appear after 2.5 minutes of mechanical treatment. The intensity of aluminum peaks decreases drastically while the intensity of Ni ref lections decreases slowly. The Ni lattice parameter remains unchanged. Only after mechanical activation for 5 min. do the diffraction patterns of aluminum disappear completely. The diffraction ref lections of nickel broaden substantially (which is especially noticeable in the large-angle region). Its lattice parameter increases (Fig. 11a), which is evidence of the incipient formation of a nickel-based solid solution. After 15 minutes of activation, the lattice parameter of the solid solution reaches its maximum a҃0.3590 nm. The coherent scattering domain of the resulting solid solution is 8Ҁ10 nm. Further mechanical activation causes a narrowing of the diffraction peaks. The formation of the Ni 3 Al phase, which is characteristic in this concentration range, was not detected at any stage of activation.
Electron microscopic investigations of the products of mechanical alloying in the mixture Ni 13.5 wt.% Al showed that the dynamics of particle morphology changes at the initial stages of the process is very similar to that observed for the composition Ni 32 wt.% Al. The formation of layered composites was observed, too. However, in this case they start to form as early as after only 30 s of mechanical alloying, which is much earlier than for the Ni 32 wt.% Al composition. The particles of the initial components are no longer detected at this moment. A substantial part of the resulting agglomerates look like plate-shaped particles with traces of strong plastic deformation on their surface. Relatively sintered and round-shaped agglomerates are also present. The maximal size of the plate-like agglomerates is 1Ҁ2 mm, while the rounded ones are between 1Ҁ2 and 40Ҁ50 µm in size. An increase of the mechanical treatment time to 15 min. leads to a decrease of the fraction of the smallest and the largest particles; they become more dense and uniform in size, almost monolithic.
A product of SHS in the mixture Ni 13.5 wt.% Al is an Ni 3 Al intermetallide. Diffraction patterns are evidence of a high order in the structure of the formed phase (Fig. 11b) of the intensity of diffraction peaks and to their substantial broadening. After 1.5 min., a non-equilibrium solid solution of aluminum in nickel is formed. Its lattice parameters are identical to those of the product formed in a mixture of initial powders mechanically alloyed for 15 min. (Fig. 11c). Further activation of this SHS product causes no change in diffraction patterns. It was stated in electron microscopic studies that the particle size decreases sharply at the very first stages of the mechanical treatment of SHS products. After mechanical activation for 1.5 min., the product particles are practically identical in size and morphology to the products obtained from the initial Ni and Al powders after mechanical alloying for 15 min.
The investigation demonstrated that metastable phases, usually produced by mechanical alloying, can be obtained from the equilibrium intermetallic compounds synthesized by SHS.
Short-term activation following the SHS process therefore allows one to homogenize intermetallic phases and to obtain metastable structures. The final products are formed in the ultrafine state with minimal contamination caused by abrasion of milling tools.
This approach can also be used in the SHS synthesis of complex oxides, nitrides, borides, and carbides, etc. However, for this method of preparing ultrafine powders, the range of compounds is limited by the possibilities of the SHS method, i.e. involved are systems with very high temperatures of the final product formation.

Metal systems
A high exothermal effect of reaction and a strong Arrhenius-type temperature dependence of the reaction rate are of decisive importance for the feasibility of SHS processes. For diffusion-controlled processes, to which SHS processes also belong, another important parameter is the disperse state of the initial components, including the uniformity of their mixing and the surface area of contacts between the components. Mechanical alloying is known as a process which leads to the formation of layered mechanocomposites formed at the initial stage of the mechanical activation of metal mixtures (see Fig. 2a). Most metals exhibit good plasticity; under mechanical action in ball planetary mills where shock and shock-with-shear occur [8], the metal undergoes plastic deformation resulting in a change to the shape of the metal particles. They f latten, adsorption films on their surface are destroyed; components come into contact with each other by atomic-pure planes [9]. ‹‹Point›› contacts of initial particles transform into f lat ones, while the contact area increases considerably [10]. As a result, a short-term preliminary activation can extend the possibilities of reaction in a self-propagating regime even for systems where the SHS process cannot be carried out without preliminary heating.
