**Section 2**

**Applications, Techniques and Mineral Formation** 

204 Crystallization – Science and Technology

Shalaev, E. & Zografi, G. 2002. *The Concept of 'Structure' in Amorphous Solids from the* 

Shinde, K.R., Patil, S.D. & Dhake, A.S. 2011. Recrystallization. *Indian Streams Research Journal*,

Skonieczny, S. 2009. Crystallization. *University of Toronto course notes for CHM249*. p. 1-18. Stieger, N., Caira, M.R., Liebenberg, W., Tiedt, L.R., Wessels, J.C. & De Villiers, M.M. 2010a.

Recrystallization of Nevirapine. *Crystal Growth & Design*, 10(9):3859-3868. Stieger, N. & Liebenberg, W. 2009. Method for Increasing the Solubility of a Transcriptase Inhibitor Composition. *Patent application PCT/IB2010/055077* (priority date 2009-11-10). Stieger, N., Liebenberg, W. & Caira, M.R. 2009. Method of Producing a Polymorph Form.

Stieger, N., Liebenberg, W., Wessels, J.C., Samsodien, H. & Caira, M.R. 2010b. Channel

Stoica, C., Verwer, P., Meekes, H., Van Hoof, P.J.C.M., Kaspersen, F.M. & Vlieg, E. 2004.

Threlfall, T. 2000. Crystallisation of Polymorphs: Thermodynamic Insight into the Role of

Tiwary, A.K. 2006. *Crystal Habit Changes and Dosage Form Performance*. (*In* Swarbrick, J. *Ed*.

Togkalidou, T., Tung, H., Sun, Y., Andrews, A. & Braatz, R.D. 2002. Solution Concentration

Tros de Ilarduya, M.C., Martin, C., Goňi, M.M. & Martinez-Uhárriz, M.C. 1997. Dissolution

Tsai, S., Kuo, S. & Lin, S. 1993. Physicochemical Characterization of 9,10-anthraquinone-2-

Tung, H., Paul, E.L., Midler, M. & McCauly, J.A. 2009. *Crystallization of Organic Compounds:* 

Yu, L. 2001. Amorphous Pharmaceutical Solids: Preparation, Characterization and

Zaccaro, J., Matic, J., Myerson, A.S. & Garetz, B.A. 2001. Nonphotochemical, Laser-Induced

Zeitler, J.A., Taday, P.F., Gordon, K.C., Pepper, M. & Rades, T. 2007. New Insights into the

Zhou, G.X., Fujiwara, M., Woo, X.Y., Rusli, E., Tung, H., Starbuck, C., Davidson, O., Ge, Z. &

Resolved Terahertz Spectroscopy. *ChemPhysChem*, 8(13):1924-1927.

through Concentration Control. *Crystal Growth & Design*, 6(4):892-898.

Nucleation of Supersaturated Aqueous Glycine Produces Unexpected γ-

Solid-State Transition Mechanism in Carbamazepine Polymorphs by Time-

Braatz, R.D. 2006. Direct Design of Pharmaceutical Antisolvent Crystallization

carboxilic acid. *Journal of Pharmaceutical Sciences*, 82(12):1250-1254.

*An Industrial Perspective*. New Jersey: John Wiley & Sons. 289 p.

Stabilization. *Advanced Drug Delivery Reviews*, 48(1):27-42.

Drug Delivery Design*. Journal of Crystal Growth*, 211(1-4):122-136.

*Patent application PCT/IB2010/055808* (priority date 2009-12-17).

Solvent. *Organic Process Research & Development*, 4(5):384-390.

*Development and Industrial Pharmacy*, 23:1095-1098.

Polymorph. *Crystal Growth & Design*, 1(1):5-8.

1(4).

*Design*, 4(4):765-768.

Healthcare.) p. 822.

6(3):317-322.

*Perspective of the Pharmaceutical Sciences*. (*In* Levine, H. *Ed*. Amorphous Food and Pharmaceutical Systems. Cambridge: The Royal Society of Chemistry.) p. 11-30. Shekunov, B.Y. & York, P. 2000. Crystallization Processes in Pharmaceutical Technology and

Influence of the Composition of Water/Ethanol Mixtures on the Solubility and

Inclusion of Primary Alcohols in Isostructural Solvates of the Antiretroviral Nevirapine: an X-Ray an Thermal Analysis Study. *Structural Chemistry*, 21(4):771-777.

Understanding the Effect of a Solvent on the Crystal Habit. *Crystal Growth &* 

Encyclopedia of Pharmaceutical Technology. 3rd edition. New York: Informa

Prediction for Pharmaceutical Crystallization Processes Using Robust Chemometrics and ATR FTIR Spectroscopy. *Organic Process Research Development*,

Rates of Polymorphs and Two New Pseudopolymorphs of Sulindac. *Drug* 

**8** 

*1,2México 3Japan* 

**Preparation of Selected Ceramic Compounds by** 

The main aim of this chapter is to review some particular aspects related with the hydrothermal crystallization process for preparing some selected oxide and non-oxide powders, namely functional compound with electric, piezoelectric, ionic conducting and catalytic properties. The literature on the hydrothermal crystallization of oxide and nonoxide powders is vast; the references cited here are those more appropriated to illustrate this review. A particular attempt is made to broaden the traditional concepts of processing perovskite powders with controlled chemical composition and particle morphology, those aspects will be discussed based on the chemical reactivity of the precursor reactants (gels). Furthermore, an additional approach that takes into account the solubility of the solid species (mineral reagents) that are employed as a precursor in the hydrothermal systems on controlling the crystallization process of oxide particles was further investigated for preparing titanates oxide ceramic. The specific reaction pathways and kinetic aspects are discussed and illustrated by experimental setups for the solution of selected problems in hydrothermal crystallization. Also the chapter includes recent work on the formation of inorganic salts of Sr or Ba, under ordinary hydrothermal treatments via the anionic replacement in sulphate minerals, because these particular reactions have promoted peculiar microstructure on the crystallized material that preserves the bulk geometrical features of particular mineral specie. Therefore, this method could be attractive for preparing net-shaped materials with controlled porosity, with optimized functional properties, because can be used as gas sensor, substrates

for porous catalytic materials, filters, among other potential applications.

**2.1 Brief history of the hydrothermal technique development** 

**2. Hydrothermal crystallization as technique for oxide and non-oxide** 

The studies recorded at the scientific annals indicate that the pioneering research on hydrothermal systems was initiated in the middle of the 19th century (Schafthaul, 1845, as

**1. Introduction** 

**synthesis** 

**Controlled Crystallization Under** 

Juan Carlos Rendón-Angeles1, Zully Matamoros-Veloza2 and

*1Research Institute for Advanced Studies of the NPI Saltillo Campus* 

*3Research Laboratory for Hydrothermal Chemistry, Kochi University* 

**Hydrothermal Conditions** 

Kazumichi Yanagisawa3

*2Technological Institute of Saltillo* 

## **Preparation of Selected Ceramic Compounds by Controlled Crystallization Under Hydrothermal Conditions**

Juan Carlos Rendón-Angeles1, Zully Matamoros-Veloza2 and Kazumichi Yanagisawa3 *1Research Institute for Advanced Studies of the NPI Saltillo Campus 2Technological Institute of Saltillo 3Research Laboratory for Hydrothermal Chemistry, Kochi University 1,2México 3Japan* 

## **1. Introduction**

The main aim of this chapter is to review some particular aspects related with the hydrothermal crystallization process for preparing some selected oxide and non-oxide powders, namely functional compound with electric, piezoelectric, ionic conducting and catalytic properties. The literature on the hydrothermal crystallization of oxide and nonoxide powders is vast; the references cited here are those more appropriated to illustrate this review. A particular attempt is made to broaden the traditional concepts of processing perovskite powders with controlled chemical composition and particle morphology, those aspects will be discussed based on the chemical reactivity of the precursor reactants (gels). Furthermore, an additional approach that takes into account the solubility of the solid species (mineral reagents) that are employed as a precursor in the hydrothermal systems on controlling the crystallization process of oxide particles was further investigated for preparing titanates oxide ceramic. The specific reaction pathways and kinetic aspects are discussed and illustrated by experimental setups for the solution of selected problems in hydrothermal crystallization. Also the chapter includes recent work on the formation of inorganic salts of Sr or Ba, under ordinary hydrothermal treatments via the anionic replacement in sulphate minerals, because these particular reactions have promoted peculiar microstructure on the crystallized material that preserves the bulk geometrical features of particular mineral specie. Therefore, this method could be attractive for preparing net-shaped materials with controlled porosity, with optimized functional properties, because can be used as gas sensor, substrates for porous catalytic materials, filters, among other potential applications.

## **2. Hydrothermal crystallization as technique for oxide and non-oxide synthesis**

## **2.1 Brief history of the hydrothermal technique development**

The studies recorded at the scientific annals indicate that the pioneering research on hydrothermal systems was initiated in the middle of the 19th century (Schafthaul, 1845, as

Preparation of Selected Ceramic Compounds by

**2.3.1 Hydrothermal single crystal growth** 

(Suchanek et al., 2004).

Controlled Crystallization Under Hydrothermal Conditions 209

has remarkably reduced the processing parameters such as: reaction time, temperature, and pressure for hydrothermal crystallization of several oxide and non-oxide materials (T<200 ºC, P<1.5 MPa). At present, the recent scientific and technological achievements have made hydrothermal synthesis more economical, because powder preparation can be carried out by a single step cost-effective process, in advanced pressure reactor technology coupled with processing methodologies proposed for a wide number of inorganic compounds

Regardless of the numerous investigations on hydrothermal single crystal growth, quartz crystal is one of the materials extensively investigated up to now. Nowadays, the electronic industry requires the use of pure large single crystals of synthetic origin, because natural quartz crystals are generally irregular in shape, and is difficult to obtain large-scale single crystals wafers by automatic cutting. The crystal growth of oxide species, namely -quartz; is conducted by the conventional hydrothermal temperature gradient method, thus, the autoclaves frequently used consist of two chambers were different process occurs. One of the important parameter to consider for the single crystal growth of quartz is the selection of the proper nutrient material; among the most employed are small particle size quartz, silica glass, high quality silica sand, or silica gel (Byrappa 2005). The nutrient reactant is placed at the liner chamber (vessel bottom, see Figure 1) made up of iron, silver or titanium that is less employed, with a suitable baffle and a frame (at the top of the vessel) holding the seed of the material that is been grown. The mineralizer solution is other factor to select coupled with a definite molarity; this is poured into the liner to make the required filling volume and achieves the dissolution of the nutrient. Under these conditions, matter transport proceeds from the nutrient chamber and the crystallization and growth of the crystals is achieved due to the temperature difference at the top-seeded part of the autoclave. The particular optimum conditions determined at Bell Laboratories for growing quartz single crystals are dissolution temperature of 425 ºC at a pressure in the range of 100– 175 MPa, the crystallization and simultaneous growth proceeds at 375 ºC whilst the temperature gradient from the nutrient to the growing chamber was of 50 ºC. The mineralizer employed was an alkaline solution of NaOH with concentration varying between 0.5 up to 1 M, and the volume of poured mineralizer solution varied in the range of

78–85 % of the total volume of the vessel (Laudies, 1970; as cited in Byrappa, 2005).

The crystallization and growth of other single crystals rather than SiO2 was firstly investigated for oxide species, namely TiO2, ZrO2, HfO2 and some related perovskite oxides PbTiO3 and PbZrO3. Experimental results evidenced that the mineralizer employed for dissolving and transport the nutrient reactant markedly affected the crystallization of the oxides and the crystal growth. Thus, the single oxide species were found to dissolve and recrystallize faster in fluoride solutions (NaF, KF and NH4F) in comparison with alkalis, KOH and K2CO3, because of the high chemical reactivity of these nutrient oxides in fluoride solutions. However, the complete dissolution and transport of the nutrient was preferentially achieved in NH4F solutions rather than NaF or KF mineralizers, even at a low temperature of 470 ºC for a reaction interval up to 6 days (Kuznetzov, 1968), and the TiO2, ZrO2, HfO2 single crystal growth was achieved at a temperature of 520 ºC with a positive temperature gradient

**2.3 Hydrothermal processes for synthesis of inorganic compounds** 

cited in Suchanek et. al., 2004). The term of hydrothermal is purely of geological origin; this was firstly used by the geologist Sir Roderick Murchinson (1792–1871), whom analysed the reactivity of water at elevated pressure and temperature, to explain the mineral formation of several rocks and minerals (Byrappa & Yoshimura, 2001). Continuous progress in material synthesis was initially accelerated by notorious developments in hydrothermal pressure vessels apparatus. Hence, at the first decade of the 20th century, geologist achieved a marked technological expansion on the hydrothermal research field, namely efforts were directed towards for designing the first metal vessels and carried out preliminary experiments at laboratory scale related to material synthesis area, mainly in the field of inorganic single crystal growth (Riman et. al., 2002; Suchanek et. al., 2004). However, the development of chemical reactions via hydrothermal processing was further limited because of the severe treatment (supercritical) conditions, which were normally required for single crystal growth. This resulted on discourage extensive research and commercialization of various materials. Hydrothermal epitaxial growth is one example, it was popular during the 1970s, but it did not reach commercial success due to the high temperatures (> 500 ºC) and pressures (> 100 MPa), which were involved to achieve the epitaxial crystal growth process. In the middle of the 1980s, the commercial interest in the hydrothermal technology was revived, because the steadily increasing of a large group of materials, manly ceramic powder, had emerged that can be produced under more environmentally friendly conditions (T < 350 ºC, P < 100 MPa). At present, the major developments on the hydrothermal synthesis technology including in particular the hydrothermal crystallization has been accomplished in several countries around the world like: China, Japan, USA, UK, Germany and some others with minor contributions (Byrappa & Yoshimura, 2001).

#### **2.2 Definition of hydrothermal synthesis**

The term "*hydrothermal*" is difficult to define, based in the etymologic root of the Greek word, "*hydrous*" means water and "*thermal*" means heat. One of the accepted statements for hydrothermal defines it as any heterogeneous chemical reaction that occurs in the presence of a solvent media at above the room temperature (> 25 ºC) and pressure levels greater than 0.1 MPa in a closed system, at these conditions does not matter whether the solvent is aqueous or non-aqueous. Hitherto, there is still a bit of confusion regarding the correct use of the term hydrothermal, because in the case of the chemistry field, the chemists prefer to use a different term, namely *solvothermal*, which means any chemical reaction conducted with a non-aqueous solvent or solvent in supercritical conditions. Additionally, another similar related terms widely used amongst the physicist, chemist and material scientist communities are: *glycothermal*, *alcothermal*, *ammonothermal*, and so on. However, some researchers also use hydrothermal for describing processes conducted at ambient conditions. The *crystallization* process of solid phases under hydrothermal conditions is usually conducted at autogeneous pressure, achieving a particular saturated vapour pressure of the solution at the specified temperature and composition of the hydrothermal solution. In this concern, in terms of industrial commercial processes mild operating conditions are preferred, for example treatment temperatures below than 350 ºC and pressures > 50 MPa (Byrappa, 2005). The limit that indicates the transition from mild to severe conditions during a hydrothermal treatment is normally determined by the strength of the inner vessel materials, which at severe treatment conditions undergoes into a corrosion process. The continuous research in this field has led the way to a better understanding for controlling chemical reactions in various hydrothermal medias, which

cited in Suchanek et. al., 2004). The term of hydrothermal is purely of geological origin; this was firstly used by the geologist Sir Roderick Murchinson (1792–1871), whom analysed the reactivity of water at elevated pressure and temperature, to explain the mineral formation of several rocks and minerals (Byrappa & Yoshimura, 2001). Continuous progress in material synthesis was initially accelerated by notorious developments in hydrothermal pressure vessels apparatus. Hence, at the first decade of the 20th century, geologist achieved a marked technological expansion on the hydrothermal research field, namely efforts were directed towards for designing the first metal vessels and carried out preliminary experiments at laboratory scale related to material synthesis area, mainly in the field of inorganic single crystal growth (Riman et. al., 2002; Suchanek et. al., 2004). However, the development of chemical reactions via hydrothermal processing was further limited because of the severe treatment (supercritical) conditions, which were normally required for single crystal growth. This resulted on discourage extensive research and commercialization of various materials. Hydrothermal epitaxial growth is one example, it was popular during the 1970s, but it did not reach commercial success due to the high temperatures (> 500 ºC) and pressures (> 100 MPa), which were involved to achieve the epitaxial crystal growth process. In the middle of the 1980s, the commercial interest in the hydrothermal technology was revived, because the steadily increasing of a large group of materials, manly ceramic powder, had emerged that can be produced under more environmentally friendly conditions (T < 350 ºC, P < 100 MPa). At present, the major developments on the hydrothermal synthesis technology including in particular the hydrothermal crystallization has been accomplished in several countries around the world like: China, Japan, USA, UK, Germany and some others with minor

The term "*hydrothermal*" is difficult to define, based in the etymologic root of the Greek word, "*hydrous*" means water and "*thermal*" means heat. One of the accepted statements for hydrothermal defines it as any heterogeneous chemical reaction that occurs in the presence of a solvent media at above the room temperature (> 25 ºC) and pressure levels greater than 0.1 MPa in a closed system, at these conditions does not matter whether the solvent is aqueous or non-aqueous. Hitherto, there is still a bit of confusion regarding the correct use of the term hydrothermal, because in the case of the chemistry field, the chemists prefer to use a different term, namely *solvothermal*, which means any chemical reaction conducted with a non-aqueous solvent or solvent in supercritical conditions. Additionally, another similar related terms widely used amongst the physicist, chemist and material scientist communities are: *glycothermal*, *alcothermal*, *ammonothermal*, and so on. However, some researchers also use hydrothermal for describing processes conducted at ambient conditions. The *crystallization* process of solid phases under hydrothermal conditions is usually conducted at autogeneous pressure, achieving a particular saturated vapour pressure of the solution at the specified temperature and composition of the hydrothermal solution. In this concern, in terms of industrial commercial processes mild operating conditions are preferred, for example treatment temperatures below than 350 ºC and pressures > 50 MPa (Byrappa, 2005). The limit that indicates the transition from mild to severe conditions during a hydrothermal treatment is normally determined by the strength of the inner vessel materials, which at severe treatment conditions undergoes into a corrosion process. The continuous research in this field has led the way to a better understanding for controlling chemical reactions in various hydrothermal medias, which

contributions (Byrappa & Yoshimura, 2001).

**2.2 Definition of hydrothermal synthesis** 

has remarkably reduced the processing parameters such as: reaction time, temperature, and pressure for hydrothermal crystallization of several oxide and non-oxide materials (T<200 ºC, P<1.5 MPa). At present, the recent scientific and technological achievements have made hydrothermal synthesis more economical, because powder preparation can be carried out by a single step cost-effective process, in advanced pressure reactor technology coupled with processing methodologies proposed for a wide number of inorganic compounds (Suchanek et al., 2004).

## **2.3 Hydrothermal processes for synthesis of inorganic compounds**

## **2.3.1 Hydrothermal single crystal growth**

Regardless of the numerous investigations on hydrothermal single crystal growth, quartz crystal is one of the materials extensively investigated up to now. Nowadays, the electronic industry requires the use of pure large single crystals of synthetic origin, because natural quartz crystals are generally irregular in shape, and is difficult to obtain large-scale single crystals wafers by automatic cutting. The crystal growth of oxide species, namely -quartz; is conducted by the conventional hydrothermal temperature gradient method, thus, the autoclaves frequently used consist of two chambers were different process occurs. One of the important parameter to consider for the single crystal growth of quartz is the selection of the proper nutrient material; among the most employed are small particle size quartz, silica glass, high quality silica sand, or silica gel (Byrappa 2005). The nutrient reactant is placed at the liner chamber (vessel bottom, see Figure 1) made up of iron, silver or titanium that is less employed, with a suitable baffle and a frame (at the top of the vessel) holding the seed of the material that is been grown. The mineralizer solution is other factor to select coupled with a definite molarity; this is poured into the liner to make the required filling volume and achieves the dissolution of the nutrient. Under these conditions, matter transport proceeds from the nutrient chamber and the crystallization and growth of the crystals is achieved due to the temperature difference at the top-seeded part of the autoclave. The particular optimum conditions determined at Bell Laboratories for growing quartz single crystals are dissolution temperature of 425 ºC at a pressure in the range of 100– 175 MPa, the crystallization and simultaneous growth proceeds at 375 ºC whilst the temperature gradient from the nutrient to the growing chamber was of 50 ºC. The mineralizer employed was an alkaline solution of NaOH with concentration varying between 0.5 up to 1 M, and the volume of poured mineralizer solution varied in the range of 78–85 % of the total volume of the vessel (Laudies, 1970; as cited in Byrappa, 2005).

The crystallization and growth of other single crystals rather than SiO2 was firstly investigated for oxide species, namely TiO2, ZrO2, HfO2 and some related perovskite oxides PbTiO3 and PbZrO3. Experimental results evidenced that the mineralizer employed for dissolving and transport the nutrient reactant markedly affected the crystallization of the oxides and the crystal growth. Thus, the single oxide species were found to dissolve and recrystallize faster in fluoride solutions (NaF, KF and NH4F) in comparison with alkalis, KOH and K2CO3, because of the high chemical reactivity of these nutrient oxides in fluoride solutions. However, the complete dissolution and transport of the nutrient was preferentially achieved in NH4F solutions rather than NaF or KF mineralizers, even at a low temperature of 470 ºC for a reaction interval up to 6 days (Kuznetzov, 1968), and the TiO2, ZrO2, HfO2 single crystal growth was achieved at a temperature of 520 ºC with a positive temperature gradient

Preparation of Selected Ceramic Compounds by

1999).

Controlled Crystallization Under Hydrothermal Conditions 211

single crystals with pyrochlore structure Pb1.83Mg0.29Nb1.71O6.39 and Pb1.83Sc0.29Nb1.71O6.39, as shown in Figure 2 (Yanagisawa et. al., 1999, 2000). This incongruence has not been determined thermodynamically and experimentally for the PZT, because this compound has been used to grown thin and thick single crystals on different seeds (SrTiO3) with excellent compositional control (Oledzka et. al., 2003). In general, the crystallization and crystal growth of a specific compound can be carried out using different mineralizer or solvents, e.g., water, soluble salts, acid solutions, non aqueous solutions; but some physicochemical factors must be taking into account to fulfil suitable conditions that facilitate both processes.

Fig. 2. Pyrochlore single crystals, of (a) PMN, (b) PSNT and (c) PSN; grown at the top of platinum capsule at 600 ºC with a gradient temperature gradient of 40 ºC in KF solutions 5, 4.2 and 6.2 M for reaction intervals of 3 (a) and 5 days (c,d), respectively (Yanagisawa et. al.,

Remarkable achievements for developing the *Hydrothermal Microwave Assisted Synthesis* process have been conducted at Pennsylvania State University (Komarneni et. al., 1992); this method enhances solid crystallization kinetics 1–4 order of magnitude faster than that occurring on the conventional hydrothermal processing for a wide solutions. In addition, other advantages of the hydrothermal microwave assisted technique are very high heating

**2.3.2 Advanced hydrothermal processing methods** 

(30 ºC). This is a popular method that promotes the crystal growth and has been widely used on the preparation of berlinite (AlPO4) crystals (Byrappa & Yoshimura 2001).

Fig. 1. Scheme of the conventional autoclave employed for single crystal growth under hydrothermal conditions (Schubert, 2000; as cited in Byrappa & Yoshimura, 2001).

A particular emphasize on the hydrothermal crystal growth research has been focused for the crystallization of materials that melts incongruently because such materials cannot be grown with compositional and phase uniformity. The relevant examples of compounds that melts incongruently are those belonging to lead family oxides titanate (PbTiO3), zirconate (PbZrO3), single crystals with millimetre size of both compounds were found to grown satisfactorily in NH4F hydrothermal solution by seedless grown crystallizing conditions at temperatures between 500–600 ºC for several days (Kuznetzov, 1968). However, the hydrothermal crystal grown of some lead related solid solutions, PbZr1-xTixO3 (PZT), PbMg1/3Nb2/3O3 (PMN) and PbSc0.5Nb0.5O3 (PSN), is limited due to the low chemical stability of the perovskite structure even in KF solution, and results on the grown of small

(30 ºC). This is a popular method that promotes the crystal growth and has been widely used

Fig. 1. Scheme of the conventional autoclave employed for single crystal growth under hydrothermal conditions (Schubert, 2000; as cited in Byrappa & Yoshimura, 2001).

A particular emphasize on the hydrothermal crystal growth research has been focused for the crystallization of materials that melts incongruently because such materials cannot be grown with compositional and phase uniformity. The relevant examples of compounds that melts incongruently are those belonging to lead family oxides titanate (PbTiO3), zirconate (PbZrO3), single crystals with millimetre size of both compounds were found to grown satisfactorily in NH4F hydrothermal solution by seedless grown crystallizing conditions at temperatures between 500–600 ºC for several days (Kuznetzov, 1968). However, the hydrothermal crystal grown of some lead related solid solutions, PbZr1-xTixO3 (PZT), PbMg1/3Nb2/3O3 (PMN) and PbSc0.5Nb0.5O3 (PSN), is limited due to the low chemical stability of the perovskite structure even in KF solution, and results on the grown of small

on the preparation of berlinite (AlPO4) crystals (Byrappa & Yoshimura 2001).

single crystals with pyrochlore structure Pb1.83Mg0.29Nb1.71O6.39 and Pb1.83Sc0.29Nb1.71O6.39, as shown in Figure 2 (Yanagisawa et. al., 1999, 2000). This incongruence has not been determined thermodynamically and experimentally for the PZT, because this compound has been used to grown thin and thick single crystals on different seeds (SrTiO3) with excellent compositional control (Oledzka et. al., 2003). In general, the crystallization and crystal growth of a specific compound can be carried out using different mineralizer or solvents, e.g., water, soluble salts, acid solutions, non aqueous solutions; but some physicochemical factors must be taking into account to fulfil suitable conditions that facilitate both processes.

Fig. 2. Pyrochlore single crystals, of (a) PMN, (b) PSNT and (c) PSN; grown at the top of platinum capsule at 600 ºC with a gradient temperature gradient of 40 ºC in KF solutions 5, 4.2 and 6.2 M for reaction intervals of 3 (a) and 5 days (c,d), respectively (Yanagisawa et. al., 1999).

## **2.3.2 Advanced hydrothermal processing methods**

Remarkable achievements for developing the *Hydrothermal Microwave Assisted Synthesis* process have been conducted at Pennsylvania State University (Komarneni et. al., 1992); this method enhances solid crystallization kinetics 1–4 order of magnitude faster than that occurring on the conventional hydrothermal processing for a wide solutions. In addition, other advantages of the hydrothermal microwave assisted technique are very high heating

Preparation of Selected Ceramic Compounds by

industrial scale (Shi & Hwang, 2003).

**materials** 

Controlled Crystallization Under Hydrothermal Conditions 213

(Ca10(PO4)6(OH)2, AlPO4, InSb, CdS) and thin films (Li2B4O7, Ba2TiSi2O8).while the average

The advanced hydrothermal techniques discussed in this section, even though wet chemical synthesis is offer in conjunction with so many significant advantages over the conventional method, might be able to have wide application in the industry. Compared to solid-state materials processing, these technologies might be more facile for scaling up. This situation, however, still cannot guarantee successful application of all these technologies to industries. Hence, much effort is needed to obtain a more comprehensive understanding of several physicochemical phenomena related to each of the new technologies hybridizing the conventional hydrothermal synthesis, in order to establish a relationship between science and technology that could lead to optimize these methods for their employment at a

**2.4 Conventional hydrothermal crystallization process for advanced ceramic** 

**2.4.1 Factors that affect optimizing hydrothermal crystallization experiments** 

In general, a majority of the hydrothermal synthesis research work that has been done in the past was based on Edisonian trial and error design for process development, but this is not the best experimental approach for discerning between processes that are controlled by either thermodynamics or kinetics. In contrast, much effort has been paid to use thermodynamic modelling for processing design; this approach is based on fundamental physicochemical principles instead of the Edisonian methods (Riman et. al., 2002). Therefore, a great number of hydrothermal fundamental works for some particular solidaqueous systems have provided sufficient experimental hydrothermal physical chemistry data. An important point derived from these works is related with the behaviour of the solvent under hydrothermal conditions, because it has a relationship with aspects like structure at critical, supercritical and sub-critical conditions, solution dielectric constant, pH variation, viscosity, expansion coefficient, density, etc., all these parameters depend markedly on thermodynamic variables such as pressure and temperature. At present, hydrothermal crystallization process is the only one where a fundamental understanding of kinetics is lacking due to the absence of physicochemical data of the intermediate phases forming in specific aqueous solutions. Although, fundamental research works related to synthesis of specific compounds demonstrated the importance of crystallization kinetics, however, a better understanding of crystallization kinetics still in an early stage of development. In this case, due to the absence of predictive methodology models, ones must estimated on terms of chemical equilibrium of the reaction the effect of temperature, pressure, precursor, and time to achieve solid crystallization and improve the reaction kinetics. Insight into this would enable us to understand how to control the formation of ionic species in the solution, the crystallization of solid phases and the rate of their growth. In the last decade, much effort has directed towards for developing thermochemical models based in fundamental knowledge of thermodynamic and the Hegelson-Kirkham-Flowers-Tanger equation of state, which allows to represent standard-state properties across substantial temperature and pressure ranges in order to estimate chemical reactions path ways under hydrothermal conditions (Lencka & Riman, 2002). Modelling can be

temperature of the reactors is maintained close to room temperature.

rates and the synthesis of novel phases. A great variety of ceramic powders with particular morphologies and controlled particle size, have been produced, e.g. TiO2, ZrO2, Fe2O3, BaTiO3, Ca10(PO4)6(OH)2, etc (Komarneni et. al., 1992; Roy, 1994; Komarneni et. al., 2002).

Another variation implemented to the hydrothermal conventional process resulted on the development of a different hybrid method denominated *Hydrothermal-Electrochemical Synthesis*. This technique was tailored to deposit polycrystalline oxide films on reactive metal substrates. This technique becomes very important when the crystallization of oxide products from supersaturated hydrothermal solutions is hinder in absence of an applied electrical potential. Nowadays, highly crystallized ceramic thin films, such as a BaTiO3, SrTiO3, LiNiO2, PbTiO3, CaWO4 and BaMoO4 can be deposited on metal substrates from aqueous solutions at relatively low temperature 50–200 ºC for several hours under a continuous applied voltage charge (Yoshimura & Suchanek, 1997). Additionally, semiconductor thin layers of GaAs, CdTe, CdSe and CdS have been successfully prepared by electrochemical atomic layer epitaxy growth, this mechanism is analogous to molecular epitaxy, however, the crystallization and growth of the layer is enhanced from a saturated hydrothermal media instead of a vapour phase that transport the growing species (Colletti et. al., 1998; as cited in Suchanek et. al., 2004).

Among other tailored techniques, the high energy milling technique has been challenged to hybridize the conventional hydrothermal processing, resulting in a coupled process that involves the classical powder mechanochemical and hydrothermal syntheses. The *Mechanochemical-Hydrothermal* route utilizes the solvency of an aqueous solution, which facilities the crystallization of solid species due to the pressure environment generated inside the mechano-chemical autoclave; it mainly accelerates the rate-determining steps of those factors, for instance: interfacial reaction, solute dissolution or dehydroxylation; that limit hydrothermal chemical reactions at low temperature. These reactions enhance the crystallization to occur locally at the particle surface because of the perturbation of superficial bonded species in the solid coupled with high temperature gradients (400–700 ºC) and pressure localized zones that are generated during the mechanical activation of slurries. This is mainly promoted by the friction and adiabatic heating of gas bubbles while maintaining the average temperature close to room temperature (Kosova et. al., 1997; as cited in Suchanek et. al., 2004). Hitherto, this technique has been used for preparing bioceramic materials such as hydroxyapatite (Suchanek et. al., 2002; Chen et. al., 2004), another tailored material like PbTiO3 was also crystallized by this technique.

Since Rustum Roy reported that the use of ultrasonic devices are feasible for improving low temperature inorganic syntheses because reaction kinetics is two orders of magnitude faster than that for the conventional wet chemistry synthesis methods (Roy, 1994). Some attempts to adapt ultrasonic devices emitting acoustic signals of 20 kHz up to 10 MHz have been conducted, because the acoustic signal produce very sharp temperature gradients with localized peak temperature zones, speculated as high as 5000 K, as well as, localized pressure zones of up to 100's MPa. The sonochemical environment also alters the molecular chemistry (chemical bond scission, generate excited states and accelerate electron transfer steps in chemical reactions), and enhances mass transport and crystallization kinetics due to the high convection of the fluid. (Peters, 1996, as cited in Riman et. al., 2002). *Hydrothermalsonochemical synthesis* method have been used for preparing various ceramic powders

rates and the synthesis of novel phases. A great variety of ceramic powders with particular morphologies and controlled particle size, have been produced, e.g. TiO2, ZrO2, Fe2O3, BaTiO3, Ca10(PO4)6(OH)2, etc (Komarneni et. al., 1992; Roy, 1994; Komarneni et. al., 2002).

Another variation implemented to the hydrothermal conventional process resulted on the development of a different hybrid method denominated *Hydrothermal-Electrochemical Synthesis*. This technique was tailored to deposit polycrystalline oxide films on reactive metal substrates. This technique becomes very important when the crystallization of oxide products from supersaturated hydrothermal solutions is hinder in absence of an applied electrical potential. Nowadays, highly crystallized ceramic thin films, such as a BaTiO3, SrTiO3, LiNiO2, PbTiO3, CaWO4 and BaMoO4 can be deposited on metal substrates from aqueous solutions at relatively low temperature 50–200 ºC for several hours under a continuous applied voltage charge (Yoshimura & Suchanek, 1997). Additionally, semiconductor thin layers of GaAs, CdTe, CdSe and CdS have been successfully prepared by electrochemical atomic layer epitaxy growth, this mechanism is analogous to molecular epitaxy, however, the crystallization and growth of the layer is enhanced from a saturated hydrothermal media instead of a vapour phase that transport the growing species (Colletti

Among other tailored techniques, the high energy milling technique has been challenged to hybridize the conventional hydrothermal processing, resulting in a coupled process that involves the classical powder mechanochemical and hydrothermal syntheses. The *Mechanochemical-Hydrothermal* route utilizes the solvency of an aqueous solution, which facilities the crystallization of solid species due to the pressure environment generated inside the mechano-chemical autoclave; it mainly accelerates the rate-determining steps of those factors, for instance: interfacial reaction, solute dissolution or dehydroxylation; that limit hydrothermal chemical reactions at low temperature. These reactions enhance the crystallization to occur locally at the particle surface because of the perturbation of superficial bonded species in the solid coupled with high temperature gradients (400–700 ºC) and pressure localized zones that are generated during the mechanical activation of slurries. This is mainly promoted by the friction and adiabatic heating of gas bubbles while maintaining the average temperature close to room temperature (Kosova et. al., 1997; as cited in Suchanek et. al., 2004). Hitherto, this technique has been used for preparing bioceramic materials such as hydroxyapatite (Suchanek et. al., 2002; Chen et. al., 2004),

another tailored material like PbTiO3 was also crystallized by this technique.

Since Rustum Roy reported that the use of ultrasonic devices are feasible for improving low temperature inorganic syntheses because reaction kinetics is two orders of magnitude faster than that for the conventional wet chemistry synthesis methods (Roy, 1994). Some attempts to adapt ultrasonic devices emitting acoustic signals of 20 kHz up to 10 MHz have been conducted, because the acoustic signal produce very sharp temperature gradients with localized peak temperature zones, speculated as high as 5000 K, as well as, localized pressure zones of up to 100's MPa. The sonochemical environment also alters the molecular chemistry (chemical bond scission, generate excited states and accelerate electron transfer steps in chemical reactions), and enhances mass transport and crystallization kinetics due to the high convection of the fluid. (Peters, 1996, as cited in Riman et. al., 2002). *Hydrothermalsonochemical synthesis* method have been used for preparing various ceramic powders

et. al., 1998; as cited in Suchanek et. al., 2004).

(Ca10(PO4)6(OH)2, AlPO4, InSb, CdS) and thin films (Li2B4O7, Ba2TiSi2O8).while the average temperature of the reactors is maintained close to room temperature.

The advanced hydrothermal techniques discussed in this section, even though wet chemical synthesis is offer in conjunction with so many significant advantages over the conventional method, might be able to have wide application in the industry. Compared to solid-state materials processing, these technologies might be more facile for scaling up. This situation, however, still cannot guarantee successful application of all these technologies to industries. Hence, much effort is needed to obtain a more comprehensive understanding of several physicochemical phenomena related to each of the new technologies hybridizing the conventional hydrothermal synthesis, in order to establish a relationship between science and technology that could lead to optimize these methods for their employment at a industrial scale (Shi & Hwang, 2003).

#### **2.4 Conventional hydrothermal crystallization process for advanced ceramic materials**

### **2.4.1 Factors that affect optimizing hydrothermal crystallization experiments**

In general, a majority of the hydrothermal synthesis research work that has been done in the past was based on Edisonian trial and error design for process development, but this is not the best experimental approach for discerning between processes that are controlled by either thermodynamics or kinetics. In contrast, much effort has been paid to use thermodynamic modelling for processing design; this approach is based on fundamental physicochemical principles instead of the Edisonian methods (Riman et. al., 2002). Therefore, a great number of hydrothermal fundamental works for some particular solidaqueous systems have provided sufficient experimental hydrothermal physical chemistry data. An important point derived from these works is related with the behaviour of the solvent under hydrothermal conditions, because it has a relationship with aspects like structure at critical, supercritical and sub-critical conditions, solution dielectric constant, pH variation, viscosity, expansion coefficient, density, etc., all these parameters depend markedly on thermodynamic variables such as pressure and temperature. At present, hydrothermal crystallization process is the only one where a fundamental understanding of kinetics is lacking due to the absence of physicochemical data of the intermediate phases forming in specific aqueous solutions. Although, fundamental research works related to synthesis of specific compounds demonstrated the importance of crystallization kinetics, however, a better understanding of crystallization kinetics still in an early stage of development. In this case, due to the absence of predictive methodology models, ones must estimated on terms of chemical equilibrium of the reaction the effect of temperature, pressure, precursor, and time to achieve solid crystallization and improve the reaction kinetics. Insight into this would enable us to understand how to control the formation of ionic species in the solution, the crystallization of solid phases and the rate of their growth.

In the last decade, much effort has directed towards for developing thermochemical models based in fundamental knowledge of thermodynamic and the Hegelson-Kirkham-Flowers-Tanger equation of state, which allows to represent standard-state properties across substantial temperature and pressure ranges in order to estimate chemical reactions path ways under hydrothermal conditions (Lencka & Riman, 2002). Modelling can be

Preparation of Selected Ceramic Compounds by

Suchanek, 1997; Riman et. al., 2002).

SrTiO3, BaTiO3 and La1-XMXCr1-YNYO3.

(Marchizo, 2009).

Controlled Crystallization Under Hydrothermal Conditions 215

Moreover the model shows that secondary nucleation is indeed very important but particle aggregation cannot be neglected, these variables were not considered in the previous model

Despite of the remarkable achievement to develop new strategies based on fundamental principles of thermodynamics and chemistry (reaction equilibrium and kinetics), caution should bear in mind for applying the proposed models to diverse hydrothermal systems and experimental situations. Hence, the estimation of kinetic parameters using population balance derived models could not be correct, these parameters can exhibit variations due to differences on temperature and concentration in the fluid media, which would lead to erroneous conclusions, because the reaction mechanism at certain conditions might change. Likewise, the population balance derived models assume that the system is well mixed and the crystallization rate is uniquely controlled by the chemical kinetics. These assumptions are generally valid for some small-scale laboratory reactors, but they fail in larger industrial plants where mixing is not well controlled. A mixing-limited precipitation rate is one of the

problems commonly encountered in the scale-up of solid precipitation processes.

In the past, the term *hydrothermal crystallization* was used to referred it as nonconventional chemical process, this process involves heating an aqueous suspension of insoluble salts in an autoclave at temperatures and pressures greater than 100 ºC and 0.1 MPa, respectively; resulting in the crystallization of the desired well-crystallized phases. This process, however, can be analogous to the term *hydrothermal synthesis* broadly used for physicists and chemists involved in this type of research, because both terms are related with the *genesis* of a specific compound. The crystallization process can mainly occur under very precise conditions of reaction, i.e. temperature, pressure, pH, concentration mineralizer or solvent solution. Thus, the hydrothermal crystallization method using inexpensive and chemicals easy to handle, can produce single or multicomponent oxides. The advantages of hydrothermal crystallization are the reduced energy costs due to the mild temperatures sufficient to achieve chemical reactions, less pollution, simplicity in the process equipment, and the fast rate of solid precipitation reactions (Vivekanandan & Kutty, 1989; Yoshimura &

There are thousands of research works in the literature that include a vast experimental data related with the hydrothermal crystallization of metal oxides or inorganic compounds. The most popular among the metal oxides are those of perovskite oxides group, because of their wide application at the electronic industry. Hence, fundamental principles of crystallization of these materials are discussed in detail and compared with the present research experimental of the present authors regarding the crystallization of perovskite particles of

**2.4.3 Hydrothermal crystallization of perovskite derived oxides from gel precursors**  In this section, we focus on the hydrothermal crystallization of perovskite-type structure oxide materials, because this particular group of oxides offers stupendous functional properties for a wide number of applications, in particular electronic applications. Perovskite compounds have the general formula ABO3, and Perovskite is cubic structure in

**2.4.2 Hydrothermal crystallization process of selected oxide compounds** 

successfully applied to very complex aqueous electrolytes over specific ranges of temperature and reactant and solvent concentration and non-aqueous systems can also be model as well. Practical computer software for conducting thermochemical modelling studies was recently developed by OLI System Inc., USA. This software can be used to conduct studies for chemical reactions for hydrothermal systems within the temperature range of -50 to 300 ºC, at pressures ranging from 0 to 150 MPa and concentration of 0 up to 30 m in molar ionic strength; for the non-aqueous systems the temperature range covered is from 0 to 1200 ºC and pressure from 0 to 150 MPa with species concentration from 0 to 1.0 mole fraction. Major predictions using the thermodynamic model have been done to determine the optimum hydrothermal conditions for achieving the crystallization of a wide variety of oxides such as, BaTiO3, PbTiO3, CaTiO3, SrZrO3 and (BaSr)TiO3 (Lencka & Riman, 1993, 1995; Gersten et. al., 2004).

Currently, the research efforts on the thermochemical-modelling topic are being focused to establish an overall rational engineering-based methodology that will speed up process development. The proposed methodology for conducting this study involves the following four steps:


Recently, population balance modelling has received much attention from both academic and industrial areas because of its applicability to a wide variety of particular processes. In general, a population balance model can be proposed by the collective phenomenology contained in entities displacement through their state space and the birth-and-death processes that terminate existing entities and produce new entities. The phenomenology concerns the behaviour of any single entity in conjunction with other entities, which is for the population balance modelling a reasonable description of the system (Ramkrishna & Mahoney, 2002). Regarding the hydrothermal crystallization process, a rigorous kinetic model for the solution precipitation and hydrothermal synthesis of BaTiO3 particles based on mass and population balances has been recently proposed (Testino et. al., 2005). The population balances considered three elementary chemical kinetic processes: primary nucleation, secondary nucleation, and diffusion-controlled particle growth. Secondary nucleation accounts for the acceleration of the formation kinetics after an initial slow crystallization of BaTiO3 particles. The time evolution of yield and particle size is represented by means of discretized mass and population balance equations. Hence, the algorithm is capable for calculating a generic number of conditions involved in the chemical reaction, which can optimize the crystallization and control the growth of BaTiO3 particles. Another different approach based in a bi-variate population balance equations have prove to give more accurate results for the modelling of the barium titanate hydrothermal crystallization. The results obtained from the proposed population balance mathematical model clearly showed that such an approach can overcome the limitations of previous modelling work, and provide a useful tool for more detailed kinetic parameters estimation.

successfully applied to very complex aqueous electrolytes over specific ranges of temperature and reactant and solvent concentration and non-aqueous systems can also be model as well. Practical computer software for conducting thermochemical modelling studies was recently developed by OLI System Inc., USA. This software can be used to conduct studies for chemical reactions for hydrothermal systems within the temperature range of -50 to 300 ºC, at pressures ranging from 0 to 150 MPa and concentration of 0 up to 30 m in molar ionic strength; for the non-aqueous systems the temperature range covered is from 0 to 1200 ºC and pressure from 0 to 150 MPa with species concentration from 0 to 1.0 mole fraction. Major predictions using the thermodynamic model have been done to determine the optimum hydrothermal conditions for achieving the crystallization of a wide variety of oxides such as, BaTiO3, PbTiO3, CaTiO3, SrZrO3 and (BaSr)TiO3 (Lencka & Riman,

Currently, the research efforts on the thermochemical-modelling topic are being focused to establish an overall rational engineering-based methodology that will speed up process development. The proposed methodology for conducting this study involves the following

1. Compute thermodynamic equilibrium as a function of chemical processing variables. 2. Generate equilibrium diagrams to draw the process variable space for the phases of

4. Utilize the processing variables to explore opportunities for controlling reactions and

Recently, population balance modelling has received much attention from both academic and industrial areas because of its applicability to a wide variety of particular processes. In general, a population balance model can be proposed by the collective phenomenology contained in entities displacement through their state space and the birth-and-death processes that terminate existing entities and produce new entities. The phenomenology concerns the behaviour of any single entity in conjunction with other entities, which is for the population balance modelling a reasonable description of the system (Ramkrishna & Mahoney, 2002). Regarding the hydrothermal crystallization process, a rigorous kinetic model for the solution precipitation and hydrothermal synthesis of BaTiO3 particles based on mass and population balances has been recently proposed (Testino et. al., 2005). The population balances considered three elementary chemical kinetic processes: primary nucleation, secondary nucleation, and diffusion-controlled particle growth. Secondary nucleation accounts for the acceleration of the formation kinetics after an initial slow crystallization of BaTiO3 particles. The time evolution of yield and particle size is represented by means of discretized mass and population balance equations. Hence, the algorithm is capable for calculating a generic number of conditions involved in the chemical reaction, which can optimize the crystallization and control the growth of BaTiO3 particles. Another different approach based in a bi-variate population balance equations have prove to give more accurate results for the modelling of the barium titanate hydrothermal crystallization. The results obtained from the proposed population balance mathematical model clearly showed that such an approach can overcome the limitations of previous modelling work, and provide a useful tool for more detailed kinetic parameters estimation.

3. Design hydrothermal experiments to test and validate the computed diagrams.

1993, 1995; Gersten et. al., 2004).

crystallization kinetics.

four steps:

interest.

Moreover the model shows that secondary nucleation is indeed very important but particle aggregation cannot be neglected, these variables were not considered in the previous model (Marchizo, 2009).

Despite of the remarkable achievement to develop new strategies based on fundamental principles of thermodynamics and chemistry (reaction equilibrium and kinetics), caution should bear in mind for applying the proposed models to diverse hydrothermal systems and experimental situations. Hence, the estimation of kinetic parameters using population balance derived models could not be correct, these parameters can exhibit variations due to differences on temperature and concentration in the fluid media, which would lead to erroneous conclusions, because the reaction mechanism at certain conditions might change. Likewise, the population balance derived models assume that the system is well mixed and the crystallization rate is uniquely controlled by the chemical kinetics. These assumptions are generally valid for some small-scale laboratory reactors, but they fail in larger industrial plants where mixing is not well controlled. A mixing-limited precipitation rate is one of the problems commonly encountered in the scale-up of solid precipitation processes.

## **2.4.2 Hydrothermal crystallization process of selected oxide compounds**

In the past, the term *hydrothermal crystallization* was used to referred it as nonconventional chemical process, this process involves heating an aqueous suspension of insoluble salts in an autoclave at temperatures and pressures greater than 100 ºC and 0.1 MPa, respectively; resulting in the crystallization of the desired well-crystallized phases. This process, however, can be analogous to the term *hydrothermal synthesis* broadly used for physicists and chemists involved in this type of research, because both terms are related with the *genesis* of a specific compound. The crystallization process can mainly occur under very precise conditions of reaction, i.e. temperature, pressure, pH, concentration mineralizer or solvent solution. Thus, the hydrothermal crystallization method using inexpensive and chemicals easy to handle, can produce single or multicomponent oxides. The advantages of hydrothermal crystallization are the reduced energy costs due to the mild temperatures sufficient to achieve chemical reactions, less pollution, simplicity in the process equipment, and the fast rate of solid precipitation reactions (Vivekanandan & Kutty, 1989; Yoshimura & Suchanek, 1997; Riman et. al., 2002).

There are thousands of research works in the literature that include a vast experimental data related with the hydrothermal crystallization of metal oxides or inorganic compounds. The most popular among the metal oxides are those of perovskite oxides group, because of their wide application at the electronic industry. Hence, fundamental principles of crystallization of these materials are discussed in detail and compared with the present research experimental of the present authors regarding the crystallization of perovskite particles of SrTiO3, BaTiO3 and La1-XMXCr1-YNYO3.

## **2.4.3 Hydrothermal crystallization of perovskite derived oxides from gel precursors**

In this section, we focus on the hydrothermal crystallization of perovskite-type structure oxide materials, because this particular group of oxides offers stupendous functional properties for a wide number of applications, in particular electronic applications. Perovskite compounds have the general formula ABO3, and Perovskite is cubic structure in

Preparation of Selected Ceramic Compounds by

hydrothermal methods.

with the dissolution process.

TiOH3+ + H2O = Ti(OH)2

Controlled Crystallization Under Hydrothermal Conditions 217

Ti(OH)22+ + H2O = Ti(OH)3+ + H+ (7)

Ti(OH)3+ + H2O = Ti(OH)4(aq) + H+ (8)

Ti(OH)4(aq) = TiO2(s) + 2H2O (9)

BaCl2(aq) + TiCl4(aq) + 6NaOH(aq) = BaTiO3(s) + 6NaCl(aq) + 3H2O(l) (10)

Ba(OH)2(aq) + TiO2•xH2O(s) + NaOH(aq) = BaTiO3(s) + NaOH(aq) + xH2O(l) (11)

 Ba(OH)2(aq) + Ti(OH)4•xH2O(aq) + NaOH(aq) = BaTiO3(s) + NaOH(aq) + 4xH2O(l) (12) The chemical synthesis of BT particles has been studied over a broad range of experimental hydrothermal conditions; this compound is preferentially formed in highly concentrated alkaline solutions (pH > 10), as indicated by thermodynamic calculations in BT phase stability diagram proposed elsewhere (Lencka et. al., 1993; Riman et. al., 2002). The normal intervals of reaction time and temperature are 2–96 h and 120–250 ºC where the crystallization is normally achieved, thus, the preparation of this type of powder is carried out in typical stainless steel 304 Teflon-lined autoclaves. Additional parametric variable studies recently conducted accounts for other variables that also have an effect on the formation of the BT pure phase like the molar ratio of precursor alkaline (KOH or NaOH) solvent (Vivekanandan & Kutty, 1989; Lee et. al, 2003) and the reactants (Ba(OH)2 or titania source) (Moon et. al., 2003, Qi et. al., 2004), the Ba/Ti molar ratio of the raw materials (Wada et. al., 1995). Another specific studies recently conducted lead to remarkable analysis of the kinetic process related to the crystallization of BT particles and their growth (Walton et. al., 2001), and the control of the particle size and shape of the BT particles have been carried out by stirring the autoclave (Kubo et. al., 2009) and microwave assisted (Moreira et. al., 2008)

Hitherto, in accordance with the analyses conducted based on experimental observations, indicate that the crystallization mechanism of BT particles is dissolution-precipitation in nature, and this mechanism is mainly controlled by the dissolving rate of the reactant solid specie, namely titania (oxide or amorphous gel). Thus, the presence of a high concentration of OH- ions in the hydrothermal system is required to produce the hydrolysis of aqueous species, and these ions also seemed to act as catalysts to accelerate the transition from Ba-OH bonds, usually resulting on the crystallization of BT via precipitation that proceeds via the chemical reaction (13), resulting from the previous reaction (10–12) that is carried out during the earlier and intermediate stages of the hydrothermal treatment and are linked

 BaOH+(aq) + Ti(OH)40(aq) = BaTiO3(s) + 2H2O + H+(aq) (13) Once the solution is supersaturated due to dissolution of the precursor, precipitation of BT from the homogeneous solutions spontaneously occurs, yielding an abundant number of nuclei. Therefore, the dissolution of either titania hydrous amorphous or crystalline powders is the rate limiting step that controls the early stage of BT nucleation and growth

2+ + H+ (6)

nature which is shown in Figure 3, where the large A cations (Ba2+, Sr2+, Ca2+, Pb2+, La3+, Bi3+ and K+), but low in electric charge, are surrounded by 12 oxygens, whilst B ions relatively small in size (Ti4+, Zr4+, Sn4+, W6+, Nb5+, Mn3+, Mg2+) are coordinated by 6 oxygens. The crystallization of the most representative compounds in this group, namely BaTiO3 (BT), SrTiO3 (ST) and solid solutions of Ba1-XSrXTiO3 (BST) or PbZr1-XTiXO3 (PZT) has been extensively studied under hydrothermal conditions.

Fig. 3. Schematic representation of the unit lattice cell of the Perovskite cubic structure.

#### **2.4.3.1 Hydrothermal crystallization of perovskite barium and strontium titanate oxides**

BaTiO3 (BT) and SrTiO3 (ST) are the metal earth alkaline perovskite-like structures most prepared species under hydrothermal conditions, these compounds has similarity on the crystallization process as that proposed for PT. In terms of the chemical reaction equilibrium that can be produced on the hydrothermal system, and in accordance with the chemical reagent reactants that have been used the chemical reactions that are related with the crystallization of BT and ST powders are as follows:

Where M on chemical equations 3 and 4 is related to earth alkaline metals of the group IIA of the periodic table of elements, *viz.* Ca, Sr and Ba. The chemical reactions from 1 to 9 represent all the reaction equilibrium that are able to proceed in the system M-Ti-Na-H2O and these reactions serve to produce the crystallization process, because they stem from the principal chemical reactions 1–12 that have widely investigated on the crystallization of BT under hydrothermal alkaline conditions and can operate on the case of the ST compound as well.

$$\rm{H}\_{2}\rm{O}=\rm{H}^{\*}+\rm{OH}^{-}\tag{1}$$

$$\text{H}\_2\text{O}(\text{g}) = \text{H}\_2\text{O} \tag{2}$$

$$\rm{MTiO\_3(s) + H\_2O = M^{2+} + 2OH^- + TiO\_2(s)}\tag{3}$$

$$\mathrm{MOH^{\cdot}} = \mathrm{M^{2\cdot}} + \mathrm{OH^{\cdot}} \tag{4}$$

$$\text{Ti}4\* + \text{H}\_2\text{O} = \text{TiOH} \\ \text{H}\_2\* + \text{H}\_2\* \tag{5}$$

nature which is shown in Figure 3, where the large A cations (Ba2+, Sr2+, Ca2+, Pb2+, La3+, Bi3+ and K+), but low in electric charge, are surrounded by 12 oxygens, whilst B ions relatively small in size (Ti4+, Zr4+, Sn4+, W6+, Nb5+, Mn3+, Mg2+) are coordinated by 6 oxygens. The crystallization of the most representative compounds in this group, namely BaTiO3 (BT), SrTiO3 (ST) and solid solutions of Ba1-XSrXTiO3 (BST) or PbZr1-XTiXO3 (PZT) has been

Fig. 3. Schematic representation of the unit lattice cell of the Perovskite cubic structure.