In order to establish the lower limit of the conditions for the SHS process to occur with mechanical activation involved, we investigated the Ni-Al, Ni-Si, Ni-Ge systems, which are SHS systems and for which the lower concentration limits for silicon and aluminum in SHS without preliminary mechanical activation are known [11]. In the Ni-Al system, the limiting solid-phase solubility of aluminum in nickel is 10.3 wt.% at 1385°C and decreases to 3.85 wt.% at 500°C, but SHS cannot be performed in this concentration range. The minimal aluminum concentration at which SHS can be performed by the traditional method is 13 wt.% Al [3].
Preliminary mechanical activation in the Ni-Al system allows one to decrease the minimal content of Al in the initial mixture for the SHS process. This minimal Al content is 7 wt.%. We succeeded in performing SHS in the sample compacted from the powder after mechanochemical treatment. The Ni 3 Al, Ni 2 Al 3 intermetallides and unreacted nickel were detected in the synthesis products. Subsequent mechanochemical activation homogenizes the product and leads to the formation of a solid solution only.
In the Ni-Si system, similarly to the Ni-Al system, the concentration range of the solid solution is strongly dependent on temperature; at room temperature, the solubility of silicon in nickel is ȁ5 wt.%, while at 1150°C it is ȁ9.3 wt.% [12,13]. It is not possible to realize an SHS process in the solid solution range without preliminary mechanical activation in this system.
Our investigations showed that in the solid solution range, the layered composites of nickel and silicon could be formed under activation within 1 min., though they do not look so dense as in the Ni-Al system. The SHS process was also performed in these mixtures containing 9 wt.% Si, the products being a mixture of the Ni 2 Si, Ni 5 Si 2 intermetallides and unreacted nickel. Similarly to the case of the Ni-Al system, subsequent mechanical activation results in a singlephase product which is a supersaturated solid solution.
In the Ni-Ge system, in which the limiting solubility of germanium in nickel is 13.8 wt.% [5,14], we investigated the possibility of performing SHS also in the solid solution concentration range. Composites are formed in a Ni 13 wt.% Ge mixture after mechanical activation for 1.5 min. The Ni 5 Ge 2 and NiGe intermetallides and unreacted nickel are the products of SHS. The formation of a single-phase solid solution is achieved by a subsequent short-term activation.
The investigations therefore showed that preliminary mechanical activation, during which layered composites are formed from the initial elements, allows one to broaden the concentration limits of SHS processes up to the solid solution region, perhaps due to reagent dispersion and an increase of contact area. The products of this synthesis are mixtures of intermetallides, the doping elements being completely consumed for their formation, and an excess of solvent metal. The conservation of the layered structure in SHS products allows us to assume that intermetallic compounds are formed locally inside the layered structure, its framework being built of solvent metal.
Using the discovered effect of preliminary mechanical activation on the possibility of performing SHS even in the concentration range of solid solutions, we tested it in the systems in which SHS cannot be performed without preliminary heating of the reaction mixture even in the range of intermetallide existence [3]. As an example, Figure 12 shows the combustion rate and the maximum combustion temperature plotted versus time of the preliminary mechanical activation for the Ni 45 wt.% Ti composition. The samples start burning at room temperature after 2.5 min. of mechanical activation. Similarly to all the systems considered above, an increase of the burning rate is connected to the formation of dense layered composites. A photograph of the sample after mechanical activation for 2.5 min. is shown in Figure 13. Increasing the time of mechanical activation also leads to a further increase of the density of the composites, a decrease of the grain size to 0.1 µm, and to the appearance of particles with traces of strong plastic deformation on the surface. The decrease of the combustion rate is connected to the start of destruction of the largest composites. Only NiTi (the main phase) and Ti 2 Ni lines are observed in the diffraction patterns of the SHS products at any time of mechanical activation. No titanium or nickel lines are observed! The relative content of these phases in the products remains practically independent of the time of mechanical activation.

Oxide systems
The synthesis of nanocrystalline mixed oxides is among the major problems of advanced ceramic technology. Conventional processes for manufacturing multi component mixed-oxide ceramics involve hightemperature reactions between metal oxides or between metal oxides and carbonates. These diffusion-controlled processes require the application of high temperatures and the use of highly dispersed precursor powders; mixed oxides are formed in particles of 1 to 5 µm in size. Mechanochemical alloying partly removes the diffusion control.  able heat of the chemical reaction provides the high rate of the mechanochemical reaction. For example, for most reactions of mechanochemically assisted metal oxidation in a gaseous phase, the oxidation rate correlates with variations in Gibbs' free energy [15]. For the exothermal reactions, the mechanochemical reaction rate of solid solution formation is a function of the enthalpy of formation of intermetallic compounds in equilibrium [16]. Therefore, if reactions with small decreases in Gibbs' free energy are used in thermodynamically controlled mechanochemical synthesis, a large power supply and a long mechanical activation time would be necessary, by analogy with mechanochemical synthesis of mixed oxides from binary oxides, either alone or mixed with carbonates [17]. Efficient mechanochemical synthesis is not feasible unless in energy-intensive activators (steel drums and steel balls), and contamination of the product is inevitable. Therefore, mechanochemistry is almost useless for commercial ceramic processes with their extremely high purity requirements. The mechanochemical approach becomes practicable only with high rates of promoted reactions, i.e., those that are not only thermodynamically allowed, but that also give an energy gain. High rates of mechanochemical processes result in final products which are highly dispersed, and this can significantly inf luence the properties of the resultant ceramics. The change of Gibbs' energy in the synthesis of complex oxides from simple ones shows that all these reactions are thermodynamically allowed, while the majority of reactions with the participation of carbonates are available only at high temperatures ( Table 1).
Some of these reactions were performed mechanochemically when a sufficient amount of energy was applied [17]. All the values obtained for the systems under consideration are approximately of the same order of magnitude, so the conditions for these reactions to proceed should be rather similar.
The use of peroxide compounds and oxides as the initial components of mixtures, as proposed by some authors [17], does not bring substantial changes to the ∆G of the reaction of complex oxide synthesis.
Among the methods to obtain oxides, the energetically most profitable ones are the direct oxidation of metals (or metal mixtures) by oxygen ( Table 2). However, in reality one can hardly perform this type of synthesis mechanochemically due to the high plasticity of metals.
The most promising way seems to be the oxidation of metals by peroxide compounds, especially if we take into account that there is a rather large class of stable metal peroxides that can provide a couple for a proper metal to synthesize a complex oxide. In such cases, one should keep in mind that the mechanochemical interaction of metal peroxide with metals can follow two routes: either with the formation of complex oxides or with the formation of a mixture of KONA Table 1 Changes of Gibbs' energy in the reactions of the mixtures of metal oxides and the metal oxides with metal carbonates resulting in the formation of complex oxides. simple ones, because both reactions are profitable from the thermodynamic viewpoint, and the change of Gibbs' energy is much higher than that for the synthesis from oxides and carbonates. The calculation of ∆G o 298 for the BaO 2 interaction with metals demonstrated that the decrease of Gibbs' energy was larger by 30Ҁ40 kcal/mol than when the products of the same reaction would be a sum of simple oxides. One can assume that synthesis will proceed to the formation of complex oxides.
According to X-ray diffraction data, the mechanochemical interaction of barium peroxide with titanium for 5 min. results in the formation of a mixture of barium titanates. The ref lections corresponding to BaO 2 and Ti disappear and ref lections corresponding to Ba 2 TiO 4 , BaTiO 3 [18], etc. appear. No diffraction ref lections from simple oxides of barium and titanium are observed.
The complex oxides were obtained for the system BaO 2 ѿMe (Me҃Zr, Al).
For the BaO 2 -Ti mixture, an antisymmetric absorption band peaking at 700 cm Ҁ1 and a very weak peak at 775 cm Ҁ1 appear after 1 min. of activation. After 5 min. of activation, a shoulder of the 700 cm Ҁ1 band appears at 550 cm Ҁ1 (Fig. 14b). This band is assignable [21-23] to ν (Ti-O) stretches in [Ti0 4 ] tetrahedra of barium titanates [23]. The positions of the peaks and the band shape do not correspond to ν (Ti-O) vibrations in various TiO 2 polymorphs [22,24].
The BaO 2 -Zr mixture mechanically activated for 5 min. exhibits an absorption band with a peak at 550 cm Ҁ1 (Fig. 14c). This band can be assigned to barium zirconate [21,23].