**2.4.3.1 Hydrothermal crystallization of perovskite barium and strontium titanate oxides**  BaTiO3 (BT) and SrTiO3 (ST) are the metal earth alkaline perovskite-like structures most prepared species under hydrothermal conditions, these compounds has similarity on the crystallization process as that proposed for PT. In terms of the chemical reaction equilibrium that can be produced on the hydrothermal system, and in accordance with the chemical reagent reactants that have been used the chemical reactions that are related with the

Where M on chemical equations 3 and 4 is related to earth alkaline metals of the group IIA of the periodic table of elements, *viz.* Ca, Sr and Ba. The chemical reactions from 1 to 9 represent all the reaction equilibrium that are able to proceed in the system M-Ti-Na-H2O and these reactions serve to produce the crystallization process, because they stem from the principal chemical reactions 1–12 that have widely investigated on the crystallization of BT under hydrothermal alkaline conditions and can operate on the case of the ST compound as

H2O = H+ + OH- (1)

H2O(g) = H2O (2)

MOH+ = M2+ + OH- (4)

Ti4+ + H2O = TiOH3+ + H+ (5)

+ TiO2(s) (3)

extensively studied under hydrothermal conditions.

crystallization of BT and ST powders are as follows:

MTiO3(s) + H2O = M2+ + 2OH-

well.

$$\text{TiOH} 3^{\ast} + \text{H}\_{2}\text{O} = \text{Ti(OH)} 5^{\ast} + \text{H}^{\ast} \tag{6}$$

$$\text{Ti(OH)}\_{2}\text{2\*} + \text{H}\_{2}\text{O} = \text{Ti(OH)}\text{3\*} + \text{H}^{\*} \tag{7}$$

$$\text{Ti(OH)}\text{\textsuperscript{}+} + \text{H}\_{2}\text{O}=\text{Ti(OH)}\text{\textsuperscript{}}\text{(aq)} + \text{H}^\*\tag{8}$$

$$\text{Ti(OH)}\_{4}\text{(aq)} = \text{TiO}\_{2}\text{(s)} + 2\text{H}\_{2}\text{O} \tag{9}$$

$$\text{BaCl}\_2(\text{aq}) + \text{TiCl}\_4(\text{aq}) + 6\text{NaOH}(\text{aq}) = \text{BaTiO}\_3(\text{s}) + 6\text{NaCl}(\text{aq}) + 3\text{H}\_2\text{O}(\text{l})\tag{10}$$

$$\text{Ba(OH)}\_{2}\text{(aq)} + \text{TiO}\_{2} \bullet \text{H}\_{2}\text{O(s)} + \text{NaOH(aq)} = \text{BaTiO}\_{3}\text{(s)} + \text{NaOH(aq)} + \text{xH}\_{2}\text{O(l)}\tag{11}$$

$$\text{Ba(OH)}\_{2}\text{(aq)} + \text{Ti(OH)}\_{4}\bullet \text{xH}\_{2}\text{O(aq)} + \text{NaOH(aq)} = \text{BaTiO}\_{2}\text{(s)} + \text{NaOH(aq)} + 4\text{xH}\_{2}\text{O(l)}\text{ (12)}$$

The chemical synthesis of BT particles has been studied over a broad range of experimental hydrothermal conditions; this compound is preferentially formed in highly concentrated alkaline solutions (pH > 10), as indicated by thermodynamic calculations in BT phase stability diagram proposed elsewhere (Lencka et. al., 1993; Riman et. al., 2002). The normal intervals of reaction time and temperature are 2–96 h and 120–250 ºC where the crystallization is normally achieved, thus, the preparation of this type of powder is carried out in typical stainless steel 304 Teflon-lined autoclaves. Additional parametric variable studies recently conducted accounts for other variables that also have an effect on the formation of the BT pure phase like the molar ratio of precursor alkaline (KOH or NaOH) solvent (Vivekanandan & Kutty, 1989; Lee et. al, 2003) and the reactants (Ba(OH)2 or titania source) (Moon et. al., 2003, Qi et. al., 2004), the Ba/Ti molar ratio of the raw materials (Wada et. al., 1995). Another specific studies recently conducted lead to remarkable analysis of the kinetic process related to the crystallization of BT particles and their growth (Walton et. al., 2001), and the control of the particle size and shape of the BT particles have been carried out by stirring the autoclave (Kubo et. al., 2009) and microwave assisted (Moreira et. al., 2008) hydrothermal methods.

Hitherto, in accordance with the analyses conducted based on experimental observations, indicate that the crystallization mechanism of BT particles is dissolution-precipitation in nature, and this mechanism is mainly controlled by the dissolving rate of the reactant solid specie, namely titania (oxide or amorphous gel). Thus, the presence of a high concentration of OH- ions in the hydrothermal system is required to produce the hydrolysis of aqueous species, and these ions also seemed to act as catalysts to accelerate the transition from Ba-OH bonds, usually resulting on the crystallization of BT via precipitation that proceeds via the chemical reaction (13), resulting from the previous reaction (10–12) that is carried out during the earlier and intermediate stages of the hydrothermal treatment and are linked with the dissolution process.

$$\rm BaOH^{\*} \text{(aq)} + Ti \text{(OH)} \mu \text{(aq)} = BaTiO\_{3} \text{(s)} + 2H\_{2}O + H^{\*} \text{(aq)}\tag{13}$$

Once the solution is supersaturated due to dissolution of the precursor, precipitation of BT from the homogeneous solutions spontaneously occurs, yielding an abundant number of nuclei. Therefore, the dissolution of either titania hydrous amorphous or crystalline powders is the rate limiting step that controls the early stage of BT nucleation and growth

Preparation of Selected Ceramic Compounds by

**oxides using mineral precursors** 

propose an economical effective processing method.

Controlled Crystallization Under Hydrothermal Conditions 219

Fig. 4. Schematic sketch of the dissolution/precipitation mechanism that conduces to the

**2.4.3.2 Single step hydrothermal crystallization of perovskite strontium and barium** 

In accordance to the former literature, in the last decade, the interest for using pure mineral species to produce pure synthetic inorganic compounds has been rising. In general, the synthesis of SrTiO3 or BaTiO3 has been broadly conducted by employing Ti(OH)4 gel (Tigel), while several strontium soluble salts had been used as a source of Sr2+ ions. However, the challenge for employing low-grade chemical reagent precursors for preparing ABO3 particles has not been considered yet. The approach for employing a pure mineral ore, like celestite (SrSO4) or barite (BaSO4) with a low grade of impurities; to produce strontium or barium compounds under hydrothermal conditions, was considered based on the analysis of the proposed chemical methods for producing functional ceramic materials and the former information related to the ionic substitution on mineral species, which are analogous to the mechanistic principles of hydrothermal crystallization discussed in this section. The employment of a low cost precursor may provide an additional advantage in order to

Recently, the present authors have conducted an exhaustive efforts to develop a simple single step reaction method for the preparation of SrTiO3, which involves the employment of a mineral SrSO4 crystal plate (0.2 ± 0.0010 g, 6 ± 1 mm side and 2 ± 0.5 mm thick) with Ti(OH)4•4.5H2O gel (1 g, stoichiometric ratio Sr/Ti=1) under hydrothermal conditions, at various temperatures (150–250 ºC) for different reaction intervals (0.08–96 h) in KOH solutions with different concentrations (5–10 M). The hydrothermal treatments were carried out in stainless steel micro-autoclaves lined with Teflon with a filling volume ratio of 50% of the total inner volume (30 ml). This process involves a complex solute dissolution stage because SrSO4 is chemically stable even at acid, neutral and mild basic conditions. However,

hydrothermal crystallization of BaTiO3 (Ecker et. al. 1996, Zhang et. al., 2004).

during the hydrothermal crystallization process, for anhydrous TiO2 precursor, Ti-O bonds must be broken via hydrolytic attack to for hydroxyl-titanium complexes (Ti(OH)X4-X) capable of dissolution and reaction with barium ions or complexes (Ba2+ or BaOH+) in solution to precipitate BT. Subsequently, the nuclei grow rapidly, resulting in an accelerated crystallization kinetics rate at the initial step of the process. In the solvent, the solute concentration decreases to below the supersaturation point as a result of this event, but remains sufficiently high for the particles to growth, avoiding a secondary nucleation. The particle coarsening proceeds during intermediate and final steps of the crystallization process, and mainly depends on the dissociation of the remaining reactants, or in some cases the separation of terminal organic groups that are liked to the Ti or Ba, e.g. acetiylacetone or acetate, these crystallization barriers serves to slow, if not halt, the kinetic rate in the final stage of the hydrothermal treatment (Eckert et. al., 1996). The crystallization studies indicates that the activation energy fro barium titanate crystallization under hydrothermal conditions varies on the range of 21-105 kJmol-1, and the value of activation energy specifically depends on the type of compound used as titania source (Eckert et. al.; 1996; Walton et. al., 2001), but these values of activation are within the range of activation energy values for chemical reaction that proceed on the liquid stage.

On the other hand, major attempts recently conducted to investigate the crystallization of SrTiO3 particles under hydrothermal conditions, have been designed by considering the fundamental principles derived from the synthesis of its related compounds, *viz.* BT or PT perovskite-like structure. A broad type of strontium titanate particle shapes, e.g. nanotube, dendrite, cuboidal and spherical; had been produced through optimizing the hydrothermal crystallization conditions by controlling parameter such as: the pH of the hydrothermal alkaline media in the range between 10–12 (Wendelbo et. al, 2006), reaction temperature and time. Furthermore, different chemical reagents including inorganic and organometallic had been employed as source of Sr2+ (Sr(NO3)2, SrCl2, Sr(OH)2•6H2O) and Ti4+ (TiCl4, TiO2•H2O, Ti(SO4)2, Ti(OC4H9)4, Ti(OPr*<sup>i</sup>* )4) for producing ST particles (Moon et. al., 1999; Zhang et. al., 2001; Wang et. al., 2009). Another factor that has been studied is related with the use of organic dispersants, namely polyvinyl alcohol, for controlling the crystallization of fine regular shaped ST particles. The dispersant that have large polymeric changes can operate as micelles that attract the hydrolysed ionic species that are formed during the hydrothermal treatment, this lead to and homogeneous nucleation and limited particle growth, resulting in the preparation of nanometer sized ST regular particles (Wei et. al., 2008). Regarding the mechanism correlated with the crystallization of ST particles, one model derived from further experimental data states that the dissolution-crystallization coupled with a second aggregative growth-recrystallization mechanisms are related with the bulk crystallization stage of ST particles under alkaline hydrothermal conditions, but is similar to that proposed for the formation of BT particles (Fig. 4). The second mechanism is achieved when heterogeneous nucleation promotes the formation of ST particles, the surface of the TiO2 particles acts as the precipitation site for ST nuclei, and when the size of TiO2 particles is reduced by the progressive dissolution, the aggregation growth of the ST particles proceeds in the reaction media. This phenomenon produces a marked particle agglomeration and also the coarsening of the particles can be promoted by the Ostwald recrystallization mechanism (Zhang et. al., 2004).

during the hydrothermal crystallization process, for anhydrous TiO2 precursor, Ti-O bonds must be broken via hydrolytic attack to for hydroxyl-titanium complexes (Ti(OH)X4-X) capable of dissolution and reaction with barium ions or complexes (Ba2+ or BaOH+) in solution to precipitate BT. Subsequently, the nuclei grow rapidly, resulting in an accelerated crystallization kinetics rate at the initial step of the process. In the solvent, the solute concentration decreases to below the supersaturation point as a result of this event, but remains sufficiently high for the particles to growth, avoiding a secondary nucleation. The particle coarsening proceeds during intermediate and final steps of the crystallization process, and mainly depends on the dissociation of the remaining reactants, or in some cases the separation of terminal organic groups that are liked to the Ti or Ba, e.g. acetiylacetone or acetate, these crystallization barriers serves to slow, if not halt, the kinetic rate in the final stage of the hydrothermal treatment (Eckert et. al., 1996). The crystallization studies indicates that the activation energy fro barium titanate crystallization under hydrothermal conditions varies on the range of 21-105 kJmol-1, and the value of activation energy specifically depends on the type of compound used as titania source (Eckert et. al.; 1996; Walton et. al., 2001), but these values of activation are within the range of activation energy

On the other hand, major attempts recently conducted to investigate the crystallization of SrTiO3 particles under hydrothermal conditions, have been designed by considering the fundamental principles derived from the synthesis of its related compounds, *viz.* BT or PT perovskite-like structure. A broad type of strontium titanate particle shapes, e.g. nanotube, dendrite, cuboidal and spherical; had been produced through optimizing the hydrothermal crystallization conditions by controlling parameter such as: the pH of the hydrothermal alkaline media in the range between 10–12 (Wendelbo et. al, 2006), reaction temperature and time. Furthermore, different chemical reagents including inorganic and organometallic had been employed as source of Sr2+ (Sr(NO3)2, SrCl2, Sr(OH)2•6H2O) and Ti4+ (TiCl4, TiO2•H2O,

2001; Wang et. al., 2009). Another factor that has been studied is related with the use of organic dispersants, namely polyvinyl alcohol, for controlling the crystallization of fine regular shaped ST particles. The dispersant that have large polymeric changes can operate as micelles that attract the hydrolysed ionic species that are formed during the hydrothermal treatment, this lead to and homogeneous nucleation and limited particle growth, resulting in the preparation of nanometer sized ST regular particles (Wei et. al., 2008). Regarding the mechanism correlated with the crystallization of ST particles, one model derived from further experimental data states that the dissolution-crystallization coupled with a second aggregative growth-recrystallization mechanisms are related with the bulk crystallization stage of ST particles under alkaline hydrothermal conditions, but is similar to that proposed for the formation of BT particles (Fig. 4). The second mechanism is achieved when heterogeneous nucleation promotes the formation of ST particles, the surface of the TiO2 particles acts as the precipitation site for ST nuclei, and when the size of TiO2 particles is reduced by the progressive dissolution, the aggregation growth of the ST particles proceeds in the reaction media. This phenomenon produces a marked particle agglomeration and also the coarsening of the particles can be promoted by the Ostwald

)4) for producing ST particles (Moon et. al., 1999; Zhang et. al.,

values for chemical reaction that proceed on the liquid stage.

Ti(SO4)2, Ti(OC4H9)4, Ti(OPr*<sup>i</sup>*

recrystallization mechanism (Zhang et. al., 2004).

Fig. 4. Schematic sketch of the dissolution/precipitation mechanism that conduces to the hydrothermal crystallization of BaTiO3 (Ecker et. al. 1996, Zhang et. al., 2004).

#### **2.4.3.2 Single step hydrothermal crystallization of perovskite strontium and barium oxides using mineral precursors**

In accordance to the former literature, in the last decade, the interest for using pure mineral species to produce pure synthetic inorganic compounds has been rising. In general, the synthesis of SrTiO3 or BaTiO3 has been broadly conducted by employing Ti(OH)4 gel (Tigel), while several strontium soluble salts had been used as a source of Sr2+ ions. However, the challenge for employing low-grade chemical reagent precursors for preparing ABO3 particles has not been considered yet. The approach for employing a pure mineral ore, like celestite (SrSO4) or barite (BaSO4) with a low grade of impurities; to produce strontium or barium compounds under hydrothermal conditions, was considered based on the analysis of the proposed chemical methods for producing functional ceramic materials and the former information related to the ionic substitution on mineral species, which are analogous to the mechanistic principles of hydrothermal crystallization discussed in this section. The employment of a low cost precursor may provide an additional advantage in order to propose an economical effective processing method.

Recently, the present authors have conducted an exhaustive efforts to develop a simple single step reaction method for the preparation of SrTiO3, which involves the employment of a mineral SrSO4 crystal plate (0.2 ± 0.0010 g, 6 ± 1 mm side and 2 ± 0.5 mm thick) with Ti(OH)4•4.5H2O gel (1 g, stoichiometric ratio Sr/Ti=1) under hydrothermal conditions, at various temperatures (150–250 ºC) for different reaction intervals (0.08–96 h) in KOH solutions with different concentrations (5–10 M). The hydrothermal treatments were carried out in stainless steel micro-autoclaves lined with Teflon with a filling volume ratio of 50% of the total inner volume (30 ml). This process involves a complex solute dissolution stage because SrSO4 is chemically stable even at acid, neutral and mild basic conditions. However,

Preparation of Selected Ceramic Compounds by

Controlled Crystallization Under Hydrothermal Conditions 221

Fig. 5. Morphologies of ST particles obtained after hydrothermal treatments of SrSO4 crystal plates, carried out at temperatures of (a) 200 ºC, (b) 250 ºC; and (c) EDX spectra of ST (*i*) peanut-like and (*ii*) cubic particles shown in Fig. 3b (Rangel-Hernandez et. al., 2009).

 Fig. 6. SrTiO3 particles hydrothermally produced using SrSO4 powders at 250 ºC in 5 M KOH solution for intervals of (a) 0.08 and (b) 24 h (Rangel-Hernandez et. al., 2009).

The bulk crystallization process of ST particles is dissolution-precipitation in nature, and the reaction path involving this mechanism that produces nuclei formation and crystal growth, is represented physically in the Figures 7a and 7b. These Figures gave the set of the original

the use of a mineral single crystal favours to analyse systematically the effect of the dissolution of this reactant coupled with the Ti(OH)4 dehydration and serve to control the synthesis and crystallization kinetics of ST particles. Thus, the complete crystallization of ST particles in a single step of reaction occurred with the complete dissolution of the SrSO4 crystal obtained at 250 ºC for 96 h in a 5 M KOH solution, resulting in the formation of SrTiO3 particles with two different shapes (peanut-like and cubic) as is shown in Figure 4. The parameter that has a marked effect on control the particle size and morphology is the temperature rather than the interval of reaction. ST particles having a bimodal size distribution (0.5–1 m and 0.2–0.5 m) were prepared at different temperatures for 96 h in a 5 M KOH solution from SrSO4 crystals. The ST powders produced at mild temperatures (200 ºC) were constituted by a large amount of agglomerated peanut-like shape irregular particles (Fig. 5a), together with a small amount of bulky crystals with a regular pseudocubic shape (1–4 m). In contrast, at high temperature (250 ºC), the pseudo-cubic (3–6 m) ST particles exhibited a coarsening, while the amount of peanut-like shape (length = 1–5 m and width = 0.5–1 m) particles was significantly reduced (Fig. 5b). Hence, the coarsening of pseudo-cubic shaped particles is attributed to the Ostwald ripening particle growth mechanism (Peterson & Slamovich, 1999). Indeed, the growth of the aggregated particles is due to a recrystallization mechanism, which involves the dissolution of the primary small ST peanut-like particles produced at lower temperatures (< 200 ºC). This process is promoted due to lower chemical stability that exhibits the strontium titanate oxide in highly concentrated (> 5 M) KOH solvent solutions at 250 ºC. Another point that deserves to be emphasized is that related with the structural features of either peanut-like or cubic-like shaped particles. EDX spectra obtained on both particles did not show a marked difference on the molar Sr/Ti ratio, because the particles contain similar Sr and Ti amounts as it is suggested by the corresponding EDX spectrum in Fig. 5c (Rangel-Hernandez et. al., 2009).

In contrast, the increase on the surface area of the precursor SrSO4 produces the control of morphology and size homogeneity for the ST particles. In general, the use of SrSO4 powders with a narrow particle size distribution in the range from 25–38 m, favoured the crystallization of very fine cubic ST particles, even at very short reaction intervals (0.08 h, Fig. 6a). However, ST powders with a bimodal size distribution, consisting of a large amount of pseudo cubic particles (0.75 m) and a small quantity of large cubic particles (1.5 m) were produced during intermediate reaction intervals between 3-12 h. In contrast, at the longest reaction interval of 24 h, the preferential formation of very fine pseudospherical ST particles (average size of 0.4 m) was observed (Fig. 6b). The formation of the cubic particles is due to the crystallographic habit growth, which proceeded when the alkaline solvent reaches a supersaturation steady-state with the ionic species Sr(OH)+ and Ti(OH)40. Hence, a massive homogeneous nuclei formation and fast growth of the cubic particles proceeds in the hydrothermal system during intermediate reaction intervals. However, the ST cubic particles underwent into a preferential dissolution by increasing the reaction interval at long reaction period (24 h). The ST particle dissolution proceeds at high-energy faceted edges on the fine cubic particles, because of the relative low chemical stability of the ST perovskite structure in the concentrated alkaline media. This fact leads to further dissolution of the cubic particles and the recrystallization of the pseudospherical shaped particles from the solution.

the use of a mineral single crystal favours to analyse systematically the effect of the dissolution of this reactant coupled with the Ti(OH)4 dehydration and serve to control the synthesis and crystallization kinetics of ST particles. Thus, the complete crystallization of ST particles in a single step of reaction occurred with the complete dissolution of the SrSO4 crystal obtained at 250 ºC for 96 h in a 5 M KOH solution, resulting in the formation of SrTiO3 particles with two different shapes (peanut-like and cubic) as is shown in Figure 4. The parameter that has a marked effect on control the particle size and morphology is the temperature rather than the interval of reaction. ST particles having a bimodal size distribution (0.5–1 m and 0.2–0.5 m) were prepared at different temperatures for 96 h in a 5 M KOH solution from SrSO4 crystals. The ST powders produced at mild temperatures (200 ºC) were constituted by a large amount of agglomerated peanut-like shape irregular particles (Fig. 5a), together with a small amount of bulky crystals with a regular pseudocubic shape (1–4 m). In contrast, at high temperature (250 ºC), the pseudo-cubic (3–6 m) ST particles exhibited a coarsening, while the amount of peanut-like shape (length = 1–5 m and width = 0.5–1 m) particles was significantly reduced (Fig. 5b). Hence, the coarsening of pseudo-cubic shaped particles is attributed to the Ostwald ripening particle growth mechanism (Peterson & Slamovich, 1999). Indeed, the growth of the aggregated particles is due to a recrystallization mechanism, which involves the dissolution of the primary small ST peanut-like particles produced at lower temperatures (< 200 ºC). This process is promoted due to lower chemical stability that exhibits the strontium titanate oxide in highly concentrated (> 5 M) KOH solvent solutions at 250 ºC. Another point that deserves to be emphasized is that related with the structural features of either peanut-like or cubic-like shaped particles. EDX spectra obtained on both particles did not show a marked difference on the molar Sr/Ti ratio, because the particles contain similar Sr and Ti amounts as it is suggested by the corresponding EDX spectrum in Fig. 5c (Rangel-Hernandez et. al., 2009).

In contrast, the increase on the surface area of the precursor SrSO4 produces the control of morphology and size homogeneity for the ST particles. In general, the use of SrSO4 powders with a narrow particle size distribution in the range from 25–38 m, favoured the crystallization of very fine cubic ST particles, even at very short reaction intervals (0.08 h, Fig. 6a). However, ST powders with a bimodal size distribution, consisting of a large amount of pseudo cubic particles (0.75 m) and a small quantity of large cubic particles (1.5 m) were produced during intermediate reaction intervals between 3-12 h. In contrast, at the longest reaction interval of 24 h, the preferential formation of very fine pseudospherical ST particles (average size of 0.4 m) was observed (Fig. 6b). The formation of the cubic particles is due to the crystallographic habit growth, which proceeded when the alkaline solvent reaches a supersaturation steady-state with the ionic species Sr(OH)+ and Ti(OH)40. Hence, a massive homogeneous nuclei formation and fast growth of the cubic particles proceeds in the hydrothermal system during intermediate reaction intervals. However, the ST cubic particles underwent into a preferential dissolution by increasing the reaction interval at long reaction period (24 h). The ST particle dissolution proceeds at high-energy faceted edges on the fine cubic particles, because of the relative low chemical stability of the ST perovskite structure in the concentrated alkaline media. This fact leads to further dissolution of the cubic particles and the recrystallization of the pseudo-

spherical shaped particles from the solution.

Fig. 5. Morphologies of ST particles obtained after hydrothermal treatments of SrSO4 crystal plates, carried out at temperatures of (a) 200 ºC, (b) 250 ºC; and (c) EDX spectra of ST (*i*) peanut-like and (*ii*) cubic particles shown in Fig. 3b (Rangel-Hernandez et. al., 2009).

Fig. 6. SrTiO3 particles hydrothermally produced using SrSO4 powders at 250 ºC in 5 M KOH solution for intervals of (a) 0.08 and (b) 24 h (Rangel-Hernandez et. al., 2009).

The bulk crystallization process of ST particles is dissolution-precipitation in nature, and the reaction path involving this mechanism that produces nuclei formation and crystal growth, is represented physically in the Figures 7a and 7b. These Figures gave the set of the original

Preparation of Selected Ceramic Compounds by

precipitate because thermodynamically are more stable than BT.

Controlled Crystallization Under Hydrothermal Conditions 223

particle crystallization process, but the reaction was conducted by increasing the concentration of the KOH solution up to 10 M, and the complete dissolution of the BaSO4 crystal occurred for large reaction intervals of 144 h above 200 ºC. Optimum steady-state supersaturation of the solvent media with Ba2+ and Ti4+ ions that achieve the conditions for homogeneous nucleation of BT particles were found to proceed at temperatures below 200 ºC (Figure 8). Fine BT particles with pseudo-cubic, star-like and dendrite shapes were preferentially formed in 10 KOH. At sever treatment conditions of temperature 250 ºC for 144 h, the BaSO4 crystals were completely dissolved, and the formation of reaction byproducts Ba2TiO4 (needle crystals in Figure 8c) and a few amount of TiO2 simultaneously occurred with the excessive growth of polyhedral aggregated BT crystals. This particular behaviour is attributed to the differences in the chemical reactivity of the alkaline solution resulting in different dissolving rates for solid species, this can shift the reaction equilibrium coupled with the saturation conditions. Therefore, other crystalline phases are able to

Fig. 7. Aspects of (a) original SrSO4 crystal plate embedded in Ti-gel, and (b) the partially reacted SrSO4 crystal and Ti-gel at 200 ºC for 6 h in a 5 M KOH solution. (c) Consumption curves () of SrSO4 crystal vs. the reaction interval obtained at 250 ºC in a 5 M KOH

solution, () without Ti-gel and () with Ti-gel additions (Rangel-Hernandez et. al., 2009).

SrSO4 crystal plate embedded in the Ti(OH)4•4.5H2O gel prior the hydrothermal treatment and after 6 h of reaction at 200 ºC. In terms of fundamental chemistry the reaction path way occurs trough the chemical equations 14 and 15:

$$\begin{aligned} \text{SrSO}\_4\text{(crystal)} &+ \text{Ti(OH)}\_4 \bullet 4.5\text{H}\_2\text{O(gel)} + 2\text{KOH} \\ &= \text{SrOH}^+\text{(aq)} + \text{K}\_2\text{SO}\_4\text{(aq)} + \text{Ti(OH)}\_4^{0}\text{(aq)} + 4.5\text{H}\_2\text{O} + \text{OH}^-\text{(aq)} \end{aligned} \tag{14}$$

$$\text{SrOH}^\*\text{(aq)} + \text{Ti(OH)}\mathbb{A}^\circ\text{(aq)} + \text{OH}\text{-(aq)} = \text{SrTiO}\_3\text{(s)} + 3\text{H}\_2\text{O} \tag{15}$$

Macroscopic aspects related with the single-step synthesis process were revealed during the early stages of the reaction at 200 ºC for 6 h in a 5 M KOH solution (Fig. 7b). In general, the reaction gradually proceeds by the SrSO4 crystal dissolution, the consumption of SrSO<sup>4</sup> occurs locally on those crystal surfaces exposed to the alkaline fluid and also on those crystal surfaces in contact with the Ti-gel. This phenomenon agrees with those results previously reported and provided evidences of the solubility of SrSO4 compound in alkaline hydrothermal solutions (5 and 10 M NaOH) (Rendón-Angeles et. al., 2006). Hence, the continuous dissolution of the SrSO4 crystal must yield the formation of SrOH+ species (Eq. 14), because of the saturated alkaline conditions of the solvent. Indeed, the solubility of the Ti-gel is low in high concentrated KOH (> 0.1 M) solutions (Wang et. al., 2009, Rangel-Hernandez et. al., 2009). However, the consumption of the SrSO4 crystal dissolution () was markedly reduced due to the presence of the Ti-gel (Fig. 7c). This fact let us to conclude that the reactivity of Ti-gel is the limiting rate variable for crystallization, because reduces the dissolution rate of the SrSO4 phase. Thus, the equation (4) is proposed for the crystallization of ST and is similar to that discussed at the early part of the present section for PT and BT compounds. The nucleation and growth processes of the ST particles occur when the alkaline (KOH) solution reaches a supersaturation steady-state of the species Sr(OH)+ and Ti(OH)40. The ST particles precipitation locally occurred at Ti-gel surface and between the spaces with the remaining SrSO4 crystal, because the mass transport was limited during the treatments, which were conducted under neither agitation nor thermal gradients. Additionally, kinetic data obtained from SrSO4 consumption curves depicted that the activation energy required for the synthesis of SrTiO3 powders from the complete consumption of an SrSO4 crystal plate under hydrothermal conditions, is 27.9 kJ mol−1. One particular fact that must be emphasised is related with the incorporation of the major impurities on the ST particles during the synthesis process, namely Ba (5.7 wt.%) and CO<sup>3</sup> -2 (0.6 wt.%), contained in the SrSO4 crystals, does not proceed during the crystallization event because neither the presence of barium nor carbonate compounds were found on the compositional analysis conducted by EDX and DRX techniques in the ST particles.