A broad asymmetric band with a maximum at 530 cm Ҁ1 appears as a result of activation of a mixture of BaO 2 with aluminum ( Fig. 14d). Magnesium aluminate exhibits a similar IR spectrum in this region [21,26,27], and the formation of barium aluminate can be assumed in the BaO 2 -Al system under mechanical activation [28].
The electron microscopic investigation of the product of mechanochemical synthesis in a mixture of BaO 2 with Zr shows that activation for 1 min. results in the formation of particles of 0.3Ҁ1 µm in size. They are composed of small blocks of 6Ҁ12 nm in size (Fig. 15). The microdiffraction picture obtained from a separate particle also points to a developed microblock structure. The diffraction spots are shaped as rings, which is characteristic of polydisperse materials, while microdiffraction from a separate block gives a point, which is evidence of the single crystal state of the substance. Diffraction ref lections in both electron diffraction patterns, though somewhat broadened, are point ref lections, which is evidence of a rather high degree of crystallinity of the substance formed in mechanochemical synthesis. A similar picture was also observed for other complex oxides, with only a small difference in microblock size.
The investigations therefore demonstrate that the mechanochemical interaction of barium peroxide with 154 KONA No.20 (2002)  metals, which proceeds with a substantial decrease of Gibbs' free energy, as it follows from thermodynamic calculations, allows one to synthesize complex oxides with nanocrystalline particles in a relatively short time.
The mechanochemical oxidation of metals by peroxide compounds allows the preparation of singlephase particles of complex oxides with nanometersized microblocks [29][30][31][32]. However, in the case of some metals such as tungsten, molybdenum, and tantalum, mechanochemical synthesis does not proceed to completion, whatever reagent ratio is taken, even after prolonged mechanical activation. Some part of the metal always remains unreacted. At the reagents molar ratio of BaO 2 :W҃1:1, the products are BaWO 4 and Ba 2 WO 5 ; unreacted W remains. At an increased barium peroxide content (up to 2:1), part of W also remains unreacted. A mixture of tungstates BaWO 4 , Ba 2 WO 5 and Ba 3 WO 6 is formed. At the reagents ratio of 3:1, only the content of complex oxide Ba 3 WO 6 increases. Mechanochemical activation of the mixtures of BaO 2 with Mo, the molybdates BaMoO 4 , Ba 2 Mo 5 and Ba 3 MoO 6 are formed. Similarly to the case of tungsten, a part of molybdenum remains unreacted.
The SHS process in these mixtures also results in the formation of a mixture of complex oxide phases; unreacted metal remains. The preliminary mechanical activation of the mixture has practically no effect on the phase composition of the SHS products although it does increase the process rate.
For BaO 2 -Me (Me҃W, Mo, Ta) systems, it was not possible to synthesize single-phase complex oxides, neither by mechanical activation nor by SHS.
IR spectroscopic investigations of the mechanochemical interaction of BaO 2 with WO 2 showed that the formation of a complex oxide starts within the first seconds of activation. The IR spectra of the initial mixture BaO 2 ѿWO 2 contain one broad band at 850Ҁ550 cm Ҁ1 , without clearly exhibited maximums. It is assigned to the stretching vibrations ν (W-O) (Fig. 16a) [21]. The ν (Ba-O) band is below 400 cm Ҁ1 . After activation for 10 s, a clear band with a maximum at 810 cm Ҁ1 is observed instead of the above-mentioned broad band; the intensity of this new band increases with increasing activation time (Fig. 16b, c). The similarity of the IR spectra of activated mixtures to the spectra of stolzite [21] allows us to assume that the activation of the BaO 2 ѿWO 2 mixture results in the formation of BaWO 4 with a spinel structure. However, according to the IR spectroscopic data, the mechanochemical reaction of BaO 2 ѿWO 2 → BaWO 4 does not proceed to completion, which is evidenced by the presence of noticeable absorption in the region 800Ҁ500 cm Ҁ1 as a shoulder of the band with a maximum at 810 cm Ҁ1 related to ν (W-O) of the lower tungsten oxide.