On the other hand, the crystallization of perovskite related compounds like ST or BT via the transformation of sulphate alkaline earth metal mineral species is hinder for the barite mineral. The crystalline transformation of BaSO4 to BaTiO3 was studied using the single step reaction route aforementioned; and the high chemical stability (low solubility) of the orthorhombic structure limited the dissolution of the barite crystals under hydrothermal conditions even in highly concentrated KOH solutions (>5 M). The reactivity of the orthorhombic structure with the alkaline solvent can be influenced by the size of the alkaline earth metal ion incorporated in the structure, these results are in agreement with those early reported on the transformation of Ca, Sr and Ba-chlorapatite crystals into their respective hydroxyapatite species (Rendón-Angeles et. al., 2000a). This factor reduces the rate of the BT

SrSO4 crystal plate embedded in the Ti(OH)4•4.5H2O gel prior the hydrothermal treatment and after 6 h of reaction at 200 ºC. In terms of fundamental chemistry the reaction path way

 SrOH+(aq) + Ti(OH)40(aq) + OH–(aq) = SrTiO3(s) + 3H2O (15) Macroscopic aspects related with the single-step synthesis process were revealed during the early stages of the reaction at 200 ºC for 6 h in a 5 M KOH solution (Fig. 7b). In general, the reaction gradually proceeds by the SrSO4 crystal dissolution, the consumption of SrSO<sup>4</sup> occurs locally on those crystal surfaces exposed to the alkaline fluid and also on those crystal surfaces in contact with the Ti-gel. This phenomenon agrees with those results previously reported and provided evidences of the solubility of SrSO4 compound in alkaline hydrothermal solutions (5 and 10 M NaOH) (Rendón-Angeles et. al., 2006). Hence, the continuous dissolution of the SrSO4 crystal must yield the formation of SrOH+ species (Eq. 14), because of the saturated alkaline conditions of the solvent. Indeed, the solubility of the Ti-gel is low in high concentrated KOH (> 0.1 M) solutions (Wang et. al., 2009, Rangel-Hernandez et. al., 2009). However, the consumption of the SrSO4 crystal dissolution () was markedly reduced due to the presence of the Ti-gel (Fig. 7c). This fact let us to conclude that the reactivity of Ti-gel is the limiting rate variable for crystallization, because reduces the dissolution rate of the SrSO4 phase. Thus, the equation (4) is proposed for the crystallization of ST and is similar to that discussed at the early part of the present section for PT and BT compounds. The nucleation and growth processes of the ST particles occur when the alkaline (KOH) solution reaches a supersaturation steady-state of the species Sr(OH)+ and Ti(OH)40. The ST particles precipitation locally occurred at Ti-gel surface and between the spaces with the remaining SrSO4 crystal, because the mass transport was limited during the treatments, which were conducted under neither agitation nor thermal gradients. Additionally, kinetic data obtained from SrSO4 consumption curves depicted that the activation energy required for the synthesis of SrTiO3 powders from the complete consumption of an SrSO4 crystal plate under hydrothermal conditions, is 27.9 kJ mol−1. One particular fact that must be emphasised is related with the incorporation of the major impurities on the ST particles during the synthesis process, namely Ba (5.7 wt.%) and CO<sup>3</sup>

(0.6 wt.%), contained in the SrSO4 crystals, does not proceed during the crystallization event because neither the presence of barium nor carbonate compounds were found on the

On the other hand, the crystallization of perovskite related compounds like ST or BT via the transformation of sulphate alkaline earth metal mineral species is hinder for the barite mineral. The crystalline transformation of BaSO4 to BaTiO3 was studied using the single step reaction route aforementioned; and the high chemical stability (low solubility) of the orthorhombic structure limited the dissolution of the barite crystals under hydrothermal conditions even in highly concentrated KOH solutions (>5 M). The reactivity of the orthorhombic structure with the alkaline solvent can be influenced by the size of the alkaline earth metal ion incorporated in the structure, these results are in agreement with those early reported on the transformation of Ca, Sr and Ba-chlorapatite crystals into their respective hydroxyapatite species (Rendón-Angeles et. al., 2000a). This factor reduces the rate of the BT

compositional analysis conducted by EDX and DRX techniques in the ST particles.

SrOH aq K SO aq Ti OH aq 4.5H O OH aq (14)

0 2 4 4 2


occurs trough the chemical equations 14 and 15:

4 2 4

SrSO crystal Ti OH •4.5H O gel 2KOH

particle crystallization process, but the reaction was conducted by increasing the concentration of the KOH solution up to 10 M, and the complete dissolution of the BaSO4 crystal occurred for large reaction intervals of 144 h above 200 ºC. Optimum steady-state supersaturation of the solvent media with Ba2+ and Ti4+ ions that achieve the conditions for homogeneous nucleation of BT particles were found to proceed at temperatures below 200 ºC (Figure 8). Fine BT particles with pseudo-cubic, star-like and dendrite shapes were preferentially formed in 10 KOH. At sever treatment conditions of temperature 250 ºC for 144 h, the BaSO4 crystals were completely dissolved, and the formation of reaction byproducts Ba2TiO4 (needle crystals in Figure 8c) and a few amount of TiO2 simultaneously occurred with the excessive growth of polyhedral aggregated BT crystals. This particular behaviour is attributed to the differences in the chemical reactivity of the alkaline solution resulting in different dissolving rates for solid species, this can shift the reaction equilibrium coupled with the saturation conditions. Therefore, other crystalline phases are able to precipitate because thermodynamically are more stable than BT.

Fig. 7. Aspects of (a) original SrSO4 crystal plate embedded in Ti-gel, and (b) the partially reacted SrSO4 crystal and Ti-gel at 200 ºC for 6 h in a 5 M KOH solution. (c) Consumption curves () of SrSO4 crystal vs. the reaction interval obtained at 250 ºC in a 5 M KOH solution, () without Ti-gel and () with Ti-gel additions (Rangel-Hernandez et. al., 2009).

Preparation of Selected Ceramic Compounds by

no traces of Ba are visible (Fig. 9d).

of the ST particles shown in (b).

Controlled Crystallization Under Hydrothermal Conditions 225

were derived due to the increase of reaction temperature, a mixture of fine particles resembling cubic and start-like shapes were formed at mild temperatures (150 and 200 ºC) for 12 h, increasing the temperature at 250 ºC bulky aggregated cubic ST particles (20 m size) were crystallized (Figure 9). During the hydrothermal crystallization of the ST compound the incorporation of the Ba2+ ions was avoided to occur in the structure of ST particles, this is supported by the EDX spectra of the cubic and star-like ST particles where

Two main factors are associated with; one is related to the marked compositional gradient differences of the metal ions Sr2+ and Ba2+ produced in the solvent media during mineral dissolution, coupled with the stability of Ba2+ to form complexes ions in alkaline solutions (BaOH+, Eq. 4). Indeed, wet chemical analyses of the remaining solutions after the treatments gave evidences that support the above inference, because the Ba2+ ions content gradually increased as far as the complete mineral dissolution was concluded for 96 h at 250 ºC. The content of Ba determined by in the solution was 20.0 ± 0.8 wt%, which is nearly the same, measured in the original barite-celestite crystals (20.8 wt%). The ST particles crystallization proceeds by the same model established previously on the case of the transformation of high pure celestite crystal to ST particles, the control of morphology and

Fig. 9. SrTiO3 particles crystallized from barite-celestite mineral plates in a 5 M KOH solution for 12 h at different temperatures of (a) 150, (b) 200 and (c) 250 ºC. (d) EDX spectra

#### **2.4.3.3 Elimination of mineral impurities during the synthesis of strontium titanate particles under hydrothermal conditions**

Regarding the presence of high content of impurities during the crystallization of ST particles from mineral celestite ores. The experimental research work conducted recently by the authors of the present chapter; clearly contribute to demonstrate that refining of the major metal ion contained in the mineral ore proceeds during the hydrothermal crystallization of ST powders. Thus, the chemical stability of the barite-celestite mineral specie was investigated to elucidate the feasibility to conduct simultaneously the crystallization of ST particles and the release of major impurities such as Ba, and the preparation of BT powders as well. In particular, the barite-celestite mineral consists in a solid solution with a general chemical formula of Sr0.70Ba0.30SO4, which was determined by wet chemistry using ICP, the total content of major constituents SrSO4 and BaSO4 was 64.7 wt.% and 35.3 wt.%, respectively. Additional microprobe wavelength X-ray diffraction analyses conducted by scanning electron microscope observations indicated that the microstructure of the mineral consists of two major solid solutions, one rich in strontium Sr0.95Ba0.05SO4 and a less amount of an intermediate Sr0.75Ba0.25SO4. The hydrothermal treatments of barite-celestite crystal plate samples were conducted in accordance with the experimental procedure explained before for high pure celestite crystals.

Fig. 8. BaTiO3 particles crystallized from barite mineral plates in a 10 M KOH solution for 24 h at temperatures of (a) 150, (b) 200 and (c) 250 ºC.

The synthesis of ST particles was preferentially obtained using the barite-celestite crystals in a feedstock alkaline 5 M KOH solution; the transformation of the mineral specie into the perovskite oxide is strongly affected by the temperature and concentration of the solvent media rather than the reaction interval. Marked morphological and particle size differences

**2.4.3.3 Elimination of mineral impurities during the synthesis of strontium titanate** 

Regarding the presence of high content of impurities during the crystallization of ST particles from mineral celestite ores. The experimental research work conducted recently by the authors of the present chapter; clearly contribute to demonstrate that refining of the major metal ion contained in the mineral ore proceeds during the hydrothermal crystallization of ST powders. Thus, the chemical stability of the barite-celestite mineral specie was investigated to elucidate the feasibility to conduct simultaneously the crystallization of ST particles and the release of major impurities such as Ba, and the preparation of BT powders as well. In particular, the barite-celestite mineral consists in a solid solution with a general chemical formula of Sr0.70Ba0.30SO4, which was determined by wet chemistry using ICP, the total content of major constituents SrSO4 and BaSO4 was 64.7 wt.% and 35.3 wt.%, respectively. Additional microprobe wavelength X-ray diffraction analyses conducted by scanning electron microscope observations indicated that the microstructure of the mineral consists of two major solid solutions, one rich in strontium Sr0.95Ba0.05SO4 and a less amount of an intermediate Sr0.75Ba0.25SO4. The hydrothermal treatments of barite-celestite crystal plate samples were conducted in accordance with the

Fig. 8. BaTiO3 particles crystallized from barite mineral plates in a 10 M KOH solution for 24

The synthesis of ST particles was preferentially obtained using the barite-celestite crystals in a feedstock alkaline 5 M KOH solution; the transformation of the mineral specie into the perovskite oxide is strongly affected by the temperature and concentration of the solvent media rather than the reaction interval. Marked morphological and particle size differences

h at temperatures of (a) 150, (b) 200 and (c) 250 ºC.

experimental procedure explained before for high pure celestite crystals.

**particles under hydrothermal conditions** 

were derived due to the increase of reaction temperature, a mixture of fine particles resembling cubic and start-like shapes were formed at mild temperatures (150 and 200 ºC) for 12 h, increasing the temperature at 250 ºC bulky aggregated cubic ST particles (20 m size) were crystallized (Figure 9). During the hydrothermal crystallization of the ST compound the incorporation of the Ba2+ ions was avoided to occur in the structure of ST particles, this is supported by the EDX spectra of the cubic and star-like ST particles where no traces of Ba are visible (Fig. 9d).

Fig. 9. SrTiO3 particles crystallized from barite-celestite mineral plates in a 5 M KOH solution for 12 h at different temperatures of (a) 150, (b) 200 and (c) 250 ºC. (d) EDX spectra of the ST particles shown in (b).

Two main factors are associated with; one is related to the marked compositional gradient differences of the metal ions Sr2+ and Ba2+ produced in the solvent media during mineral dissolution, coupled with the stability of Ba2+ to form complexes ions in alkaline solutions (BaOH+, Eq. 4). Indeed, wet chemical analyses of the remaining solutions after the treatments gave evidences that support the above inference, because the Ba2+ ions content gradually increased as far as the complete mineral dissolution was concluded for 96 h at 250 ºC. The content of Ba determined by in the solution was 20.0 ± 0.8 wt%, which is nearly the same, measured in the original barite-celestite crystals (20.8 wt%). The ST particles crystallization proceeds by the same model established previously on the case of the transformation of high pure celestite crystal to ST particles, the control of morphology and

Preparation of Selected Ceramic Compounds by

powders under hydrothermal conditions.

(Rivas-Vázquez et. al., 2004, 2006; Rendón-Angeles et. al., 2009).

Controlled Crystallization Under Hydrothermal Conditions 227

ºC) for a reaction interval between 0.5 and 2 h. After the treatments, the precipitates were well washed with distilled water, decanted and then dried in an oven at 100 ºC overnight

Fig. 10. Typical Hastelloy C type lined microautoclave and scheme of the cross section of the vessel, which was employed for conducting the synthesis of perovskite lanthanum chromite

The minimum temperature for conducting the crystallization of the pure LC powders was 375 ºC for a reaction interval of 1 h. This parameter, however, is mainly affected by the incorporation of the metal ions in both A and B sites of the perovskite structure ABO3, and also with the amount of metal dopant ion inserted. Thus, the crystallization of powders corresponding to the solid solutions La1-XMXCr1-YNYO3 (where M = Ca or Sr and N = Al or Ni) was observed to proceeds at temperatures above 400 ºC, when only a 10 mol% of Ca2+ or Sr2+ were partially substituting La3+ site. Thus, in the particular case of the La0.9Ca0.1CrO3 and La0.8Ca0.2CrO3 solid solutions, the structural analyses in Figure 11a depicted that at a constant reaction interval 1 h, the powders of La0.9Ca0.1CrO3 were produced at a temperature of 400 ºC without contaminant formation, in contrast with the La0.8Ca0.2CrO3 powders that were produced together with a marked amount of secondary crystalline phases of La(OH)3, CrOOH, and CaCrO4, these phases were eliminated by increasing the reaction temperature up to 425 ºC resulting in the crystallization of the pure phase La0.8Ca0.2CrO3. This behaviour was also determined for the crystallization of the La0.9Ca0.1CrO3 and La0.8Ca0.2CrO3 powders. The minimum temperature that achieves the perovskite powders formation was found to increase at 450 ºC, when Al3+ partially substituted Cr3+ sites. This situation is more critical when two different metal dopants were simultaneously incorporated at the same time to form a complex oxide such as; LaCa0.2Cr0.95Al0.05O3 and LaCa0.2Cr0.9Al0.1O3, the powders corresponding to these LC solid solutions were obtained at minimum temperature of 475 ºC, which is a significant increase (100 ºC) when compared with the low temperature needed to form the pure LaCrO3. The significant variation on the processing parameters

particle size for the ST powders can be optimized by limiting the dissolution-recrystallization mechanism that operates at sever treatment conditions of temperature (T> 200 ºC) and long periods (t> 72 h). Therefore, these results probe that the hydrothermal technique combined with the use of mineral (pure and contaminated) as precursor for the synthesis of inorganic materials can be an attractive technique to be explored at an industrial scale.

## **2.4.3.4 Hydrothermal crystallization of perovskite lanthanum chromite oxides**

Lanthanum chromite (LaCrO3, LC) powders substituted with alkaline metals (Ca or Sr) have been widely accepted as the candidate for interconnection materials in Solid Oxide Fuel Cells (SOFCs) and also particular solid solutions can be used as oxygen sensor at high temperature (Chakraborty et. al., 2000) The partial substitution of lanthanum ions with alkaline metal ions, Ca or Sr in the A site of La, increases the chemical stability and electric conductivity, whilst Al or Ni in the B site reduces the thermal expansion coefficient at high temperature, when compared with the properties of pure lanthanum chromite (Ianculescu, et. al., 2001). Hitherto, various chemical routes have been used to process lanthanum chromite powders (Bliger. et. al., 1997). However, these chemical processes involve heat treatments at temperatures beyond 700 °C, in order to obtain the crystalline phase pure phase and their solid solutions with Ca or Sr in the A site and Ni or Al in the B in the sites of the perovskite-like structure ABO3. Hence, the present authors recently determined through exhaustive research work, the aspects related to the crystallization of precursor complex gel of the La1-XMXCr1-YNYOH6-x compound into their respective perovskite solid solutions La1-XMXCr1-YNYO3 (where M = Ca or Sr and N = Al or Ni) with orthorhombic crystalline structure under hydrothermal conditions, at a temperature range between 350–500 °C for short reaction intervals (0.5–2 h).

The precursor complex gel was prepared by the alkaline coprecipitation method widely used to prepare this type of materials; the details of the chemical preparative method are given elsewhere (Inagaki et. al., 1990). Precursor lanthanum chromite complex gels were prepared by employing reagent grade chemicals of: LaCl3•7H2O (99.998%), Cr(NO3)3•9H2O (99.9%), CaCl2•2H2O (99%), SrCl2•2H2O (99%), Ni(NO3)2•6H2O, Al(NO3)3•9H2O and NaOH (99.998%) (Wako Pure Chemical Industries, Ltd., Japan). Aqueous solutions with a concentration 0.05 M of LaCl3, Cr(NO3)3 and CaCl2 were prepared with deionized water, and a solution of 0.5 M of NaOH was employed as coprecipitation media. In a typical procedure, a volume of 475 ml of the precipitating solution (NaOH) was poured in a biker, and chromium or the mixtures of Cr+Al, Cr+Ni; solution (500 ml) was then mixed, which results in the formation of an opaque whitish green precipitate (Cr(OH)3, or Al(OH)3, Ni(OH)2), which was subsequently dissolved by vigorous stirring. Finally, the coprecipitation of the complex gel was carried out by the addition of the same volume (500 ml) of the solution containing the other elements, La or the mixture of La+Ca, La+Sr. The solutions were mixed in different volumetric ratios, La:M:Cr:N, 1:0:1:0, 0.9:0.1:1:0 and 0.8:0.2:1:0, 1:0:0.95:0.05; 1:0:0.9:0.1; 0.8:0.2:0.95:0.05, 0.8:0.2:0.9:0.1; which matches the compositional stoichiometric of the solid solutions, LaCrO3, La0.9Ca0.1CrO3, La0.8Ca0.2CrO3, La0.8Sr0.1CrO3, La0.8Sr0.2CrO3, LaCr0.95Al0.05O3, LaCr0.9Al0.1O3, LaCa0.2Cr0.95Al0.05O3, LaCa0.2Cr0.9Al0.1O3, LaSr0.2Cr0.95Al0.05O3, LaSr0.2Cr0.9Al0.1O3, LaSr0.2Cr0.95Ni0.05O3, LaSr0.2Cr0.9Ni0.1O3. The coprecipitated gel was centrifuged and a volume of 20 ml was then poured into a hydrothermal Hastelloy C-lined microautoclave (40 ml capacity, Figure 10). The vessel was heated at a constant rate of 20 ºC/min up at various temperatures (350–500

particle size for the ST powders can be optimized by limiting the dissolution-recrystallization mechanism that operates at sever treatment conditions of temperature (T> 200 ºC) and long periods (t> 72 h). Therefore, these results probe that the hydrothermal technique combined with the use of mineral (pure and contaminated) as precursor for the synthesis of inorganic

Lanthanum chromite (LaCrO3, LC) powders substituted with alkaline metals (Ca or Sr) have been widely accepted as the candidate for interconnection materials in Solid Oxide Fuel Cells (SOFCs) and also particular solid solutions can be used as oxygen sensor at high temperature (Chakraborty et. al., 2000) The partial substitution of lanthanum ions with alkaline metal ions, Ca or Sr in the A site of La, increases the chemical stability and electric conductivity, whilst Al or Ni in the B site reduces the thermal expansion coefficient at high temperature, when compared with the properties of pure lanthanum chromite (Ianculescu, et. al., 2001). Hitherto, various chemical routes have been used to process lanthanum chromite powders (Bliger. et. al., 1997). However, these chemical processes involve heat treatments at temperatures beyond 700 °C, in order to obtain the crystalline phase pure phase and their solid solutions with Ca or Sr in the A site and Ni or Al in the B in the sites of the perovskite-like structure ABO3. Hence, the present authors recently determined through exhaustive research work, the aspects related to the crystallization of precursor complex gel of the La1-XMXCr1-YNYOH6-x compound into their respective perovskite solid solutions La1-XMXCr1-YNYO3 (where M = Ca or Sr and N = Al or Ni) with orthorhombic crystalline structure under hydrothermal conditions, at a temperature range between 350–500 °C for

The precursor complex gel was prepared by the alkaline coprecipitation method widely used to prepare this type of materials; the details of the chemical preparative method are given elsewhere (Inagaki et. al., 1990). Precursor lanthanum chromite complex gels were prepared by employing reagent grade chemicals of: LaCl3•7H2O (99.998%), Cr(NO3)3•9H2O (99.9%), CaCl2•2H2O (99%), SrCl2•2H2O (99%), Ni(NO3)2•6H2O, Al(NO3)3•9H2O and NaOH (99.998%) (Wako Pure Chemical Industries, Ltd., Japan). Aqueous solutions with a concentration 0.05 M of LaCl3, Cr(NO3)3 and CaCl2 were prepared with deionized water, and a solution of 0.5 M of NaOH was employed as coprecipitation media. In a typical procedure, a volume of 475 ml of the precipitating solution (NaOH) was poured in a biker, and chromium or the mixtures of Cr+Al, Cr+Ni; solution (500 ml) was then mixed, which results in the formation of an opaque whitish green precipitate (Cr(OH)3, or Al(OH)3, Ni(OH)2), which was subsequently dissolved by vigorous stirring. Finally, the coprecipitation of the complex gel was carried out by the addition of the same volume (500 ml) of the solution containing the other elements, La or the mixture of La+Ca, La+Sr. The solutions were mixed in different volumetric ratios, La:M:Cr:N, 1:0:1:0, 0.9:0.1:1:0 and 0.8:0.2:1:0, 1:0:0.95:0.05; 1:0:0.9:0.1; 0.8:0.2:0.95:0.05, 0.8:0.2:0.9:0.1; which matches the compositional stoichiometric of the solid solutions, LaCrO3, La0.9Ca0.1CrO3, La0.8Ca0.2CrO3, La0.8Sr0.1CrO3, La0.8Sr0.2CrO3, LaCr0.95Al0.05O3, LaCr0.9Al0.1O3, LaCa0.2Cr0.95Al0.05O3, LaCa0.2Cr0.9Al0.1O3, LaSr0.2Cr0.95Al0.05O3, LaSr0.2Cr0.9Al0.1O3, LaSr0.2Cr0.95Ni0.05O3, LaSr0.2Cr0.9Ni0.1O3. The coprecipitated gel was centrifuged and a volume of 20 ml was then poured into a hydrothermal Hastelloy C-lined microautoclave (40 ml capacity, Figure 10). The vessel was heated at a constant rate of 20 ºC/min up at various temperatures (350–500

materials can be an attractive technique to be explored at an industrial scale.

short reaction intervals (0.5–2 h).

**2.4.3.4 Hydrothermal crystallization of perovskite lanthanum chromite oxides**

ºC) for a reaction interval between 0.5 and 2 h. After the treatments, the precipitates were well washed with distilled water, decanted and then dried in an oven at 100 ºC overnight (Rivas-Vázquez et. al., 2004, 2006; Rendón-Angeles et. al., 2009).

Fig. 10. Typical Hastelloy C type lined microautoclave and scheme of the cross section of the vessel, which was employed for conducting the synthesis of perovskite lanthanum chromite powders under hydrothermal conditions.

The minimum temperature for conducting the crystallization of the pure LC powders was 375 ºC for a reaction interval of 1 h. This parameter, however, is mainly affected by the incorporation of the metal ions in both A and B sites of the perovskite structure ABO3, and also with the amount of metal dopant ion inserted. Thus, the crystallization of powders corresponding to the solid solutions La1-XMXCr1-YNYO3 (where M = Ca or Sr and N = Al or Ni) was observed to proceeds at temperatures above 400 ºC, when only a 10 mol% of Ca2+ or Sr2+ were partially substituting La3+ site. Thus, in the particular case of the La0.9Ca0.1CrO3 and La0.8Ca0.2CrO3 solid solutions, the structural analyses in Figure 11a depicted that at a constant reaction interval 1 h, the powders of La0.9Ca0.1CrO3 were produced at a temperature of 400 ºC without contaminant formation, in contrast with the La0.8Ca0.2CrO3 powders that were produced together with a marked amount of secondary crystalline phases of La(OH)3, CrOOH, and CaCrO4, these phases were eliminated by increasing the reaction temperature up to 425 ºC resulting in the crystallization of the pure phase La0.8Ca0.2CrO3. This behaviour was also determined for the crystallization of the La0.9Ca0.1CrO3 and La0.8Ca0.2CrO3 powders. The minimum temperature that achieves the perovskite powders formation was found to increase at 450 ºC, when Al3+ partially substituted Cr3+ sites. This situation is more critical when two different metal dopants were simultaneously incorporated at the same time to form a complex oxide such as; LaCa0.2Cr0.95Al0.05O3 and LaCa0.2Cr0.9Al0.1O3, the powders corresponding to these LC solid solutions were obtained at minimum temperature of 475 ºC, which is a significant increase (100 ºC) when compared with the low temperature needed to form the pure LaCrO3. The significant variation on the processing parameters

Preparation of Selected Ceramic Compounds by

Rendón-Angeles et. al., 2009).

La0.8M0.2CrO3.

Controlled Crystallization Under Hydrothermal Conditions 229

chemical stability of the crystalline phase, which must be low and therefore the dissolution proceeds rapidly. Indeed, this inference is supported by the fact that the amount of regular La0.8Sr0.2CrO3 particles is further decreased in this particular powder. Although, the particles exhibited a non-peculiar morphology when compared with those used for preparing dense advanced ceramic materials, it was experimentally probed that these fine powders have better sinterability features even in oxidizing atmospheres (Rivas-Vázquez et. al., 2004, 2006;

 Fig. 12. Transmission electron micrographs of hydrothermally produced powders for 1 h at

The details of the reaction pathway that are linked to the crystallization of the final tailored LC composition were determined during the preparation of La0.9Ca0.1CrO3 solid solution. Preliminary observations conducted even during the heating stage of treatment confirmed that the hydrothermal crystallization involves a preliminary reaction stage, which is related with the dehydration process of the complex gel and proceeds at temperatures above 300 ºC, resulting on the autogeneous formation of hydrothermal solvent inside the reaction vessel. The pH of the remaining solutions after the hydrothermal treatment varied in the range between 7.2-7.8, which confirms that the crystallization of the lanthanum chromite powders

400 ºC (a) La0.9Ca0.1CrO3 and (c) La0.9Sr0.1CrO3; and 425 ºC (b) La0.8M0.2CrO3 and (d)

experimentally determined, namely the temperature, can be associated with the thermodynamic fundamental principles, e.g. the Gibbs free energy. The calculus of the G values for the formation of Ca and Sr doped LC solid solutions over a wide range of temperature (Fig. 11b), are in a good agreement with the trend of the experimental results aforementioned. Because, in terms of energy, the crystallization reaction that required less energy consumption is that leads to the formation of pure LaCrO3, compound that exhibit the lowest G values in comparison with those of La0.9M0.1CrO3 and La0.8M0.2CrO3 solid solutions (M= Ca2+ or Sr2+).

Fig. 11. (a) X-ray diffraction patterns of (*i*) La0.9Ca0.1CrO3 and (*ii*) La0.8Ca0.2CrO3 powders produced at 400 ºC for 1 h. LaCrO3 (JCPDS 33-701) compound doted line; () La(OH)3, () CrOOH, () CaCrO4, () low intensity peaks of LaCrO3. (b) Variation of the G formation for pure LaCrO3 and La0.9M0.1CrO3 and La0.8M0.2CrO3 solid solutions (M= Ca2+ or Sr2+).

Morphological aspects of powders corresponding to the solid solutions of La0.9M0.1CrO3 and La0.8M0.2CrO3 (M= Ca2+ or Sr2+) solid solutions obtained at 400 and 425 ºC, respectively; are shown in Figure 12. Particles with submicron size and irregular morphology, which resembles a peanut-like shape, were preferentially formed on all cases investigated, it is also indicated that the particle morphology of these compounds is irrespective of the gel crystallization temperature. Furthermore, the particles showed a marked agglomeration due to its particular morphology and particle size aspect. One point that deserves emphasis is that related with the marked tendency for particle bounding that underwent the particles. These are reliable evidences that indicate the particles of the formed LC solid solutions were partially dissolved in the fluid at an intermediate reaction stage (>1 h) of the hydrothermal treatment, once the secondary crystalline phases were completely exhausted.

This phenomenon seems to occur markedly on the La0.8Sr0.2CrO3 particles (Fig. 12d) hydrothermally synthesized 1 h at 425 ºC, because these particles exhibit the smallest particle size (average 250 nm) in comparison with the other three powders that have an average particle size of 350 nm (Figs. 12a-12c). This fact can be explained based on the

experimentally determined, namely the temperature, can be associated with the thermodynamic fundamental principles, e.g. the Gibbs free energy. The calculus of the G values for the formation of Ca and Sr doped LC solid solutions over a wide range of temperature (Fig. 11b), are in a good agreement with the trend of the experimental results aforementioned. Because, in terms of energy, the crystallization reaction that required less energy consumption is that leads to the formation of pure LaCrO3, compound that exhibit the lowest G values in comparison with those of La0.9M0.1CrO3 and La0.8M0.2CrO3 solid

 Fig. 11. (a) X-ray diffraction patterns of (*i*) La0.9Ca0.1CrO3 and (*ii*) La0.8Ca0.2CrO3 powders produced at 400 ºC for 1 h. LaCrO3 (JCPDS 33-701) compound doted line; () La(OH)3, () CrOOH, () CaCrO4, () low intensity peaks of LaCrO3. (b) Variation of the G formation for pure LaCrO3 and La0.9M0.1CrO3 and La0.8M0.2CrO3 solid solutions (M= Ca2+ or Sr2+).

Morphological aspects of powders corresponding to the solid solutions of La0.9M0.1CrO3 and La0.8M0.2CrO3 (M= Ca2+ or Sr2+) solid solutions obtained at 400 and 425 ºC, respectively; are shown in Figure 12. Particles with submicron size and irregular morphology, which resembles a peanut-like shape, were preferentially formed on all cases investigated, it is also indicated that the particle morphology of these compounds is irrespective of the gel crystallization temperature. Furthermore, the particles showed a marked agglomeration due to its particular morphology and particle size aspect. One point that deserves emphasis is that related with the marked tendency for particle bounding that underwent the particles. These are reliable evidences that indicate the particles of the formed LC solid solutions were partially dissolved in the fluid at an intermediate reaction stage (>1 h) of the hydrothermal

This phenomenon seems to occur markedly on the La0.8Sr0.2CrO3 particles (Fig. 12d) hydrothermally synthesized 1 h at 425 ºC, because these particles exhibit the smallest particle size (average 250 nm) in comparison with the other three powders that have an average particle size of 350 nm (Figs. 12a-12c). This fact can be explained based on the

treatment, once the secondary crystalline phases were completely exhausted.

solutions (M= Ca2+ or Sr2+).

chemical stability of the crystalline phase, which must be low and therefore the dissolution proceeds rapidly. Indeed, this inference is supported by the fact that the amount of regular La0.8Sr0.2CrO3 particles is further decreased in this particular powder. Although, the particles exhibited a non-peculiar morphology when compared with those used for preparing dense advanced ceramic materials, it was experimentally probed that these fine powders have better sinterability features even in oxidizing atmospheres (Rivas-Vázquez et. al., 2004, 2006; Rendón-Angeles et. al., 2009).

Fig. 12. Transmission electron micrographs of hydrothermally produced powders for 1 h at 400 ºC (a) La0.9Ca0.1CrO3 and (c) La0.9Sr0.1CrO3; and 425 ºC (b) La0.8M0.2CrO3 and (d) La0.8M0.2CrO3.

The details of the reaction pathway that are linked to the crystallization of the final tailored LC composition were determined during the preparation of La0.9Ca0.1CrO3 solid solution. Preliminary observations conducted even during the heating stage of treatment confirmed that the hydrothermal crystallization involves a preliminary reaction stage, which is related with the dehydration process of the complex gel and proceeds at temperatures above 300 ºC, resulting on the autogeneous formation of hydrothermal solvent inside the reaction vessel. The pH of the remaining solutions after the hydrothermal treatment varied in the range between 7.2-7.8, which confirms that the crystallization of the lanthanum chromite powders

Preparation of Selected Ceramic Compounds by

Figure 14 for La0.8Sr0.2CrO3 powders.

Controlled Crystallization Under Hydrothermal Conditions 231

Regarding the peculiar morphology of the La0.9M0.1CrO3 and La0.8M0.2CrO3 solid solutions (M= Ca2+ or Sr2+). Two main factors promoted the formation of the peanut-like shaped particles (Figure 14); the first is related with the nucleation and growth process which proceeded very fast for reaction intervals of 0.5 up to 2 h at interval of temperatures (400– 450 ◦C) and limited the growth of La1−XMXCrO3 particles. The second one is associated with the concentration of the solvent media; the alkaline fluid autogenously formed during the hydrothermal treatment is not capable of achieving the actuated conditions that lead the crystallization of particles with cubic habit (Zheng et. al., 1999; Spooren et. al., 2003). However, the La1−XSrXCrO3 particles were partially dissolved in the solvent media at high temperature (400–450 ºC), resulting in a marked particle joining by developing necks on the surfaces of primary synthesized particles in contact, as shown in

 Fig. 14. Transmission electron micrographs of La0.8Sr0.2CrO3 powders obtained at 450 ºC for reaction intervals of (a) 1 and (b) 2 h. La0.8Sr0.2Cr0.9Al0.9O3 powders prepared at 475 ºC for 1 h

The local recrystallization of the solute is mainly promoted inside the reaction vessel, due to the fact that no fluid convection occurred during the hydrothermal treatment, because the treatments were conducted under static conditions at constant temperature, therefore, mass

in NaOH solutions of (c) 0.1 and (d) 5 M using as a precursor dried complex gel.

was conducted in low alkaline conditions. Once the treatment temperature is reached (400 ºC), as a result of the complex gel dehydration process, the crystallization of intermediate secondary phases proceeded during the earlier stages of the reaction (10–30 min). The major stable secondary phases that were produced during this reaction interval range were La(OH)3, CrOOH, which correspond to the needle and platelets, respectively (Figs. 13a– 13b); these observations are in a good consistence with the X-ray diffraction patterns of the powders produced at 400 ºC for different reaction intervals (Fig. 13d), these results also showed the presence of a slight amount of CaCrO4. The dissolution of the reaction byproducts is fast in such a diluted alkaline hydrothermal media. Hence, the nucleation of the oxide particles proceeds preferentially at the surface of the remaining complex gel during the dehydration process, because the crystallization by the process employed in this study proceeded at higher temperatures (300–450 ºC) than those determined for LaCrO3 and La0.5Sr0.5MnO3 particles under concentrated alkaline hydrothermal conditions (8 M KOH), 240 and 260 ºC, respectively (Zheng et. al., 1999; Spooren et. al., 2003). The complete crystallization of La0.9Ca0.1CrO3 particles was even promoted for a shorter interval as 1 h, resulting in the formation of irregular particles resembling peanut-like shapes (Fig. 13c).

Fig. 13. TEM micrographs of the reaction products obtained under hydrothermal conditions at 400 ºC for different intervals (a) 0.33, (b) 0.5 and (c) 1 h. (d) X-ray diffraction patterns of complex gel corresponding to the above micrographs. LaCrO3 (JCPDS 33-701) compound doted line; () La(OH)3, () CrOOH, () CaCrO4, () low intensity peaks of LaCrO3.

was conducted in low alkaline conditions. Once the treatment temperature is reached (400 ºC), as a result of the complex gel dehydration process, the crystallization of intermediate secondary phases proceeded during the earlier stages of the reaction (10–30 min). The major stable secondary phases that were produced during this reaction interval range were La(OH)3, CrOOH, which correspond to the needle and platelets, respectively (Figs. 13a– 13b); these observations are in a good consistence with the X-ray diffraction patterns of the powders produced at 400 ºC for different reaction intervals (Fig. 13d), these results also showed the presence of a slight amount of CaCrO4. The dissolution of the reaction byproducts is fast in such a diluted alkaline hydrothermal media. Hence, the nucleation of the oxide particles proceeds preferentially at the surface of the remaining complex gel during the dehydration process, because the crystallization by the process employed in this study proceeded at higher temperatures (300–450 ºC) than those determined for LaCrO3 and La0.5Sr0.5MnO3 particles under concentrated alkaline hydrothermal conditions (8 M KOH), 240 and 260 ºC, respectively (Zheng et. al., 1999; Spooren et. al., 2003). The complete crystallization of La0.9Ca0.1CrO3 particles was even promoted for a shorter interval as 1 h, resulting in the formation of irregular particles resembling peanut-like shapes (Fig. 13c).

 Fig. 13. TEM micrographs of the reaction products obtained under hydrothermal conditions at 400 ºC for different intervals (a) 0.33, (b) 0.5 and (c) 1 h. (d) X-ray diffraction patterns of complex gel corresponding to the above micrographs. LaCrO3 (JCPDS 33-701) compound doted line; () La(OH)3, () CrOOH, () CaCrO4, () low intensity peaks of LaCrO3.

Regarding the peculiar morphology of the La0.9M0.1CrO3 and La0.8M0.2CrO3 solid solutions (M= Ca2+ or Sr2+). Two main factors promoted the formation of the peanut-like shaped particles (Figure 14); the first is related with the nucleation and growth process which proceeded very fast for reaction intervals of 0.5 up to 2 h at interval of temperatures (400– 450 ◦C) and limited the growth of La1−XMXCrO3 particles. The second one is associated with the concentration of the solvent media; the alkaline fluid autogenously formed during the hydrothermal treatment is not capable of achieving the actuated conditions that lead the crystallization of particles with cubic habit (Zheng et. al., 1999; Spooren et. al., 2003). However, the La1−XSrXCrO3 particles were partially dissolved in the solvent media at high temperature (400–450 ºC), resulting in a marked particle joining by developing necks on the surfaces of primary synthesized particles in contact, as shown in Figure 14 for La0.8Sr0.2CrO3 powders.

Fig. 14. Transmission electron micrographs of La0.8Sr0.2CrO3 powders obtained at 450 ºC for reaction intervals of (a) 1 and (b) 2 h. La0.8Sr0.2Cr0.9Al0.9O3 powders prepared at 475 ºC for 1 h in NaOH solutions of (c) 0.1 and (d) 5 M using as a precursor dried complex gel.

The local recrystallization of the solute is mainly promoted inside the reaction vessel, due to the fact that no fluid convection occurred during the hydrothermal treatment, because the treatments were conducted under static conditions at constant temperature, therefore, mass

Preparation of Selected Ceramic Compounds by

(Yoshino et. al., 1985).

Controlled Crystallization Under Hydrothermal Conditions 233

This particular transformation involves the conversion of MSO4 into MCO3, M= Sr or Ba, because these carbonated compound are widely used as a precursors for the preparation of strontium and barium inorganic compounds (Suaréz-Orduña et. al., 2004a; Rendón-Angeles et. al., 2008). Preliminary evidences of the ion exchange replacement process were investigated on the mineral celestite specie, thermodynamic and kinetic details were reported elsewhere (Yoshino et. al., 1985). The exchange of SO42- ions with CO32- ions was investigated on large mineral SrSO4 bulky crystal plates, which were leached at low temperature (55 ºC). The ion replacement was achieved by two reaction mechanisms, at initial and intermediate stages, the superficial reaction and the diffusion of SO42- ions produced a dense SrCO3, the complete ion replacement process proceeded by a second mechanism by a solid-state ion exchange mechanism even under hydrothermal conditions

Recently, a better approach conducted to elucidate the source of the conversion M2+SO4 to M2+CO3, M= Sr or Ba; was proposed based on the crystalline structural differences associated to the physical bulk molar volume change, which also must occur due to the replacement of a large anion SO42- by a smaller one CO32-. Thus, the study was conducted in large single crystals of celestite SrSO4 and barite BaCO3, which were treated under alkaline hydrothermal conditions using high concentrated carbonated solutions of Na2CO3 and K2CO3. The mineral celestite single crystals (SrSO4, square plates 10 mm wide and 3 mm thick) were topotaxially converted to strontianite (SrCO3) under alkaline hydrothermal conditions. The reaction was completed in a short reaction time (such as 24 h) at a temperature of 250 C. Increasing the treatment temperature and the molar ratio CO32-/SO4

accelerated the exchange of SO42- with CO32- ions. Under these conditions the topotaxial hydrothermal conversion to strontianite is carried out with the formation of intermediate CO32--rich solid solutions in the system SrCO3–SrSO4, this is shown in the Figure 15; and these solid solutions were formed by a controlled crystallization process achieved by the cluster dissolution–recrystallization. An anisotropic dissolution gave a characteristic texture in the converted crystals, and the difference on the reactivity of the celestite crystals in Na2CO3 and K2CO3 solutions resulted in a different texture inside the pseudomorphical

Despite the aspect of the converted crystals remained without any change after the conversion process (Fig. 15c), the formation of small holes randomly distributed were produced on outside layer produced as a result of the transformed into the MCO3 phase Figure 16. However, the morphology of these holes underwent a change due to a simultaneous dissolution of the solid product, the recrystallization process of this phase also occurred at the same place where the previous dissolution occurred. One point that must be emphasized is that the two phases are separated by a sharp boundary in texture, as well as composition. From this observation it is clear that the replacement reaction begins from the surface of the mineral that was in contact with the hydrothermal media (Fig. 16a). In addition, the reaction proceeded by the incorporation of the solvent through the inner porosity. The holes form a zig-zag network inside the crystals which allows incorporating fresh solvent media at the reaction front (Suaréz-Orduña et. al., 2004a; Rendón-Angeles et. al., 2008). This reaction is similar to the ionic replacement process that was found for the conversion of chlorapatite and hydroxyapatite single crystals into

converted strontianite crystals (Rendón-Angeles et. al., 2000b).

2-

transfer due to convection was further limited (Bayrappa & Yoshimura, 2001). Hence, the crystallization mechanism is similar to that described previously for other perovskite species, but it differs on the rate of kinetics that lead to produce particles particular aspect. In contrast, the particle bonding was limited by modifying the treatment conditions, using a dried complex gel La0.8Sr0.2Cr0.9Al0.1O3 and different aqueous solvent media (water, KOH, NaOH and KF), resulting in the optimum control of dissolution and crystallization of fine particles with regular. A marked particle growth, however, was determined as a result of increasing the concentration of the alkaline solution from 0.1 to 5 M (NaOH or KOH). The formation of submicron particles monodispersed (0.5 – 0.75 μm) was found to proceed on these solvents (Figs. 14c, 14d). Furthermore, the employment of the alkaline solvents NaOH and KF leads to the crystallization of pseudo-cubic shaped and hexagonal plate particles, respectively; the differences on the particle morphology control are due to the chemical reactivity of the different solutions with the dried complex gel.

#### **2.5 Replacement reactions on minerals species under hydrothermal conditions as a new approach for preparing inorganic materials**

Another different approach similar to the mineral transformation into perovskite oxide powders was investigated to establish an alternative processing route for mineral ores via controlled dissolution-precipitation. The precursor mineral considered were the alkaline earth sulphates of SrSO4 and BaSO4, for preparing high pure strontium inorganic compounds, namely SrCO3, BaCO3, SrCrO4, SrF2, Sr(OH)2. This chemical preparative method involves the dissolution–recrystallization mechanism, similar to the process that achieves the ionic replacement reaction in mineral ores, and promotes the conversion of natural ores at the earth's crust into more chemically stable mineral species or inorganic compounds (Putnis, 2002, 2009). The development of more efficient and environmental friendly chemical routes has recently been under concern of some researchers. The new chemical techniques or the optimized conventional routes might lead to reduce the pollution grade of contaminants, which is produced during mineral processing stages. The related alkaline earth sulphate minerals, such as SrSO4 and BaSO4, have been exploited since several decades, because these are the main source of the alkaline earth metals elements, Sr and Ba, and are used as a main source for preparing inorganic compounds of Sr and Ba. During the last decade, some attention has been paid to the use of celestite ore for producing functional ceramic compounds with magnetic properties, e.g. strontium hexaferrite (SrFe12O19); two different methods were proposed, coprecipitation in aqueous solutions via mineral powder leaching and powder mechanochemical activation techniques (Hessien et. al., 2009; Tiwary, et. al. 2008). Hence, since a decade, the present authors have devoted efforts in order to investigate, from a different approach of preparative chemistry; the nature of the possible replacement reactions that can be achieved under hydrothermal conditions, as was found to proceed even in synthetic inorganic compounds (Rendón-Angeles et. al., 2000b). This could derive in an optimum method for preparing high grade Sr and Ba compounds. The most relevant aspects found regarding this topic are given in this section.

#### **2.5.1 Aspects of the compositional and structural transformation on sulphate minerals under alkaline hydrothermal conditions**

Hitherto, the ion exchange reaction of SO42- ions with CO3 2- ions has been one of the subjects of research work, which involves both mineral celestite (SrSO4) and barite (BaSO4) species.

transfer due to convection was further limited (Bayrappa & Yoshimura, 2001). Hence, the crystallization mechanism is similar to that described previously for other perovskite species, but it differs on the rate of kinetics that lead to produce particles particular aspect. In contrast, the particle bonding was limited by modifying the treatment conditions, using a dried complex gel La0.8Sr0.2Cr0.9Al0.1O3 and different aqueous solvent media (water, KOH, NaOH and KF), resulting in the optimum control of dissolution and crystallization of fine particles with regular. A marked particle growth, however, was determined as a result of increasing the concentration of the alkaline solution from 0.1 to 5 M (NaOH or KOH). The formation of submicron particles monodispersed (0.5 – 0.75 μm) was found to proceed on these solvents (Figs. 14c, 14d). Furthermore, the employment of the alkaline solvents NaOH and KF leads to the crystallization of pseudo-cubic shaped and hexagonal plate particles, respectively; the differences on the particle morphology control are due to the chemical

**2.5 Replacement reactions on minerals species under hydrothermal conditions as a** 

Another different approach similar to the mineral transformation into perovskite oxide powders was investigated to establish an alternative processing route for mineral ores via controlled dissolution-precipitation. The precursor mineral considered were the alkaline earth sulphates of SrSO4 and BaSO4, for preparing high pure strontium inorganic compounds, namely SrCO3, BaCO3, SrCrO4, SrF2, Sr(OH)2. This chemical preparative method involves the dissolution–recrystallization mechanism, similar to the process that achieves the ionic replacement reaction in mineral ores, and promotes the conversion of natural ores at the earth's crust into more chemically stable mineral species or inorganic compounds (Putnis, 2002, 2009). The development of more efficient and environmental friendly chemical routes has recently been under concern of some researchers. The new chemical techniques or the optimized conventional routes might lead to reduce the pollution grade of contaminants, which is produced during mineral processing stages. The related alkaline earth sulphate minerals, such as SrSO4 and BaSO4, have been exploited since several decades, because these are the main source of the alkaline earth metals elements, Sr and Ba, and are used as a main source for preparing inorganic compounds of Sr and Ba. During the last decade, some attention has been paid to the use of celestite ore for producing functional ceramic compounds with magnetic properties, e.g. strontium hexaferrite (SrFe12O19); two different methods were proposed, coprecipitation in aqueous solutions via mineral powder leaching and powder mechanochemical activation techniques (Hessien et. al., 2009; Tiwary, et. al. 2008). Hence, since a decade, the present authors have devoted efforts in order to investigate, from a different approach of preparative chemistry; the nature of the possible replacement reactions that can be achieved under hydrothermal conditions, as was found to proceed even in synthetic inorganic compounds (Rendón-Angeles et. al., 2000b). This could derive in an optimum method for preparing high grade Sr and Ba compounds. The most

reactivity of the different solutions with the dried complex gel.

relevant aspects found regarding this topic are given in this section.

**minerals under alkaline hydrothermal conditions** 

**2.5.1 Aspects of the compositional and structural transformation on sulphate** 

Hitherto, the ion exchange reaction of SO42- ions with CO32- ions has been one of the subjects of research work, which involves both mineral celestite (SrSO4) and barite (BaSO4) species.

**new approach for preparing inorganic materials** 

This particular transformation involves the conversion of MSO4 into MCO3, M= Sr or Ba, because these carbonated compound are widely used as a precursors for the preparation of strontium and barium inorganic compounds (Suaréz-Orduña et. al., 2004a; Rendón-Angeles et. al., 2008). Preliminary evidences of the ion exchange replacement process were investigated on the mineral celestite specie, thermodynamic and kinetic details were reported elsewhere (Yoshino et. al., 1985). The exchange of SO4 2- ions with CO32- ions was investigated on large mineral SrSO4 bulky crystal plates, which were leached at low temperature (55 ºC). The ion replacement was achieved by two reaction mechanisms, at initial and intermediate stages, the superficial reaction and the diffusion of SO42- ions produced a dense SrCO3, the complete ion replacement process proceeded by a second mechanism by a solid-state ion exchange mechanism even under hydrothermal conditions (Yoshino et. al., 1985).

Recently, a better approach conducted to elucidate the source of the conversion M2+SO4 to M2+CO3, M= Sr or Ba; was proposed based on the crystalline structural differences associated to the physical bulk molar volume change, which also must occur due to the replacement of a large anion SO42- by a smaller one CO32-. Thus, the study was conducted in large single crystals of celestite SrSO4 and barite BaCO3, which were treated under alkaline hydrothermal conditions using high concentrated carbonated solutions of Na2CO3 and K2CO3. The mineral celestite single crystals (SrSO4, square plates 10 mm wide and 3 mm thick) were topotaxially converted to strontianite (SrCO3) under alkaline hydrothermal conditions. The reaction was completed in a short reaction time (such as 24 h) at a temperature of 250 C. Increasing the treatment temperature and the molar ratio CO32-/SO4 2 accelerated the exchange of SO42- with CO32- ions. Under these conditions the topotaxial hydrothermal conversion to strontianite is carried out with the formation of intermediate CO32--rich solid solutions in the system SrCO3–SrSO4, this is shown in the Figure 15; and these solid solutions were formed by a controlled crystallization process achieved by the cluster dissolution–recrystallization. An anisotropic dissolution gave a characteristic texture in the converted crystals, and the difference on the reactivity of the celestite crystals in Na2CO3 and K2CO3 solutions resulted in a different texture inside the pseudomorphical converted strontianite crystals (Rendón-Angeles et. al., 2000b).

Despite the aspect of the converted crystals remained without any change after the conversion process (Fig. 15c), the formation of small holes randomly distributed were produced on outside layer produced as a result of the transformed into the MCO3 phase Figure 16. However, the morphology of these holes underwent a change due to a simultaneous dissolution of the solid product, the recrystallization process of this phase also occurred at the same place where the previous dissolution occurred. One point that must be emphasized is that the two phases are separated by a sharp boundary in texture, as well as composition. From this observation it is clear that the replacement reaction begins from the surface of the mineral that was in contact with the hydrothermal media (Fig. 16a). In addition, the reaction proceeded by the incorporation of the solvent through the inner porosity. The holes form a zig-zag network inside the crystals which allows incorporating fresh solvent media at the reaction front (Suaréz-Orduña et. al., 2004a; Rendón-Angeles et. al., 2008). This reaction is similar to the ionic replacement process that was found for the conversion of chlorapatite and hydroxyapatite single crystals into

Preparation of Selected Ceramic Compounds by

Controlled Crystallization Under Hydrothermal Conditions 235

hydrothermal treatment at 250 °C in a Na2CO3 solution with a molar ratio CO32-/SO42- of 10

In terms of the porosity resulted as consequence of the crystalline transformation, this factor has also been the subject of controversy, because the causes that promote its formation are not clarified yet. However, we have recently found that the formation of the residual porosity depends strongly of two principal factors: i) the differences on the molar volume associated with the crystalline structural differences and ii) the chemical stability (solubility) of the new converted crystalline phase in the hydrothermal media. These inferences were established on partially and completely converted SrCO3 specimens obtained by hydrothermal treatments at 250 ºC for a interval of 24 h with a molar ratio CO32-/SO42- = 10. The inner volume on the converted crystals was determined by helium picnometry measurements and those results are giving in Figure 17. It is clear that the volume measurements conducted on the completed converted SrCO3 plates at 250 °C for 96 h, revealed that the residual inner porosity value obtained on the specimens treated in Na2CO3 solutions is nearly similar to the theoretical value (dotted line in Fig. 17), this value (15.62 %) associated with the reduction of the molar volume, was calculated by considering the unit cell volume values of the parent (celestite, 307.06 Å3) and the product (strontianite, 259.07 Å3). This crystalline structural variation is related with the formation of the residual inner porosity, because macroscopically the crystal plate remains without any change regarding its shape and dimension, therefore, the bulk molar volume reduction does not proceed on the crystal plate and this must be compensated by which is like to proceed the residual porosity. Moreover, the control of the porosity depends on the chemical stability of the phase crystallized with the ion exchange media, once the replacement reaction was completed (Suaréz-Orduña et. al., 2004b). This inference is supported by the fact that when the conversion was carried out in K2CO3 solutions, the residual inner porosity on the completed converted SrCO3 plate was markedly increased due to a further dissolution of the converted SrCO3 crystal, in comparison with the porosity obtained the SrCO3 plates

**2.5.2 Aspects associated with the formation and control of the residual porosity** 

Fig. 16. SEM micrographs of barite crystals converted to barium carbonate after

for different reaction intervals (a) 24 h (b) 192 h (Rendón-Angeles et. al., 2008).

**during the replacement reaction of mineral sulphate species** 

transformed with highly concentrated Na2CO3 solutions at CO3

2-/SO4

2- > 5.

fluorapatite single crystals under hydrothermal conditions (Rendón-Angeles et. al., 2000b). In terms of the macroscopic aspects and crystalline structural differences associated with the replacement of a large ion (SO42- = 4.32 Å) by a smaller (CO32- = 1.55 Å). However, the completely converted MCO3 crystals are constituted for very tiny crystals randomly oriented which resemble a polycrystalline arrangement on the converted M2+CO3. Thus in this particular case, the conversion process proceeds with the formation of a converted layer that has a peculiar texture (holes, Fig. 16b) and a moving reaction front. Hence, the conversion of the M2+SO4 mineral species into their related carbonated inorganic compounds; is associated to a pseudomorphic replacement process rather than the ion-exchange process. In addition, at the hydrothermal conditions in which the mineral conversion can be achieved, the pseudomorphic replacement process is mainly achieved by a mechanism of coupled bulk dissolution and precipitation. It is well known that this mechanism promotes the formation of a great wide type of inorganic compounds (Putnis, 2002).