According to the XRD data, a growth of the ref lections from the BaWO 4 phase starts at the second minute of activation and reaches its maximum by 5 minutes (Fig. 17 a, b). But, with increasing activation time, the intensity of the diffraction peaks of the complex oxide does not increase, and the intensities of the peaks related to the initial tungsten oxide do not decrease. This means that the result of mechanochemical interaction between BaO 2 ѿWO 2 is a mixture of phases.
For the interaction of BaO 2 with MoO 2 , the IR spectra and XRD also reveal a mixture of phases.
The relatively high temperatures of formation of the complex oxides in the system involving the oxidation of the lowest tungsten oxide with barium peroxide allows one to perform these reactions by means of SHS. However, pretreatment of the barium peroxide and sometimes heating of the initial mixture are necessary [33]. This is due to the fact that BaO 2 particles entrained in air become coated with a layer of barium carbonate and hydroxide. This layer prevents an SHS reaction. One can assume that a short-term prelimi- nary mechanochemical activation of the initial mixture leads to the destruction of these barrier layers and provides a substantial increase of the area of contact between the oxides and the barium peroxide. An investigation of the effect of the preliminary mechanochemical activation of commercially available BaO 2 with WO 2 showed that the maximal temperature and rate of SHS process are achieved after activation for 2 min., and the single-phase complex oxide BaWO 4 is formed (Figs. 16d and 17c). A disadvantage of the resulting substance was the large particle size (300Ҁ500 nm) and in some places even partial agglomeration (Fig. 18a). Electron microscopic studies show that subsequent mechanical treatment of the product for 2 minutes in a highenergy activator of planetary type produces a material with a particle size of 20Ҁ30 nm (Fig. 18b) without changes of phase composition ( ! ).
Similar results were also obtained for the BaO 2 ѿ MoO 2 system, in which the maximal rate of the SHS process with the formation of the complex oxide BaMoO 4 is achieved after preliminary mechanical activation for 30 s.
The structural similarity of the higher oxide WO 3 and barium tungstate BaWO 4 and rather high temperature of the reaction BaO 2 ѿWO 3 → BaWO 4 ѿ1/2 O 2 (Ҁ128 k J/mol) allow us to assume that this reaction can be conducted mechanochemically. It follows from the analysis of the IR spectra (Fig. 19). Similarly to the case of the lower oxide, the reaction starts during the first seconds of activation. The shape of the ν (W-O) band of WO 3 (1000Ҁ500 cm Ҁ1 ) and the ratio of intensities of their maximums are changed. After activation of the mixture for 1 min., the IR spectrum of the sample contains only one intensive band with a maximum at 810 cm Ҁ1 . This band relates to ν 3 vibrations of WO 4 tetrahedrons of the reaction product BaWO 4 . X-ray phase analysis showed that the growth of the intensity of the BaWO 4 phase ref lections was accompanied by the decrease of the WO 3 ref lections intensities till their complete disappearance after activation for 5 min. (Fig. 20). Microphotographs suggest that the initial stage of the process involves intensive dispersion of the particles; their aggregation starts at the second minute of the process. After activation for 5 min., complex oxide particles are formed; the size of their blocks is 20Ҁ30 nm (Fig. 21).
The mechanochemcial activation of BaO 2 with MoO 3 results in the formation of the complex oxide BaMoO 4 . The complex oxide with the same composition (BaMoO 4 ) can be obtained via SHS between barium peroxide and the higher molybdenum oxide, both reagents being activated preliminarily for 30 s.

Conclusion
Our investigations show that a combination of the SHS process with the mechanical activation both of reagents and products could be rather attractive for   technological applications. The preliminary activation facilitates combustion, making it possible even in the concentration region where conventional SHS is never observed. In other cases, the very short mechanical activation of SHS products allows us to prepare uniform single-phase products. For complex oxides, the picture is very similar if we can only provide the energetic conditions for SHS. There is usually quite a wide variety of precursors available for ceramic synthesis, which makes it possible to find an energetically efficient root.
The complex oxides can be obtained by mechanochemical synthesis, by SHS, and by a combination of mechanical activation and SHS.
The combination of mechanical activation and SHS brings advantages to both methods. As a result, the short-term mechanochemical activation of SHS products becomes an extremely convenient method for manufacturing nanopowders.