Fig. 15. (a) SEM image of the reaction interface and (b) sulfur concentration in a partially converted celestite crystal obtained by hydrothermal treatment in a Na2CO3 solution with a molar ratio CO32-/SO42- = 10, at 250 C for 1 h. and (c) 96 h, aspects of the original SrSO4 (left) and converted SrCO3 (right). Grid size= 10 mm

fluorapatite single crystals under hydrothermal conditions (Rendón-Angeles et. al., 2000b). In terms of the macroscopic aspects and crystalline structural differences associated with the replacement of a large ion (SO42- = 4.32 Å) by a smaller (CO32- = 1.55 Å). However, the completely converted MCO3 crystals are constituted for very tiny crystals randomly oriented which resemble a polycrystalline arrangement on the converted M2+CO3. Thus in this particular case, the conversion process proceeds with the formation of a converted layer that has a peculiar texture (holes, Fig. 16b) and a moving reaction front. Hence, the conversion of the M2+SO4 mineral species into their related carbonated inorganic compounds; is associated to a pseudomorphic replacement process rather than the ion-exchange process. In addition, at the hydrothermal conditions in which the mineral conversion can be achieved, the pseudomorphic replacement process is mainly achieved by a mechanism of coupled bulk dissolution and precipitation. It is well known that this mechanism promotes the formation of a great wide type of inorganic

 Fig. 15. (a) SEM image of the reaction interface and (b) sulfur concentration in a partially converted celestite crystal obtained by hydrothermal treatment in a Na2CO3 solution with a

2-/SO42- = 10, at 250 C for 1 h. and (c) 96 h, aspects of the original SrSO4 (left)

compounds (Putnis, 2002).

molar ratio CO3

and converted SrCO3 (right). Grid size= 10 mm

Fig. 16. SEM micrographs of barite crystals converted to barium carbonate after hydrothermal treatment at 250 °C in a Na2CO3 solution with a molar ratio CO32-/SO42- of 10 for different reaction intervals (a) 24 h (b) 192 h (Rendón-Angeles et. al., 2008).

#### **2.5.2 Aspects associated with the formation and control of the residual porosity during the replacement reaction of mineral sulphate species**

In terms of the porosity resulted as consequence of the crystalline transformation, this factor has also been the subject of controversy, because the causes that promote its formation are not clarified yet. However, we have recently found that the formation of the residual porosity depends strongly of two principal factors: i) the differences on the molar volume associated with the crystalline structural differences and ii) the chemical stability (solubility) of the new converted crystalline phase in the hydrothermal media. These inferences were established on partially and completely converted SrCO3 specimens obtained by hydrothermal treatments at 250 ºC for a interval of 24 h with a molar ratio CO32-/SO42- = 10. The inner volume on the converted crystals was determined by helium picnometry measurements and those results are giving in Figure 17. It is clear that the volume measurements conducted on the completed converted SrCO3 plates at 250 °C for 96 h, revealed that the residual inner porosity value obtained on the specimens treated in Na2CO3 solutions is nearly similar to the theoretical value (dotted line in Fig. 17), this value (15.62 %) associated with the reduction of the molar volume, was calculated by considering the unit cell volume values of the parent (celestite, 307.06 Å3) and the product (strontianite, 259.07 Å3). This crystalline structural variation is related with the formation of the residual inner porosity, because macroscopically the crystal plate remains without any change regarding its shape and dimension, therefore, the bulk molar volume reduction does not proceed on the crystal plate and this must be compensated by which is like to proceed the residual porosity. Moreover, the control of the porosity depends on the chemical stability of the phase crystallized with the ion exchange media, once the replacement reaction was completed (Suaréz-Orduña et. al., 2004b). This inference is supported by the fact that when the conversion was carried out in K2CO3 solutions, the residual inner porosity on the completed converted SrCO3 plate was markedly increased due to a further dissolution of the converted SrCO3 crystal, in comparison with the porosity obtained the SrCO3 plates transformed with highly concentrated Na2CO3 solutions at CO3 2-/SO4 2- > 5.

Preparation of Selected Ceramic Compounds by

molar ratio CrO42-/SO42-=1.

and their relative solubility.

**3. Conclusions** 

Controlled Crystallization Under Hydrothermal Conditions 237

Fig. 18. Partially reacted SrSO4 crystal at 200 °C for 96 h in a hydrothermal media with a

The hydrothermal crystallization of advanced materials is an important branch of chemical science and technology, this process has advantages over conventional technologies, namely the purity of products, quality, and performance, and can be consider as a green chemistry process because is environmentally friendly, because reactions consume lesser energy and these can be carried out under controlled parameters in a closed system. Although, the study of a wide number of oxides have provided physico-chemical experimental information related to solid-aqueous interactions, that affect the dissolution-precipitation mechanism main driving force for achieving bulk solid crystallization, coupled with thermodynamic modelling analysis; have contributed to establish the optimum hydrothermal processing conditions that favours the crystallization proceeds at high yield rates. However, still crystallization research work to carry out on those solid-aqueous systems associated to a specific compound that departs from the thermodynamic modelling limits. In addition, the employment of mineral ore as reactant precursors emerges as an interesting route to explore in order to produce other advanced ceramic compounds via the hydrothermal crystallization. Among the new research routes explored for hydrothermal processing, the ionic replacement is able to become common knowledge in the near future as alternative processing route, because this reaction promotes a peculiar microstructure in the transformed material that preserves the bulk original geometrical aspects during a single step hydrothermal crystallization reaction. This processing route can lead to prepare pore net-shaped materials with controlled porosity, which depending on their functional properties of the compound, can be used as gas sensor, substrates for porous catalytic materials, filters, and other applications. The conversion process focused toward a particular preparation of a functional inorganic compound from the mineral transformation requires a proper appreciation of the physico-chemical aspects involved for the solid solution-aqueous solution system. Likewise, to evaluate the likely porosity development in any mineral replacement reaction requires knowledge of the coexisting solid and fluid phases involved

On the other hand, the remaining porosity is limited by the differences on the crystalline structure between the parent and the product. This is likely to occur when a large one replaces a small anion and a change on the structural structure also proceeds. One example was found in the case of the conversion of celestite to SrCrO4, this reaction proceeds at low temperatures (200 °C) in relatively alkaline hydrothermal conditions (K2CrO4 solution), with the formation of a peculiar phase on the surface of the partially converted crystals. It is well known that dissolution of mineral species in hydrothermal fluids normally proceeds anisotropically, producing a peculiar texture with holes inside the recrystallized mineral specie (Suaréz-Orduña et. al., 2004a; Rendón-Angeles et. al., 2008). The holes might not be inherited from etch pits produced during the dissolution process, because they did not penetrate the crystals. The factor that has a marked influence in limiting the formation of a residual porosity is related with the replacement of SO4 2- ions by CrO4 2- ions in the SrSO4 crystals. This fact is suggested from the structural change of the orthorhombic to monoclinic structure, which occurs during the conversion. In terms of the global unit cell volume, an expansion process is likely to proceed, therefore, in accordance with the differences on the unit cell volume between that for celestite (312.37 Å3) and that of SrCrO4 (354.11 Å3), a volume increase of 41.74 Å3 is attained in the converted new layer of SrCrO4 as is seen in Figure 18. Hence, the global volumetric unit cell expansion must be compensated by the formation of a continuous solid phase and the formation of some microcracks in this new phase (Rendón-Angeles et. al., 2000b). The formation of a solid phase covering the partially reacted SrSO4 might reduce the transfer of fresh ion exchange media, because under hydrothermal conditions the absence of a texture (small porosity) in the converted phase avoids the penetration of the hydrothermal fluid, coming to an abrupt halt of the replacement reaction

Fig. 17. Variation of the residual porosity on partial and completely converted SrCO3 crystal plates, under hydrothermal conditions at 250 ºC for 96 h in () Na2CO3 and () K2CO3 solutions with different molar CO32-/SO42- ratios. Dotted line= theoretical value calculated from the variation of the cell unit lattice associated with the structural conversion.

Fig. 18. Partially reacted SrSO4 crystal at 200 °C for 96 h in a hydrothermal media with a molar ratio CrO42-/SO42-=1.

## **3. Conclusions**

236 Crystallization – Science and Technology

On the other hand, the remaining porosity is limited by the differences on the crystalline structure between the parent and the product. This is likely to occur when a large one replaces a small anion and a change on the structural structure also proceeds. One example was found in the case of the conversion of celestite to SrCrO4, this reaction proceeds at low temperatures (200 °C) in relatively alkaline hydrothermal conditions (K2CrO4 solution), with the formation of a peculiar phase on the surface of the partially converted crystals. It is well known that dissolution of mineral species in hydrothermal fluids normally proceeds anisotropically, producing a peculiar texture with holes inside the recrystallized mineral specie (Suaréz-Orduña et. al., 2004a; Rendón-Angeles et. al., 2008). The holes might not be inherited from etch pits produced during the dissolution process, because they did not penetrate the crystals. The factor that has a marked influence in limiting the formation of a

crystals. This fact is suggested from the structural change of the orthorhombic to monoclinic structure, which occurs during the conversion. In terms of the global unit cell volume, an expansion process is likely to proceed, therefore, in accordance with the differences on the unit cell volume between that for celestite (312.37 Å3) and that of SrCrO4 (354.11 Å3), a volume increase of 41.74 Å3 is attained in the converted new layer of SrCrO4 as is seen in Figure 18. Hence, the global volumetric unit cell expansion must be compensated by the formation of a continuous solid phase and the formation of some microcracks in this new phase (Rendón-Angeles et. al., 2000b). The formation of a solid phase covering the partially reacted SrSO4 might reduce the transfer of fresh ion exchange media, because under hydrothermal conditions the absence of a texture (small porosity) in the converted phase avoids the penetration of the hydrothermal fluid, coming to an abrupt halt of the replacement reaction

Fig. 17. Variation of the residual porosity on partial and completely converted SrCO3 crystal plates, under hydrothermal conditions at 250 ºC for 96 h in () Na2CO3 and () K2CO3 solutions with different molar CO32-/SO42- ratios. Dotted line= theoretical value calculated

from the variation of the cell unit lattice associated with the structural conversion.

2- ions by CrO4

2- ions in the SrSO4

residual porosity is related with the replacement of SO4

The hydrothermal crystallization of advanced materials is an important branch of chemical science and technology, this process has advantages over conventional technologies, namely the purity of products, quality, and performance, and can be consider as a green chemistry process because is environmentally friendly, because reactions consume lesser energy and these can be carried out under controlled parameters in a closed system. Although, the study of a wide number of oxides have provided physico-chemical experimental information related to solid-aqueous interactions, that affect the dissolution-precipitation mechanism main driving force for achieving bulk solid crystallization, coupled with thermodynamic modelling analysis; have contributed to establish the optimum hydrothermal processing conditions that favours the crystallization proceeds at high yield rates. However, still crystallization research work to carry out on those solid-aqueous systems associated to a specific compound that departs from the thermodynamic modelling limits. In addition, the employment of mineral ore as reactant precursors emerges as an interesting route to explore in order to produce other advanced ceramic compounds via the hydrothermal crystallization. Among the new research routes explored for hydrothermal processing, the ionic replacement is able to become common knowledge in the near future as alternative processing route, because this reaction promotes a peculiar microstructure in the transformed material that preserves the bulk original geometrical aspects during a single step hydrothermal crystallization reaction. This processing route can lead to prepare pore net-shaped materials with controlled porosity, which depending on their functional properties of the compound, can be used as gas sensor, substrates for porous catalytic materials, filters, and other applications. The conversion process focused toward a particular preparation of a functional inorganic compound from the mineral transformation requires a proper appreciation of the physico-chemical aspects involved for the solid solution-aqueous solution system. Likewise, to evaluate the likely porosity development in any mineral replacement reaction requires knowledge of the coexisting solid and fluid phases involved and their relative solubility.

Preparation of Selected Ceramic Compounds by

0897-4756

434, ISSN 0884-2914

708, ISSN 0026-461X

3557, ISSN 0955-2219

ISSN 1529-6466

2003), pp. 1090-1098, ISSN 0897-4756

Controlled Crystallization Under Hydrothermal Conditions 239

Komarneni S., Roy R. & Li, Q. H. (1992). Microwave-Hydrothermal Synthesis of Ceramic Powders, *Materials Research Bulletin*, Vol. 27, pp. 1393-1405, ISSN 0025-5408 Komarneni, S. ; Komarneni Y. S.; Newalkar, B.; Stour, S. (2002). Microwave-Hydrothermal

Kubo, T.; Hogiri, M.; Kagata, H. & Nakahira, A. (2009). Synthesis of Nano-Sized BaTiO3

Lencka, M. M. & Riman, R.E . (1993). Thermodynamic Modeling of Hydrothermal Synthesis

Lencka, M.M. & Riman, R. E. (1995). Thermodynamics of the Hydrothermal Synthesis of

Lencka, M.M. & Riman, R.E. (2002). *Intelligent Systems of Smart Ceramics*, Encyclopedia of Smart Materials, Vol. 1, John Wiley & Sons, Inc., ISBN 9780-471216278, N.J, USA Marchisio, D. L. (2009). On the Use of Bi-Variate Population Balance Equations for

Moreira, M. L.; Mambrini, G. P.; Volanti, D. P.; Leite, E. R.; Orlandi, M. O.; Pizani, V. R.

Nanoparticles*, Chemistry of Materials*, Vol. 20, pp. 5381–5387*, ISSN* 0897-4756 Moon, J.; Kerchener, J. A.; Krarup, H. & Adair, H. (1999). Hydrothermal Synthesis of

Oledzka, M.; Lencka, M. M.; Pincelup, M. P.; Mikulka-Bulen, K.; McCandlish, L. & Riman, R.

Peterson, C. R. & Slamovich, E. B. (1999). Effect of Processing Parameters on the

Putnis A. (2009). Replacement Reactions, *Review Mineral Geochemistry*, Vol. 70, pp. 87-124,

Qi, L.; Lee, B. I.; Badheka, P.; Yoon, D. H.; Samuels, W. D. & Exarhos, G. J., (2004). Short-

*Ceramic Society*, Vol. 82, No. 7, (July1999), pp. 1702-1710, ISSN 0002-7820 Putnis, A. (2002). Mineral Replacement Reactions: From Macroscopic Observations to

*Materials*, Vol. 7, (September 1995), pp. 18-25, *ISSN* 0897-4756

Vol. 64, (February 2009), pp. 697-708, ISSN 0009-2509

*Society*, Vol. 92, No. S1, (January 2009), pp. S172-S176, ISSN 0002-7820 Lee, S. K.; Choi, G.J.; Hwang, U. Y.; Koo, K. K. & Park, T. J. (2003). Effect of Molar Ratio of

Vol. 3, (March 2002), pp. 1025-1032, ISSN 0025-5408

Synthesis of Al-Sustituted Tobermorite from Zeolites. *Materials Research Bulletin,*

Powders by The Rotatory-Hydrothermal Process, *Journal of the American Ceramic* 

KOH to Ti-Isopropoxide on the Formation of BaTiO3 Powders by Hydrothermal Method, *Materials Letters*, Vol. 57, No. 15, (April 2003), pp. 2201-2207, ISSN 0167-577X

of Ceramic Powders, *Chemistry of Materials*, Vol. 5, (October 1993), pp. 61-70, *ISSN*

Calcium Titanate with Reference to Other Alkaline-Earth Titanates, *Chemistry of* 

Modelling Barium Titanate Nanoparticle Precipitation, *Chemical Engineering Science,*

Mastelaro, P.S.; Paiva-Santos, C. O.; Longo, E. & Varela, J. A. (2008). Hydrothermal Microwave: A New Route to Obtain Photoluminescent Crystalline BaTiO3

Ferroelectric Perovskites from Chemically Modified Titanium Isopropoxide and Acetate Salts, *Journal of Materials Research,* Vol. 14, No. 2, (February 1999), pp. 425-

(2003). Influence of Precursor on Microstructure and Phase Composition of Epitaxial Hydrothermal PbZr0.7Ti0.3O3 Films, *Chemistry of Materials,* Vol. 15, (March

Morphology of Hydrothermally Derived PbTiO3 Powders, *Journal of the American* 

Microscopic Mechanisms, *Mineral Magazine*, Vol. 66, No.5, (October 2002), pp. 689-

range Dissolution–Precipitation Crystallization of Hydrothermal Barium Titanate*, Journal of the European Ceramic Society*, Vol. 24, No. 13, (October 2004), pp. 3553–

## **4. Acknowledgment**

One of the authors JCRA wish to acknowledge CONACYT and COECYT for the financial support trough research grants (Project 34830-U) and (FOMIX-COAH 2003-C02-02), respectively. JCRA and ZMV are indebt to the SNI for the financial support. Many thanks are particularly offered to former Ph D. students that collaborate with the experimental work, Roberto Suarez-Orduña, Laura Patricia Vazquez-Rivas and Yadira Marlen Rangel-Hernandez, also to technicians MSc. Martha Rivas-Aguilar and Eng. Felipe de Jesus Márquez-Torres, whom helped in the preparation of samples and its observation by SEM.

## **5. References**


One of the authors JCRA wish to acknowledge CONACYT and COECYT for the financial support trough research grants (Project 34830-U) and (FOMIX-COAH 2003-C02-02), respectively. JCRA and ZMV are indebt to the SNI for the financial support. Many thanks are particularly offered to former Ph D. students that collaborate with the experimental work, Roberto Suarez-Orduña, Laura Patricia Vazquez-Rivas and Yadira Marlen Rangel-Hernandez, also to technicians MSc. Martha Rivas-Aguilar and Eng. Felipe de Jesus Márquez-Torres, whom helped in the preparation of samples and its observation by SEM.

Bilger, S.; Blab, G. & Förthmann, R. (1997). Sol-gel Synthesis of Lanthanum Chromite

Byrappa, K. & Yoshimura, M. (2001). *Handbook of Hydrothermal Technology,* William Andrew

Byrappa, K. (2005). *Hydrothermal Processing,* Kirk-Othmer Encyclopedia of Chemical Technology, pp. Copyright John Wiley & Sons, Inc., ISBN 9780471238966 Chakraborty, A.; Basu, R. N. & Maiti, H. S. (2000). Low Temperature Sintering of

Chen, C. W.; Riman, R. E; TenHuisenb, K. S. & Brown. (2004). Mechanochemical–

Precursors, *Journal of Crystal Growth*, Vol. 270, pp. 615–623, ISSN 0022-0248 Eckert J. O.; Hung-Houston, C. C.; Gersten, B. L.; Lencka, M. M. & Riman, R. E. (1996).

Gersten, B. l.; Lencka, M. M. & Riman R. E. (2004). Low Temperature Hydrothermal

Ianculescu, A.; Braileanu, A.; Pasuk, I. & Zaharescu, M. (2001). Phase Formation Study of

Inagaki, M.; Yamamoto, O. & Hirohara, M., (1990). Synthesis of LaCrO3 from Complex

Kashkurov, K. F., Nikitichev, P. I., Osipov, V. V., Sizova, L. D., & Simonov, A. V. (1968).

Kobe K. A. & Deiglmeier N. J. (1943). Strontium Carbonate, Conversion from Strontium

Publishing/Noyes, ISBN 0-8155-1445-X, New York, USA

(September 2000), pp. 162-166, ISSN 0167-577X

No. 1-2, 12 (May 2009), pp. 373-378, ISSN 0925-8388

2, (November 2001), pp. 501-507, ISSN: 1388-6150

*Crystallografia*, Vol. 12, pp. 837– 839, ISSN 0038-5638

*Chemistry*, Vol. 35, No. 3, pp. 323-325, ISSN 1226-086X

Vol. 98, pp. 675–678, ISSN 1882-0743

Powder, *Journal of the European Ceramic Society*, Vol. 17, No. 8, pp. 1027-1031, ISSN

La(Ca)CrO3 Prepared by an Autoignition Process, *Materials Letters*, Vol. 45, No. 3-4,

Hydrothermal Synthesis of Hydroxyapatite from Nonionic Surfactant Emulsion

Kinetics and Mechanism of Hydrothermal Synthesis of Barium Titanate, *Journal of the American Ceramic Society* Vol. 79, No. 11, (November 1996), pp. 2929-2939, ISSN

Synthesis of Phase-Pure (Br,Sr)TiO3 Perovskite Using EDTA, *Journal of the American Ceramic Society*, Vol. 87, No. 11, (January 2004), pp. 2025-2032, ISSN 0002-7820 Hessien, M. M.; Hassan, M. M. & El-Barawy, K. (2009). Synthesis and Magnetic Properties of

Strontium Hexaferrite from Celestite Ore, *Journal of Alloys and Compounds*, Vol. 476,

Alkaline Earth-doped Lanthanum Chromites, *J. Therm. Anal. Calorim.,* Vol. 66, No.

Precipitation and its Electrical Conductivity, *Journal of the Ceramic Society of Japan*,

Growth of Large Corundum Crystals by the Hydrothermal Method, *Soviet Physics* 

Sulfate by Metathesis with Alkali Carbonate Solution, *Industrial and Engineering* 

**4. Acknowledgment** 

**5. References** 

0955-2219

0002-7820


Preparation of Selected Ceramic Compounds by

0953-8984

Exchange of SO4

Ltd., New Jersey, USA

1403, ISNN 0897-4756

12547-12555, ISSN 0002-7863

2519-2523, ISSN 0022-0248

(November 2008), pp. 11-17, ISSN 0002-7820

Controlled Crystallization Under Hydrothermal Conditions 241

Suárez-Orduña, R.; Rendón-Angeles, J. C.; Matamoros-Veloza, Z. & Yanagisawa, K. (2004b).

Suchanek, W. L.; Shuka, P.; Byrappa, K.; Riman, R. E.; TenHuisen, K. S. & Janas, V. F. (2002).

Suchanek, W. L.; Lencka M. M; & Riman, R. E. (2004). Hydrothermal Synthesis of Ceramics

Shi, S. &. Hwang, J. Y. (2003). Microwave-Assisted Wet Chemical Synthesis: Advantages,

Testino, A.; Buscaglia, V.; Buscaglia, M. T.; Viviani, M. & Nanni, P. (2005). Kinetic Modeling

Tiwary, R. K.; Narayan, S. P. & Pandey, O.P. (2008). Preparation of Strontium Hexaferrite

*and Metallurgy, Section B*, Vol. 44, (February 2008), pp. 91 - 100, ISSN 1450-5339 Vivekanandan, R. & Kutty, T. R. N. (1989). Characterization of Barium Titanate Fine

Wada, S. Suzuki T. & Noma T. (1995). Preparation of Barium Titanate Particles by

Wang, Y.; Xu, G.; Yang, L.; Ren, Z.; Wei, X.; Weng, W.; Du, P.; Shen, G. & Han, G. (2009).

Wei, X.; Xu, G.; Ren, Z.; Xu, C.; Sheng, G. & Han, G. (2008). PVA-Assited Hydrothermal

Wendelbo, R.; Akporiaye, D. E.; Karlsson, A.; Plassen, M. & Olafsen, A. (2006).

*Japan*, Vol. 103, No. 12, (December 1995) pp. 1220-1227, ISSN 1882-0743 Walton, R. I.; Millange, F. Dmith, R. I., Hansen, T. C. & O'hare, D. (2001). Real Time

*Materials*, Vol. 17, (September 2005), pp. 5346-5356*, ISSN* 0897-4756

No. 3, (March 1989), pp. 181 – 192, ISSN 1570-1468

*Characterization & Engineering*, Vol. 2, No.2, pp. 101-110, ISSN 1539-2511 Spooren, J.; Rumplecker, A.; Millange, F. & Walton, R. I. (2003). Subcritical Hydrothermal

Conditions, *Solid State Ionics,* Vol. 172, pp. 393-393, ISSN 0167-2738

Conditions, *Journal of Physics: Condensed Matter*, Vol. 16, pp. S1331-S1344, ISSN

Mechanochemical–hydrothermal synthesis of Carbonated Apatite Powders at Room Temperature, *Biomaterials*, Vol. 23, (April 2001), pp. 699–710, ISSN 0142-9612

Materials, *Aqueous System at Elevated Temperature and Pressure Physical Chemistry in Water, Steam and Hydrothermal Solutions*, Chapter 18, ISBN 0-12-544461-3, Elsevier

Significance, and Steps to Industrialization, *Journal of Mineral & Materials &* 

Synthesis of Perovskite Manganites: A Direct and Rapid Route to Complex Transition-Metal Oxides*, Chemistry of Materials* Vol. 15, (March 2003), pp. 1401–

of Aqueous and Hydrothermal Synthesis of Barium Titanate (BaTiO3), *Chemistry of* 

Magnets from Celestite and Blue Dust by Mechanochemical Route, *Journal of Mining* 

Powders Formed From Hydrothermal Crystallization, *Powder Technology*, Vol. 57,

Hydrothermal Method and Their Characterization, *Journal of the Ceramic Society of* 

Observation of the Hydrothermal Crystallization of Barium Titanate Using In Situ Neutron Powder Difracction. *Journal of the American Chemical Society*, Vol. 123, pp.

Formation of Single-Crystal SrTiO3 Dendritic Nanostructures via a Simple Hydrothermal Method, *Journal of Crystal Growth*, Vol. 311, No. 8, (April 2009), pp.

Sinthesis of SrTiO3 Nanoparticles with Enhanced Photocatalytic Activity for Degradation of RhB, *Journal of the American Ceramic Society*, Vol. 91, No. 11,

Combinatorial Hydrothermal Synthesis and Characterization of Perovskites, *Journal of the European Ceramic Society,* Vol. 26, No. 6, pp. 849-859, ISSN 0955-2219

2- Ions With F- Ions In Mineral Celestite Under Hydrothermal


Ramkrishna, D. & Mahoney, A. W. (2002) Population Balance Modeling Promise for the

Rangel-Hernandez, Y. M.; Rendón-Angeles, J. C.; Matamoros-Veloza, Z.; Pech-Canul, M. I.;

Rivas-Vázquez, L. P.; Rendón-Angeles, J. C.; Rodríguez-Galicia, J. L.; Zhu, K. & Yanagisawa,

Rivas-Vázquez, L. P.; Rendón-Angeles, J. C.; Rodríguez-Galicia, J. L.; Gutiérrez-Chavarria,

Rendón-Angeles, J. C.; Yanagisawa, K.; Ishizawa, N. and Oishi, S. (2000a). Effect of Metal

Rendón-Angeles, J. C.; Yanagisawa, K.; Ishizawa, N. and Oishi, S. (2000b). Topotaxial

Rendón-Angeles, J. C.; Pech-Canul, M. I.; López-Cuevas, J.; Matamoros-Veloza, Z. &

*Chemistry*, Vol. 179, No. 12, (December 2006), pp. 3645-3652, ISSN 0022-4596 Rendón-Angeles, J. C.; Matamoros-Veloza, Z.; López-Cuevas J.; Pech-Canul, M. I &

Rendón-Angeles, J. C., Yanagisawa, K.; Matamoros-Veloza, Z.; Pech-Canul, M. I.; Mendez-

Shi, S. &. Hwang, J. Y. (2003). Microwave-Assisted Wet Chemical Synthesis: Advantages,

*Characterization & Engineering*, Vol. 2, No.2, pp. 101-110, ISSN: 1539-2511 Suárez-Orduña, R; Rendón-Angeles, J. C.; López-Cuevas, J; & Yanagisawa, K. (2004a). The

*Society*, Vol. 26, No. 1-2, (October 2006), pp. 81-88, ISSN 0955-2219

(November 2000), pp. 569-578, ISSN 0022-4596

Schubert, U. and Hüsing, N., 2000. Synthesis of Inorganic Materials

0009-2509

pp. 15-36, ISSN 0151-9107

600, ISSN 0167-2738

0897-4756

0022-2461

Future, *Chemical Engineering Science*, Vol. 57, (February 2002), pp. 595 – 606, ISSN

Diaz-de la Torre, S. & Yanagisawa, K. (2009). One-step Synthesis of Fine SrTiO3 Particles Using SrSO4 Ore Under Alkaline Hydrothermal Conditions, *Chemical Engineering Journal*, Vol. 155, No. 2, (January 2009), pp. 483–492, ISSN 1385-8947 Riman R.E .; Suchanek, W. L. & Lencka, M. M. (2002). Hydrothermal Crystallization of

Ceramics*, Annales de Chimie Science des Materiaux*, Vol. 27, No. 6, (November 2002),

K. (2004). Hydrothermal Synthesis and Sintering of Lanthanum Chromite Powders Doped With Calcium, *Solid State Ionics*, Vol. 172, No. 1-4, (August 2004), pp. 597-

C. A.; Zhu, K. & Yanagisawa, K. (2006). Preparation of Calcium Doped LaCrO3 Fine Powders by Hydrothermal Method and its Sintering, *Journal of the European Ceramic* 

Ions of Chlorapatites on the Topotaxial Replacement by Hydroxyapatite under Hydrothermal Conditions, *Journal of Solid State Chemistry*, Vol. 154, No. 2,

Conversion of Chlorapatite and Hydroxyapatite to Fluorapatite by Hydrothermal Ion Exchange, *Chemistry of Materials* Vol. 12, No. 8, pp. 2143-2150, (June 2000), *ISSN*

Yanagisawa., K. (2006). Differences on the Conversion of Celestite in Solutions Bearing Monovalent Ions Under Hydrothermal Conditions, *Journal of Solid State* 

Yanagisawa, K. (2008). Stability and Direct Conversion of Mineral Barite Crystals in Carbonated Hydrothermal Fluids, *Journal of Material Science*, Vol. 43, pp. 2189, ISSN

Nonell, J. & Diaz-de la Torre, S. (2010). Hydrothermal Synthesis of Perovskite Strontium Doped Lanthanum Chromite Fine Powders and its Sintering, *Journal of Alloys and Compounds*, Vol. 504, No. 1, (August 2010), pp. 251-256, ISSN 0925-8388 Roy, R. (1994). Accelerating the Kinetics of Low-Temperature Inorganic Syntheses*, Journal of Solid State Chemistry,* Vol. 111, (July 1994), pp. 11-17, ISSN 0022-4596

Significance, and Steps to Industrialization, *Journal of Mineral & Materials &* 

Conversion of Mineral Celestite to Strontianite Under Alkaline Hydrothermal

Conditions, *Journal of Physics: Condensed Matter*, Vol. 16, pp. S1331-S1344, ISSN 0953-8984


**9** 

*Serbia* 

**Fe81B13Si4C2 Alloy** 

*2Military Technical Institute, Belgrade,* 

**Fe-Based Nanocomposite Formed by** 

*1University of Belgrade, Faculty for Physical Chemistry,* 

**Thermal Treatment of Rapid-Quenched** 

Dragica M. Minić1, Vladimir A. Blagojević1 and Dušan M. Minić<sup>2</sup>

Amorphous alloys have been a focus of considerable scientific interest, both from fundamental and practical point of view, ever since the first of its kind (Al75Si25) was produced by Klement, Willens and Duwez in 1960 (Klement et al., 1960). It has been shown that the amorphous alloys have features that are different from those of crystalline alloys in both alloy compositions and atomic configurations. This enabled the exhibition of various characteristics which were not obtained for conventional crystalline alloys. Their soft ferromagnetic properties (saturation magnetization, high permeability, low coercivity and loss), high corrosion resistance and good mechanical properties make them suitable for use in a variety of applications, such as power devices, information handling technology, magnetic sensors, anti-theft security systems and construction materials (Minić et al., 2007). Since amorphous alloys are meta-stable, elevated temperature or prolonged performance can induce a transformation into a crystalline state, which could lead to a loss of their advantageous physical properties limiting them to single-use applications. Commercial soft magnetic nanocrystalline materials have recently been successfully obtained by crystallization of amorphous precursors. Materials like this are characterized by a microstructure of nanocrystals embedded into an amorphous matrix, exhibiting superior soft magnetic and mechanical properties to both amorphous and crystalline magnetic alloys (Blagojević et al., 2011; Minić et al., 2011a). This dependence of functional properties on microstructure can be

used to produce functional materials with tailored properties (Maričić et al., 2012).

The production of early amorphous alloys required very high cooling rates (as much as 106 K/s) to avoid crystallization. This limited the form in which they could be produced, as one dimension had to remain small enough to allow sufficiently rapid heat extraction, in order to achieve the necessary high cooling rates. The result was that the thickness of amorphous metal specimens was limited to less than 100μm. In 1976, Liebermann and Graham developed a new method of manufacturing thin ribbons of amorphous metals on a supercooled rapidly rotating disc – the melt-spinning method (Liebermann & Graham, 1976). For the next thirty years, with the production of new materials, the required cooling rate diminished, until, in 1990s, materials were developed, whose production required cooling rates as low as 1 K/s, allowing these alloys to be cast into metallic moulds to

**1. Introduction** 


## **Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy**

Dragica M. Minić1, Vladimir A. Blagojević1 and Dušan M. Minić<sup>2</sup> *1University of Belgrade, Faculty for Physical Chemistry, 2Military Technical Institute, Belgrade, Serbia* 

## **1. Introduction**

242 Crystallization – Science and Technology

Yanagisawa, K.; Rendón-Angeles, J. C.; Kanai H. & Yamashita, Y. (1999). Stability and Single

Yoshimura, M. & Suchanek, W. (1997). In Situ Fabrication of Morphology-Controlled

Yoshino, K.; Nishino, T.; Yoshino, M. & Yoshimura, S. (1985). Exchange Reaction between

Zhang, S.; Han, Y.; Chen, B. & Song, X. (2001). The Influence of TiO2•H2O Gel on

Zhang, S.; Liu, J.; Han, Y.; Chen, B. & Li, X. (2004). Formation Mechanisms of SrTiO3

Zheng, W.; Pang, W.; Meng, G. & Peng, D. (1999). Hydrothermal Synthesis and

1033-1036, ISSN 0955-2219

ISNN 0959-9428

No. 3-4, (June 1997) pp. 197–208, ISSN 0167-2738

(November 2001), pp. 368-370, ISSN 0167-577X

*Kyokai-Shi*, Vol. 93, No. 6, pp. 334-337, ISSN 0372-7718

Vol. 110, No. 1, (June 2004), pp. 11-17, ISSN 0921-5107

Crystal Growth of Dielectric Materials Containing Lead Under Hydrothermal Conditions, *Journal of the European Ceramic Society*, Vol. 19, No. 6-7, (June 1999), pp.

Advanced Ceramic Materials by Soft Solution Processing*, Solid State Ionics*, Vol. 98,

Alkaline Earth Carbonate and Sulfate under Hydrothermal Condition, *Yogyo-*

Hydrothermal Synthesis of SrTiO3 Powders, *Materials Letters*, Vol. 51, No. 4,

Nanoparticles Under Hydrothermal Synthesis, *Materials Science and Engineering B,*

Characterization of LaCrO3, *Journal of Materials Chemistry*, Vol. 9, pp. 2833–2836,

Amorphous alloys have been a focus of considerable scientific interest, both from fundamental and practical point of view, ever since the first of its kind (Al75Si25) was produced by Klement, Willens and Duwez in 1960 (Klement et al., 1960). It has been shown that the amorphous alloys have features that are different from those of crystalline alloys in both alloy compositions and atomic configurations. This enabled the exhibition of various characteristics which were not obtained for conventional crystalline alloys. Their soft ferromagnetic properties (saturation magnetization, high permeability, low coercivity and loss), high corrosion resistance and good mechanical properties make them suitable for use in a variety of applications, such as power devices, information handling technology, magnetic sensors, anti-theft security systems and construction materials (Minić et al., 2007). Since amorphous alloys are meta-stable, elevated temperature or prolonged performance can induce a transformation into a crystalline state, which could lead to a loss of their advantageous physical properties limiting them to single-use applications. Commercial soft magnetic nanocrystalline materials have recently been successfully obtained by crystallization of amorphous precursors. Materials like this are characterized by a microstructure of nanocrystals embedded into an amorphous matrix, exhibiting superior soft magnetic and mechanical properties to both amorphous and crystalline magnetic alloys (Blagojević et al., 2011; Minić et al., 2011a). This dependence of functional properties on microstructure can be used to produce functional materials with tailored properties (Maričić et al., 2012).

The production of early amorphous alloys required very high cooling rates (as much as 106 K/s) to avoid crystallization. This limited the form in which they could be produced, as one dimension had to remain small enough to allow sufficiently rapid heat extraction, in order to achieve the necessary high cooling rates. The result was that the thickness of amorphous metal specimens was limited to less than 100μm. In 1976, Liebermann and Graham developed a new method of manufacturing thin ribbons of amorphous metals on a supercooled rapidly rotating disc – the melt-spinning method (Liebermann & Graham, 1976). For the next thirty years, with the production of new materials, the required cooling rate diminished, until, in 1990s, materials were developed, whose production required cooling rates as low as 1 K/s, allowing these alloys to be cast into metallic moulds to

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 245

The nanocrystalline phase can be identified as α-Fe on matte side and a mixture of α-Fe and Fe3Si on shiny side, with α-Fe being the major component. In addition to sharp crystalline peaks (37, 69 and 78o) in the XRD spectra of the as-prepared alloy, a broad spread halo around 53o, corresponding to domains of short-range ordering in the sample, was also observed. The position of the spread halo corresponds to approximate position of Fe3Si peak and, using the Scherrer equation, we estimated the size of these domains to be 1-1.5 nm. The entire structure could best be described as a combination of nanocrystals and short-range ordered domains embedded in an amorphous matrix. Recent theoretical studies of ironbased binary systems predict existence of short-range ordering in iron-based amorphous alloys (Lass et al., 2010). Therefore, appearance of domains of short-range ordering in

Fig. 2. SEM of as-prepared Fe81B13Si4C2 amorphous alloy (a – shiny; b – matte side)

Crystal structures of α-Fe and Fe3Si are closely related, as Fe3Si crystal system is cubic, same as α-Fe crystal system. Fe3Si lattice is composed of four sub-lattices: three composed of iron atoms and one of silicon atoms, and this leads to doubling of the unit cell (when compared

Fig. 1. X-ray diffraction spectra of as-prepared alloy

Fe81B13Si4C2 amorphous alloy could be expected.

produce specimens up to 100mm thick (Ponnambalam et al., 2004). The alloys that require cooling rates below 103 K/s are known as bulk amorphous alloys, as the production process allows for higher specimen thickness (above 1mm), classifying them as bulk materials.

## **2. Experimental procedures**

The ribbon shaped samples of Fe81B13Si4C2 amorphous alloy were obtained using the standard procedure of rapid quenching of the melt on a rotating disc (melt-spinning method). The obtained ribbon was 2 cm wide and 35 μm thick. During the preparation process of the amorphous alloy ribbon, one of the sides was in direct contact with the cooled rotating disc, while the other was in inert atmosphere. As a result, the two sides of the ribbon show an easily observable difference in reflectivity, surface morphology and structure, as can be seen in X-ray diffraction (XRD) spectra, Fig.1, as well as SEM images of the sample in the as-prepared alloy (Fig.2). The side that was in the contact with the cooled spinning disc is usually labeled as fishy or matte side and the other side, free of contact, as shiny side.

DSC was obtained using SHIMADZU DSC-50 analyzer. In this case, samples weighting several milligrams were heated in the DSC cell from the room temperature to 650oC in a stream of nitrogen with nitrogen flowing at a rate of 20 mL min-1 at the heating rates of 5, 10, 20 and 30oCmin-1.

Mössbauer spectra were taken in the standard transmission geometry using a 57Co(Rh) source at room temperature. The calibration was done against α-Fe foil data. "CONFIT" program package was used for the spectra fitting and decomposition (T. Žák, 1999).

The X-ray diffraction (XRD) patterns were recorded on an X'Pert PROMPD diffractometer from PANalytical with CoKα radiation operated at 40kVand 30mA. For routine characterization, diffraction data was collected in the range of 2θ Bragg angles (15–135o, step 0.0081). For a quantitative analysis and determination of crystallite size from XRD spectra, TOPAS V3 general profile and structure analysis software for powder diffraction data was used (BrukerAXS, general profile and structure analysis software for powder diffraction data, Karlsruhe, 2005). Dislocation density was obtained from the Rietveld analysis, while microstrain was calculated using Williamson-Hall method (Williamson, Hall 1953), using the XRD data. Lattice parameters obtained through XRD spectra were used to calculate the unit cell volumes, which were then compared to the standard values in JCPDS database.

Vickers microhardness tests were performed using MHT-10 (Anton Paar, Austria) microhardness tester, with loads of 0.4N and loading time of 10s (Minić, et al. 2011c). Up to seven measurements were performed on each individual sample, using the average value of microhardness for each sample. The measurements were performed on the cross-section of the ribbons, rather than on any of the sides. Error was calculated as standard deviation for each series of measurements.

### **2.1 Microstructure of as-prepared Fe81B13Si4C2 amorphous alloy**

XRD spectra of the as-prepared alloy ribbon, Fig.1, showed that the as-prepared alloy already had a degree of crystallinity caused by presence of α-Fe phase (JCPDS-PDF 06-0696). The degree of crystallinity is much higher on the shiny side, with peak intensity in XRD spectrum being about 8 times higher on the shiny side.

Fig. 1. X-ray diffraction spectra of as-prepared alloy

produce specimens up to 100mm thick (Ponnambalam et al., 2004). The alloys that require cooling rates below 103 K/s are known as bulk amorphous alloys, as the production process allows for higher specimen thickness (above 1mm), classifying them as bulk materials.

The ribbon shaped samples of Fe81B13Si4C2 amorphous alloy were obtained using the standard procedure of rapid quenching of the melt on a rotating disc (melt-spinning method). The obtained ribbon was 2 cm wide and 35 μm thick. During the preparation process of the amorphous alloy ribbon, one of the sides was in direct contact with the cooled rotating disc, while the other was in inert atmosphere. As a result, the two sides of the ribbon show an easily observable difference in reflectivity, surface morphology and structure, as can be seen in X-ray diffraction (XRD) spectra, Fig.1, as well as SEM images of the sample in the as-prepared alloy (Fig.2). The side that was in the contact with the cooled spinning disc is usually labeled as fishy

DSC was obtained using SHIMADZU DSC-50 analyzer. In this case, samples weighting several milligrams were heated in the DSC cell from the room temperature to 650oC in a stream of nitrogen with nitrogen flowing at a rate of 20 mL min-1 at the heating rates of 5, 10,

Mössbauer spectra were taken in the standard transmission geometry using a 57Co(Rh) source at room temperature. The calibration was done against α-Fe foil data. "CONFIT"

The X-ray diffraction (XRD) patterns were recorded on an X'Pert PROMPD diffractometer from PANalytical with CoKα radiation operated at 40kVand 30mA. For routine characterization, diffraction data was collected in the range of 2θ Bragg angles (15–135o, step 0.0081). For a quantitative analysis and determination of crystallite size from XRD spectra, TOPAS V3 general profile and structure analysis software for powder diffraction data was used (BrukerAXS, general profile and structure analysis software for powder diffraction data, Karlsruhe, 2005). Dislocation density was obtained from the Rietveld analysis, while microstrain was calculated using Williamson-Hall method (Williamson, Hall 1953), using the XRD data. Lattice parameters obtained through XRD spectra were used to calculate the unit cell volumes, which were then compared to the standard values in JCPDS database.

Vickers microhardness tests were performed using MHT-10 (Anton Paar, Austria) microhardness tester, with loads of 0.4N and loading time of 10s (Minić, et al. 2011c). Up to seven measurements were performed on each individual sample, using the average value of microhardness for each sample. The measurements were performed on the cross-section of the ribbons, rather than on any of the sides. Error was calculated as standard deviation for each

XRD spectra of the as-prepared alloy ribbon, Fig.1, showed that the as-prepared alloy already had a degree of crystallinity caused by presence of α-Fe phase (JCPDS-PDF 06-0696). The degree of crystallinity is much higher on the shiny side, with peak intensity in XRD

**2.1 Microstructure of as-prepared Fe81B13Si4C2 amorphous alloy** 

spectrum being about 8 times higher on the shiny side.

program package was used for the spectra fitting and decomposition (T. Žák, 1999).

**2. Experimental procedures** 

20 and 30oCmin-1.

series of measurements.

or matte side and the other side, free of contact, as shiny side.

The nanocrystalline phase can be identified as α-Fe on matte side and a mixture of α-Fe and Fe3Si on shiny side, with α-Fe being the major component. In addition to sharp crystalline peaks (37, 69 and 78o) in the XRD spectra of the as-prepared alloy, a broad spread halo around 53o, corresponding to domains of short-range ordering in the sample, was also observed. The position of the spread halo corresponds to approximate position of Fe3Si peak and, using the Scherrer equation, we estimated the size of these domains to be 1-1.5 nm. The entire structure could best be described as a combination of nanocrystals and short-range ordered domains embedded in an amorphous matrix. Recent theoretical studies of ironbased binary systems predict existence of short-range ordering in iron-based amorphous alloys (Lass et al., 2010). Therefore, appearance of domains of short-range ordering in Fe81B13Si4C2 amorphous alloy could be expected.

Fig. 2. SEM of as-prepared Fe81B13Si4C2 amorphous alloy (a – shiny; b – matte side)

Crystal structures of α-Fe and Fe3Si are closely related, as Fe3Si crystal system is cubic, same as α-Fe crystal system. Fe3Si lattice is composed of four sub-lattices: three composed of iron atoms and one of silicon atoms, and this leads to doubling of the unit cell (when compared

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 247

Fig. 4. Thermomagnetic scans for increasing (dotted line) and decreasing (dashed line)

Measurements of relative magnetic susceptibility were performed using a modified Maxwell method, based on the action of an inhomogeneous field on the magnetic sample. The magnetic force measurements were performed with a sensitivity of 10-6N in an argon

The temperature dependence of the relative magnetic susceptibility of the as-prepared Fe81B13Si4C2 amorphous alloy during three thermal treatments to different temperatures is presented in Fig. 5. During the first and second treatment, the decrease in the magnetic susceptibility in the temperature region from 320oC to 380oC is the result of proximity to the Curie temperature of the amorphous alloy. Before the start of the second treatment it was observed that magnetic susceptibility increased slightly. This was caused by the structural relaxation of an amorphous structure during the first treatment. During this process, internal strains and the free volume are reduced in the starting material. These changes are accompanied by subtle inter-atomic movements, causing the changes in the electron structure and leading to an increase in the number of electrons with unpaired spin in the direction of the outer magnetic field (Minić et al., 2009b). This also leads to a decrease in the number of electrons spinning in the reverse direction and causes an increase in the magnetic susceptibility upon cooling. At the same time, strains and decrease in the free volume enable greater mobility of the walls of the magnetic domains and this behavior further contributes

During the second treatment, the alloy loses its ferromagnetic properties in the temperature region from 400oC to 470oC. With further heating, the magnetic susceptibility starts to rise, and the alloy regains its ferromagnetic properties once the crystallization process starts at about 490oC. After the second heating to 490oC, the magnetic susceptibility decreases by 23 % when compared to the value of as-prepared alloy and to the value of the relaxed state of

temperature.

**2.3.2 Magnetic susceptibility** 

atmosphere (Minić, et al. 2009b).

to the increase in the magnetic susceptibility.

to α-Fe). Fe3Si lattice is, also, slightly distorted, so the value of its lattice parameter is slightly higher than twice the value of lattice parameter of α-Fe.

#### **2.2 Thermal stability of alloy**

The thermal stability of the alloy was investigated using differential scanning calorimetry (DSC) in a nitrogen atmosphere (Minić, et al., 2009a). Typical DSC scan obtained during heating and cooling cycle is presented in Fig. 3. DSC scan involves a series of endothermic and exothermic peaks indicating a stepwise process of structural stabilization of the alloy in the temperature range 170-560oC. A broad exothermic peak, indicated as (*Tsr*), in temperature range 170-400oC, coresponding to structural relaxation, is followed by endothermic hump (temperature of glass transition *Tg*) and a short supercooled liquid region before the sharp exothermic crystallization peak (*Tk*) in temperature range 500-540oC. The enthalpy of crystallization is 83.5 J/g as determined from area of corresponding peak obtained at heating rate of 20 oCmin-1.

#### **2.3 Magnetic properties**

#### **2.3.1 Thermo-magnetic behavior**

Thermally induced processes were also studied using the thermo-magnetic scan, where the sample is heated, annealed and then cooled in vacuum furnace at low magnetic field of 4 kA·m-1 while its magnetic moment is monitored (Minić, et al., 2011b). Both heating and cooling rate were 4oC·min-1, dwell time at maximum temperature of 800oC was 30 minutes. The shape of the thermo-magnetic curve, Fig. 4, reflects changes in the magnetic moment of the sample, caused by phase or structural transitions.

Most pronounced change represents the Curie point (Tc = 420oC), where the magnetization of the respective phase drops to almost zero, because the thermal motion overcomes magnetic interaction. Annealing at the temperature near 200oC is sometimes called stressrelieving (or structural relaxation). Temperatures marked by arrows in Fig. 4 have been identified as points of interest for further study of structural transformations.

Fig. 4. Thermomagnetic scans for increasing (dotted line) and decreasing (dashed line) temperature.

## **2.3.2 Magnetic susceptibility**

246 Crystallization – Science and Technology

to α-Fe). Fe3Si lattice is, also, slightly distorted, so the value of its lattice parameter is slightly

The thermal stability of the alloy was investigated using differential scanning calorimetry (DSC) in a nitrogen atmosphere (Minić, et al., 2009a). Typical DSC scan obtained during heating and cooling cycle is presented in Fig. 3. DSC scan involves a series of endothermic and exothermic peaks indicating a stepwise process of structural stabilization of the alloy in the temperature range 170-560oC. A broad exothermic peak, indicated as (*Tsr*), in temperature range 170-400oC, coresponding to structural relaxation, is followed by endothermic hump (temperature of glass transition *Tg*) and a short supercooled liquid region before the sharp exothermic crystallization peak (*Tk*) in temperature range 500-540oC. The enthalpy of crystallization is 83.5 J/g as determined from area of corresponding peak

Fig. 3. DSC scan of heating and cooling cycle in nitrogen atmosphere; heating rate 10oC/min.

Thermally induced processes were also studied using the thermo-magnetic scan, where the sample is heated, annealed and then cooled in vacuum furnace at low magnetic field of 4 kA·m-1 while its magnetic moment is monitored (Minić, et al., 2011b). Both heating and cooling rate were 4oC·min-1, dwell time at maximum temperature of 800oC was 30 minutes. The shape of the thermo-magnetic curve, Fig. 4, reflects changes in the magnetic moment of

Most pronounced change represents the Curie point (Tc = 420oC), where the magnetization of the respective phase drops to almost zero, because the thermal motion overcomes magnetic interaction. Annealing at the temperature near 200oC is sometimes called stressrelieving (or structural relaxation). Temperatures marked by arrows in Fig. 4 have been

identified as points of interest for further study of structural transformations.

higher than twice the value of lattice parameter of α-Fe.

**2.2 Thermal stability of alloy** 

obtained at heating rate of 20 oCmin-1.

**2.3 Magnetic properties** 

**2.3.1 Thermo-magnetic behavior** 

the sample, caused by phase or structural transitions.

Measurements of relative magnetic susceptibility were performed using a modified Maxwell method, based on the action of an inhomogeneous field on the magnetic sample. The magnetic force measurements were performed with a sensitivity of 10-6N in an argon atmosphere (Minić, et al. 2009b).

The temperature dependence of the relative magnetic susceptibility of the as-prepared Fe81B13Si4C2 amorphous alloy during three thermal treatments to different temperatures is presented in Fig. 5. During the first and second treatment, the decrease in the magnetic susceptibility in the temperature region from 320oC to 380oC is the result of proximity to the Curie temperature of the amorphous alloy. Before the start of the second treatment it was observed that magnetic susceptibility increased slightly. This was caused by the structural relaxation of an amorphous structure during the first treatment. During this process, internal strains and the free volume are reduced in the starting material. These changes are accompanied by subtle inter-atomic movements, causing the changes in the electron structure and leading to an increase in the number of electrons with unpaired spin in the direction of the outer magnetic field (Minić et al., 2009b). This also leads to a decrease in the number of electrons spinning in the reverse direction and causes an increase in the magnetic susceptibility upon cooling. At the same time, strains and decrease in the free volume enable greater mobility of the walls of the magnetic domains and this behavior further contributes to the increase in the magnetic susceptibility.

During the second treatment, the alloy loses its ferromagnetic properties in the temperature region from 400oC to 470oC. With further heating, the magnetic susceptibility starts to rise, and the alloy regains its ferromagnetic properties once the crystallization process starts at about 490oC. After the second heating to 490oC, the magnetic susceptibility decreases by 23 % when compared to the value of as-prepared alloy and to the value of the relaxed state of

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 249

typical for the amorphous volume of the as-prepared sample, while the sharp lines (Fig. 6b) characterize the crystalline structure with well defined position of atoms, which results from the thermally induced crystallization process during annealing of the sample (Minić, et al. 2011b). The computer processing of Mössbauer spectra yielded intensities, *I*, of components, their hyperfine inductions, *Bhf,* isomer shifts, *δ*, and quadrupole splitting, *σ* (T. Žák & Y. Jirásková, 2006). The contents of the iron-containing phases are determined as proportional to the relative areas of the corresponding spectral components. The phase percentages

> α-Fe(Si) at%

as-prepared alloy 0.95 0.03 — — 0.02 — — 200oC/30 min. 0.94 0.02 — — 0.02 0.02 — 450oC/30 min. 0.83 0.14 — 0.03 — — — 500oC/30 min. — 0.42 0.42 0.15 — — 0.01 550 oC/30 min. — 0.52 0.47 — — — 0.01 600oC/30 min. — 0.54 0.45 — — — 0.01 700oC/30 min. — 0.55 0.44 — — — 0.01

Table 1. Mössbauer tentative phase analysis (distribution of Mössbauer iron atoms among

In the as-prepared alloy, the amorphous structure, having a high-field and a low-field component is accompanied by a small amount of α-Fe(Si) solid solution and a FeB phase. Mössbauer phase analysis at higher temperatures reveals α-Fe(Si) solid solution and Fe2B phase to be the most important final crystallization products, although metastable phase Fe3B, was detected initially at 450oC and in higher percentage at 500oC. Amount of iron atoms in paramagnetic positions is almost below the sensitivity threshold. Content of silicon in the α-Fe(Si) solid solution seems to be about 9 at.%, which is close to 7 at.% published in

The electrical resistance of the ribbon was measured using the four-point method within a temperature interval of 20–630oC in an argon atmosphere (Minić, et al., 2011d). Fig. 7a shows the temperature dependence of the electrical resistivity of the alloy in the temperature range of 25-630oC. The dependence clearly shows each structural stabilization step which causes the change in the ordering of the investigated material. These changes are

The slow increase of electrical resistivity was caused by the structural relaxation process in the temperature range of 200-380oC. This process is followed by an increase of electrical resistivity in the vicinity of Curie temperature *Tc* at 420oC, corresponding to the first maximum of the differential curve. At this point the effect that scattering of conductive electrons had on the magnons disappeared (I. Balberg & J. S. Helman, 1978; G. Bohnke et al., 1983) and the amorphous alloy loses its ferromagnetic properties (D. M. Minić et al. 2010). This is in excellent agreement with the results of the thermo-magnetic measurements (Fig.

Fe2B at%

Fe3B at%

FeB at% α-Fe at%

Fe para at%

correspond to the distribution of Mössbauer iron atoms among phases (Table 1).

at%

Annealing temperature Amorphous

phases) (Minić, et al. 2011e)

(Saegusa & Morrish, 1982).

**2.4 Electrical properties 2.4.1 Electrical resistance** 

more obvious in the derivative curve (Fig. 7b).

the lattice after the first heating cycle. During the third treatment, above the crystallization temperature, the alloy maintains its ferromagnetic features in the whole temperature region, whereas the maximum change in the magnetic susceptibility occurs at about 190oC as a consequence of further phase transformation of the crystallized alloy.

Fig. 5. Temperature dependence of relative magnetic susceptibility of as-prepared Fe81B13Si4C2 amorphous alloy during three thermal treatments up to different temperatures: a) 420oC; b) 500oC; c) 630oC.

#### **2.3.3 Mössbauer spectra**

In the Fig. 6, spectra illustrate the ability of Mössbauer effect to distinguish between individual iron-containing phases of different structure. Broad-line components (Fig. 6a) are

Fig. 6. Mössbauer spectra of the as-prepared material (left) and of material after final annealing at 700°C (right), including components of iron containing phases.

typical for the amorphous volume of the as-prepared sample, while the sharp lines (Fig. 6b) characterize the crystalline structure with well defined position of atoms, which results from the thermally induced crystallization process during annealing of the sample (Minić, et al. 2011b). The computer processing of Mössbauer spectra yielded intensities, *I*, of components, their hyperfine inductions, *Bhf,* isomer shifts, *δ*, and quadrupole splitting, *σ* (T. Žák & Y. Jirásková, 2006). The contents of the iron-containing phases are determined as proportional to the relative areas of the corresponding spectral components. The phase percentages correspond to the distribution of Mössbauer iron atoms among phases (Table 1).


Table 1. Mössbauer tentative phase analysis (distribution of Mössbauer iron atoms among phases) (Minić, et al. 2011e)

In the as-prepared alloy, the amorphous structure, having a high-field and a low-field component is accompanied by a small amount of α-Fe(Si) solid solution and a FeB phase. Mössbauer phase analysis at higher temperatures reveals α-Fe(Si) solid solution and Fe2B phase to be the most important final crystallization products, although metastable phase Fe3B, was detected initially at 450oC and in higher percentage at 500oC. Amount of iron atoms in paramagnetic positions is almost below the sensitivity threshold. Content of silicon in the α-Fe(Si) solid solution seems to be about 9 at.%, which is close to 7 at.% published in (Saegusa & Morrish, 1982).

## **2.4 Electrical properties**

248 Crystallization – Science and Technology

the lattice after the first heating cycle. During the third treatment, above the crystallization temperature, the alloy maintains its ferromagnetic features in the whole temperature region, whereas the maximum change in the magnetic susceptibility occurs at about 190oC as a

consequence of further phase transformation of the crystallized alloy.

Fig. 5. Temperature dependence of relative magnetic susceptibility of as-prepared

a) 420oC; b) 500oC; c) 630oC.

**2.3.3 Mössbauer spectra** 

Fe81B13Si4C2 amorphous alloy during three thermal treatments up to different temperatures:

In the Fig. 6, spectra illustrate the ability of Mössbauer effect to distinguish between individual iron-containing phases of different structure. Broad-line components (Fig. 6a) are

Fig. 6. Mössbauer spectra of the as-prepared material (left) and of material after final

annealing at 700°C (right), including components of iron containing phases.

### **2.4.1 Electrical resistance**

The electrical resistance of the ribbon was measured using the four-point method within a temperature interval of 20–630oC in an argon atmosphere (Minić, et al., 2011d). Fig. 7a shows the temperature dependence of the electrical resistivity of the alloy in the temperature range of 25-630oC. The dependence clearly shows each structural stabilization step which causes the change in the ordering of the investigated material. These changes are more obvious in the derivative curve (Fig. 7b).

The slow increase of electrical resistivity was caused by the structural relaxation process in the temperature range of 200-380oC. This process is followed by an increase of electrical resistivity in the vicinity of Curie temperature *Tc* at 420oC, corresponding to the first maximum of the differential curve. At this point the effect that scattering of conductive electrons had on the magnons disappeared (I. Balberg & J. S. Helman, 1978; G. Bohnke et al., 1983) and the amorphous alloy loses its ferromagnetic properties (D. M. Minić et al. 2010). This is in excellent agreement with the results of the thermo-magnetic measurements (Fig.

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 251

thermocouple made by coupling a copper conductor to the amorphous alloy (D.M. Minić, et al. 2009b). The alloy sample was mechanically attached to a copper conductor, forming the Cu – Fe81B13Si4C2 thermocouple, which was placed into a specially designed furnace, while the other end of the sample was submerged into a mixture of water and ice. The TEMF produced by the thermocouple during the heating process was measured by a voltmeter

The temperature dependence of a thermo-electromotor force (Fig. 9) shows three linear regions corresponding to the structural transformations of the alloy. Different slopes of these linear dependences correspond to structural changes involving a structural relaxation, the loss of ferromagnetic properties, and the crystallization, respectively. The temperature

> 1 2 2 1 2

*k N N eN N*

where *k* is Boltzmann's constant, *e* is electron charge, N1(Ef) is the electron state density in

The electron density of states in copper remained unchanged during heating to 680oC, meaning that the change in the temperature coefficient during the heating of the thermocouple was caused only by the change of the electron density of states at the Fermi level of the alloy. Based on the slope of the first linear segment in Fig. 9, temperature coefficient *α1*= 9.4 µV/K, and the relative change in the electron density of states of the alloy

caused by the structural relaxation process was determined to be <sup>1</sup> 3.53% *<sup>N</sup>*

temperature coefficient for the second linear segment is *α2*= 8.36 µV/K, and <sup>2</sup> 5.33% *<sup>N</sup>*

 

*F F F F E E E E*

(1)

*N*

**.** The

*N*

coefficient of TEMF is a function of the electron state density at the Fermi level:

copper and N2(Ef) is the electron state density in the alloy.

Fig. 9. Temperature dependence of thermo-electromotor force.

with the sensitivity of 10-5V.

4). The beginning of crystallization at about 520oC caused a sharp decrease of electrical resistivity. The appearance of two clearly separated maxima, *Tk1* and *Tk2* (490 and 510oC respectively) on the differential curve of electrical resistivity (Fig. 7b), suggests that crystallization of the amorphous alloy is a complex process, occurring in two steps, which appear as a single overlapping peak in the DSC scans, Fig. 3.

Fig. 7. Temperature dependence of the electrical resistivity of amorphous alloy.

The electrical resistivity of the crystalline alloy is lower than that of the amorphous alloy of the same composition, as a result of the increase in electron free path. The linear change of electrical resistivity with increasing temperature during the second treatment shows that crystallization was completed during the first heating cycle (Fig. 8).

Fig. 8. Temperature dependence of the electrical resistivity of second thermal treatment

#### **2.4.2 Thermo-electromotor force**

The structural relaxation processes, as well as the crystallization, in the temperature interval of 25-680oC, were also investigated by measuring the thermo-electromotor force (TEMF) of a

4). The beginning of crystallization at about 520oC caused a sharp decrease of electrical resistivity. The appearance of two clearly separated maxima, *Tk1* and *Tk2* (490 and 510oC respectively) on the differential curve of electrical resistivity (Fig. 7b), suggests that crystallization of the amorphous alloy is a complex process, occurring in two steps, which

Fig. 7. Temperature dependence of the electrical resistivity of amorphous alloy.

crystallization was completed during the first heating cycle (Fig. 8).

**2.4.2 Thermo-electromotor force** 

The electrical resistivity of the crystalline alloy is lower than that of the amorphous alloy of the same composition, as a result of the increase in electron free path. The linear change of electrical resistivity with increasing temperature during the second treatment shows that

Fig. 8. Temperature dependence of the electrical resistivity of second thermal treatment

The structural relaxation processes, as well as the crystallization, in the temperature interval of 25-680oC, were also investigated by measuring the thermo-electromotor force (TEMF) of a

appear as a single overlapping peak in the DSC scans, Fig. 3.

thermocouple made by coupling a copper conductor to the amorphous alloy (D.M. Minić, et al. 2009b). The alloy sample was mechanically attached to a copper conductor, forming the Cu – Fe81B13Si4C2 thermocouple, which was placed into a specially designed furnace, while the other end of the sample was submerged into a mixture of water and ice. The TEMF produced by the thermocouple during the heating process was measured by a voltmeter with the sensitivity of 10-5V.

The temperature dependence of a thermo-electromotor force (Fig. 9) shows three linear regions corresponding to the structural transformations of the alloy. Different slopes of these linear dependences correspond to structural changes involving a structural relaxation, the loss of ferromagnetic properties, and the crystallization, respectively. The temperature coefficient of TEMF is a function of the electron state density at the Fermi level:

$$\alpha = \frac{k}{2e} \left( \frac{N\_{1E\_{\text{F}}}}{N\_{2E\_{\text{F}}}} - \frac{N\_{2E\_{\text{F}}}}{N\_{1E\_{\text{F}}}} \right) \tag{1}$$

*N*

where *k* is Boltzmann's constant, *e* is electron charge, N1(Ef) is the electron state density in copper and N2(Ef) is the electron state density in the alloy.

The electron density of states in copper remained unchanged during heating to 680oC, meaning that the change in the temperature coefficient during the heating of the thermocouple was caused only by the change of the electron density of states at the Fermi level of the alloy. Based on the slope of the first linear segment in Fig. 9, temperature coefficient *α1*= 9.4 µV/K, and the relative change in the electron density of states of the alloy caused by the structural relaxation process was determined to be <sup>1</sup> 3.53% *<sup>N</sup> N* **.** The temperature coefficient for the second linear segment is *α2*= 8.36 µV/K, and <sup>2</sup> 5.33% *<sup>N</sup>* 

Fig. 9. Temperature dependence of thermo-electromotor force.

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 253

and below, and then grow faster than they do on the shiny side. As a consequence, crystal sizes of Fe3Si and α-Fe, respectively, at the end of crystallization are almost the same on both sides of the ribbon. The evolution of average crystal size of Fe2B shows the same trend of smaller initial crystal size and then faster crystal growth on the matte side, with increase in temperature. Metastable Fe3B phase shows the same trend with regards to growth, except, it has a higher average crystal size on the matte side, after the samples are treated at 500oC.

Fig. 10. XRD spectra of alloy samples after thermal treatment (left – shiny side; right – matte

In order to further comprehend the changes in phase composition, unit cell volumes for individual phases were determined using XRD data, and compared with standard values in JCPDS database (Minić, et al. 2011c). This way, we estimated the lattice distortion caused by the presence of boron in α-Fe and Fe3Si lattices. The change in unit cell volume was negative for α-Fe and positive for Fe3Si phase. The diagram (Fig 12) shows that the distortion of the unit cell of α-Fe was the greatest before the crystallization started; it decreased during the crystallization

Fig. 11. XRD spectrum of alloy sample treated at 600oC with individual peak assignments

(left); Rietveld analysis of XRD spectrum of alloy sample treated at 700oC (right)

and remained relatively stable after treatment at 500oC and higher temperatures.

side)

was determined to be 5.33% and for the third linear segment *α3*= 7.12 µV/K and <sup>3</sup> 7.81% *<sup>N</sup> N* was 7.81%. The overall change in the electron density of states at the Fermi level caused by the structural transformations during heating the alloy in temperature range

The increase in the electron density of states at the Fermi level and the above mentioned increase in free path of the electron combine to diminish resistivity of the crystalline alloy (Fig. 8).

#### **2.5 Structural transformations induced by thermal treatment**

25-680oC is the sum of the three ΔN/N values and equals 16.67%.

X-ray diffraction (XRD) spectra of the alloy ribbon samples (Fig. 10) show that thermal treatment (200-700oC) caused a series of structural transformations of the amorphous alloy leading to formation of more than one crystalline phase: the stable α-Fe, Fe3Si and Fe2B as well as metastable Fe3B. Crystallization initially leads to formation of a nanocomposite structure of nanocrystals dispersed in the amorphous matrix. After thermal treatment at 700oC, the alloy ribbon sample was fully crystallized and composed of interdispersed nanocrystals of three crystalline phases: α-Fe, Fe3Si and Fe2B. The analysis of crystallite orientation on the two sides of the ribbon showed that Fe3Si and α-Fe crystallites exhibit a degree of preferential orientation after thermal treatment. In the as-prepared alloy, Fe3Si and α-Fe crystallites on the shiny side are oriented preferentially in [100] direction, while those on the matte side are not. This is probably the reason for higher reflectivity of the shiny side of the ribbon and the consistently high intensity of peak around 78o in the XRD spectra of the shiny side. After the treatment at 500oC, the degree of preferential orientation increases on the shiny side and forms on the matte side, suggesting asymmetric growth of the crystallites. After treatment at 700oC, amplitudes of the degree of preferential orientation on both sides decreased, suggesting that the crystallites were growing in more symmetrical manner. All of these changes are much more pronounced on the shiny side than on the matte side.

Rietveld analysis (Minić, et al. 2011c, Fig. 11) of XRD spectra yielded phase composition of the crystalline portion of the samples, average crystal sizes, dislocation density and microstrain for individual phases (Table 2). R2 was greater than 0.98 for all XRD measurements. Individual phase contributions of α-Fe and Fe3Si could not be completely separated, because of the overlap of their peaks. The phase composition data shows that shiny side has higher percentage of metastable Fe3B phase, but lower percentage of Fe2B phase. After heating at 700oC, the final phase composition on the matte side shows more crystalline Fe2B than on the shiny side. The phase content of Fe3B declines in two steps, with increase in heating temperature, to disappear completely after thermal treatment at 650oC. The first sharp decrease in percentage of Fe3B (550oC) coincides with an increase in percentage of combined α-Fe and Fe3Si. After thermal treatment at 650oC, Fe3B phase disappears completely, and this coincides with increase in phase content of Fe2B. Phase content of Fe2B phase shows an increase after treatment at temperatures above 550oC.

The evolution of average crystal sizes for the respective phases shows that while there is some difference between shiny and matte side of the ribbon at the onset of crystallization, those differences become almost negligible after heating at 700oC. On closer examination, Fe3Si and α-Fe have lower crystal sizes on the matte side, at heating temperatures of 500oC

was determined to be 5.33% and for the third linear segment *α3*= 7.12 µV/K and

was 7.81%. The overall change in the electron density of states at the Fermi

level caused by the structural transformations during heating the alloy in temperature range

The increase in the electron density of states at the Fermi level and the above mentioned increase in free path of the electron combine to diminish resistivity of the crystalline alloy

X-ray diffraction (XRD) spectra of the alloy ribbon samples (Fig. 10) show that thermal treatment (200-700oC) caused a series of structural transformations of the amorphous alloy leading to formation of more than one crystalline phase: the stable α-Fe, Fe3Si and Fe2B as well as metastable Fe3B. Crystallization initially leads to formation of a nanocomposite structure of nanocrystals dispersed in the amorphous matrix. After thermal treatment at 700oC, the alloy ribbon sample was fully crystallized and composed of interdispersed nanocrystals of three crystalline phases: α-Fe, Fe3Si and Fe2B. The analysis of crystallite orientation on the two sides of the ribbon showed that Fe3Si and α-Fe crystallites exhibit a degree of preferential orientation after thermal treatment. In the as-prepared alloy, Fe3Si and α-Fe crystallites on the shiny side are oriented preferentially in [100] direction, while those on the matte side are not. This is probably the reason for higher reflectivity of the shiny side of the ribbon and the consistently high intensity of peak around 78o in the XRD spectra of the shiny side. After the treatment at 500oC, the degree of preferential orientation increases on the shiny side and forms on the matte side, suggesting asymmetric growth of the crystallites. After treatment at 700oC, amplitudes of the degree of preferential orientation on both sides decreased, suggesting that the crystallites were growing in more symmetrical manner. All of

these changes are much more pronounced on the shiny side than on the matte side.

Rietveld analysis (Minić, et al. 2011c, Fig. 11) of XRD spectra yielded phase composition of the crystalline portion of the samples, average crystal sizes, dislocation density and microstrain for individual phases (Table 2). R2 was greater than 0.98 for all XRD measurements. Individual phase contributions of α-Fe and Fe3Si could not be completely separated, because of the overlap of their peaks. The phase composition data shows that shiny side has higher percentage of metastable Fe3B phase, but lower percentage of Fe2B phase. After heating at 700oC, the final phase composition on the matte side shows more crystalline Fe2B than on the shiny side. The phase content of Fe3B declines in two steps, with increase in heating temperature, to disappear completely after thermal treatment at 650oC. The first sharp decrease in percentage of Fe3B (550oC) coincides with an increase in percentage of combined α-Fe and Fe3Si. After thermal treatment at 650oC, Fe3B phase disappears completely, and this coincides with increase in phase content of Fe2B. Phase content of Fe2B phase shows an increase after treatment at temperatures above 550oC.

The evolution of average crystal sizes for the respective phases shows that while there is some difference between shiny and matte side of the ribbon at the onset of crystallization, those differences become almost negligible after heating at 700oC. On closer examination, Fe3Si and α-Fe have lower crystal sizes on the matte side, at heating temperatures of 500oC

25-680oC is the sum of the three ΔN/N values and equals 16.67%.

**2.5 Structural transformations induced by thermal treatment** 

<sup>3</sup> 7.81% *<sup>N</sup>*

*N*

(Fig. 8).

and below, and then grow faster than they do on the shiny side. As a consequence, crystal sizes of Fe3Si and α-Fe, respectively, at the end of crystallization are almost the same on both sides of the ribbon. The evolution of average crystal size of Fe2B shows the same trend of smaller initial crystal size and then faster crystal growth on the matte side, with increase in temperature. Metastable Fe3B phase shows the same trend with regards to growth, except, it has a higher average crystal size on the matte side, after the samples are treated at 500oC.

Fig. 10. XRD spectra of alloy samples after thermal treatment (left – shiny side; right – matte side)

In order to further comprehend the changes in phase composition, unit cell volumes for individual phases were determined using XRD data, and compared with standard values in JCPDS database (Minić, et al. 2011c). This way, we estimated the lattice distortion caused by the presence of boron in α-Fe and Fe3Si lattices. The change in unit cell volume was negative for α-Fe and positive for Fe3Si phase. The diagram (Fig 12) shows that the distortion of the unit cell of α-Fe was the greatest before the crystallization started; it decreased during the crystallization and remained relatively stable after treatment at 500oC and higher temperatures.

Fig. 11. XRD spectrum of alloy sample treated at 600oC with individual peak assignments (left); Rietveld analysis of XRD spectrum of alloy sample treated at 700oC (right)

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 255

Fig. 12. Change in unit cell volumes of α-Fe and Fe3Si against change in phase contents of

Microhardness was measured on the cross-section of the alloy ribbon samples, rather than on any of the sides, giving a good measure of the average properties of the alloy samples. This was done due to extreme brittleness of the ribbon samples after crystallization. In order to clearly show the influence of the change in microstructure, induced by thermal treatment, on microhardness, we presented the microhardness data in combination with DSC scan (Fig. 12). In terms of microstructure, the as-prepared alloy contains a small percentage (less than 5%) of crystalline α-Fe phase in form of nanocrystals dispersed in the amorphous matrix (Fig. 1) and small domains (1-1.5 nm in size) of short-range crystalline ordering. This structure can be best described as a nanocomposite of nanocrystals and nanoclusters embedded in the amorphous matrix. The nanocomposite structure combines with chemical composition, involving significant percentage of boron, silicon and carbon, all of which are known to increase hardness in iron alloys, to produce high hardness of the as-prepared alloy

The change of microhardness with respect to heating temperature showed three distinct temperature regions with completely different behavior. Before the onset of crystallization around 450oC, microhardness exhibited slight growth, from 909 to 931HV, which corresponds to lattice relaxation, as shown by the broad exothermic peaks in 200-400oC region in the DSC (Fig. 13a). In the second region, after thermal treatment at 450oC, it increased to 951HV and then to 1250HV, after thermal treatment at 500oC. This corresponds to crystallization involving formation of disordered Fe3Si crystal phase (around 450oC), stable crystalline phases Fe3Si, Fe2B, α-Fe and metastable Fe3B (exothermic peak around 512oC). These phases form a nanocomposite of nanocrystals dispersed in the amorphous matrix, which is the main cause of the increased microhardness. In the last temperature region, after the sample was treated at 700oC, microhardness decreased to 908HV, as a consequence of further crystal growth, leading to the change in the nature of the interfaces.

Fe2B and Fe3B

sample (909HV).

**2.6 Mechanical properties** 

The change in Fe3B percentage after treatment at 550oC was accompanied by a decreased distortion of the unit cell of both α-Fe and Fe3Si and increase in their combined phase content. This indicates that iron from Fe3B was transformed into these two phases, with boron being incorporated back into the α-Fe/Fe3Si matrix. As Fe3Si showed greater change, Fe3B probably transformed more to Fe3Si than it did to α-Fe.


Table 2. Phase composition and dislocation density in the samples, obtained by analysis of XRD data

Reduced lattice distortion in these phases is probably due to the fact that mass percentage of boron in Fe3B (6.06%) is lower than the mass percentage of boron in the as-prepared alloy (13%). Therefore, transformation of Fe3B would lead to dilution of boron in α-Fe/Fe3Si matrix, leading to stabilization of its crystal structure. The change in unit cell volumes of Fe3Si and α-Fe showed increase at higher heating temperatures and it is possible that the distortion of the lattice in α-Fe and Fe3Si is caused by the increasing occurrence of crystal/crystal interfaces between α-Fe and Fe3Si on one side and Fe2B on the other.

Fig. 12. Change in unit cell volumes of α-Fe and Fe3Si against change in phase contents of Fe2B and Fe3B

#### **2.6 Mechanical properties**

254 Crystallization – Science and Technology

The change in Fe3B percentage after treatment at 550oC was accompanied by a decreased distortion of the unit cell of both α-Fe and Fe3Si and increase in their combined phase content. This indicates that iron from Fe3B was transformed into these two phases, with boron being incorporated back into the α-Fe/Fe3Si matrix. As Fe3Si showed greater change,

Phase composition (% mass of crystalline phase) 500 84± 4 8± 2 8± 4 83± 5 6± 2 11± 2 550 88± 3 8± 2 4± 1 88± 3 8± 2 4± 1 600 86± 3 10± 2 4± 1 89± 3 8± 2 3± 1

700 87± 2 13± 2 - 90± 2 10± 2 Average crystal size (nm)

700 44.6 71.9 107.5 - 44.9 67.5 133.5 Dislocation density (1015 m-2)

700 1.51 0.58 0.26 - 1.49 0.66 0.17 Microstrain (%)

Table 2. Phase composition and dislocation density in the samples, obtained by analysis of

Reduced lattice distortion in these phases is probably due to the fact that mass percentage of boron in Fe3B (6.06%) is lower than the mass percentage of boron in the as-prepared alloy (13%). Therefore, transformation of Fe3B would lead to dilution of boron in α-Fe/Fe3Si matrix, leading to stabilization of its crystal structure. The change in unit cell volumes of Fe3Si and α-Fe showed increase at higher heating temperatures and it is possible that the distortion of the lattice in α-Fe and Fe3Si is caused by the increasing occurrence of

crystal/crystal interfaces between α-Fe and Fe3Si on one side and Fe2B on the other.

500 26.7 22.3 49.3 16 33.1 35.1 37.3 15.8 550 32.1 35.4 90.6 46.4 36.4 45.6 82.7 44.8 600 38.6 51.6 97.7 56.6 39.0 52.7 111.1 67.7

500 4.21 6.03 1.23 11.72 2.74 2.43 2.16 12.02 550 2.91 2.39 0.37 1.39 2.26 1.44 0.44 1.50 600 2.01 1.13 0.31 0.94 1.97 1.08 0.24 0.66

25 26.2 36.9 200 24.6 36.5 450 19.7 32.9

25 4.37 2.2 200 4.96 2.25 450 7.73 2.77

25 2.19 2.47 200 1.85 2.63

500 2.53 3.56 4.73 550 2.00 4.27 2.26 3.76 600 1.94 4.03 1.70 5.69 700 1.26 3.47 1.37 2.66

Matte Shiny α-Fe Fe3Si Fe2B Fe3B α-Fe Fe3Si Fe2B Fe3B

Fe3B probably transformed more to Fe3Si than it did to α-Fe.

Temperature oC

XRD data

Microhardness was measured on the cross-section of the alloy ribbon samples, rather than on any of the sides, giving a good measure of the average properties of the alloy samples. This was done due to extreme brittleness of the ribbon samples after crystallization. In order to clearly show the influence of the change in microstructure, induced by thermal treatment, on microhardness, we presented the microhardness data in combination with DSC scan (Fig. 12). In terms of microstructure, the as-prepared alloy contains a small percentage (less than 5%) of crystalline α-Fe phase in form of nanocrystals dispersed in the amorphous matrix (Fig. 1) and small domains (1-1.5 nm in size) of short-range crystalline ordering. This structure can be best described as a nanocomposite of nanocrystals and nanoclusters embedded in the amorphous matrix. The nanocomposite structure combines with chemical composition, involving significant percentage of boron, silicon and carbon, all of which are known to increase hardness in iron alloys, to produce high hardness of the as-prepared alloy sample (909HV).

The change of microhardness with respect to heating temperature showed three distinct temperature regions with completely different behavior. Before the onset of crystallization around 450oC, microhardness exhibited slight growth, from 909 to 931HV, which corresponds to lattice relaxation, as shown by the broad exothermic peaks in 200-400oC region in the DSC (Fig. 13a). In the second region, after thermal treatment at 450oC, it increased to 951HV and then to 1250HV, after thermal treatment at 500oC. This corresponds to crystallization involving formation of disordered Fe3Si crystal phase (around 450oC), stable crystalline phases Fe3Si, Fe2B, α-Fe and metastable Fe3B (exothermic peak around 512oC). These phases form a nanocomposite of nanocrystals dispersed in the amorphous matrix, which is the main cause of the increased microhardness. In the last temperature region, after the sample was treated at 700oC, microhardness decreased to 908HV, as a consequence of further crystal growth, leading to the change in the nature of the interfaces.

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 257

where *α* is the fractional extent of reaction (conversion degree), *t* is time and the function *f*(α)

The temperature dependence of the rate conversion is introduced by replacing *k*(*T*) with the

<sup>d</sup>

where *A* (pre-exponential factor) and *E* (activation energy) are the Arrhenius parameters

Kinetic description of solid state transformations usually includes a kinetic triplet, involving Arrhenius parameters (activation energy, *E* and pre-exponential factor, *A*) as well as an algebraic expression of the conversion function, *f*(*α*) (presented in Table 3), which describes

In solid state reactions, the constant value of activation energy can be expected only for a single-step reaction, therefore *E* in equation (3) becomes an apparent quantity (*Ea*), based on a quasi-single-step reaction. In non-isothermal measurements at constant heating rate, *β,* the

<sup>d</sup>

*Ea A f <sup>T</sup> RT*

<sup>d</sup> ( ) ( ) ( ) *AEa g p <sup>x</sup> f R*

 

where *p(x)* is the temperature integral for *x=Ea/RT* which does not have analytical solution.

exp <sup>d</sup>

The integral form of the reaction model can be obtained by integration of equation (4)

0

 

 

(5)

(4)

(3)

*<sup>E</sup> A f <sup>t</sup> RT*

exp <sup>d</sup>

the dependence of the reaction rate on the conversion degree, *α*.

depends on the particular crystallization mechanism.

Arrhenius equation, which gives the relation

and *R* is the gas constant.

equation (3) transforms to:

where dα/d*t* ≡ *β* (dα/d*T*).

Fig. 14. DSC curves at different heating rates

Fig. 13. Microhardness shown with DSC scan (a) and average crystal size on matte side (b).

These changes in microhardness can also be correlated with change in microstructure, induced by thermal treatment, through evolution of average crystal size of the crystalline phases in the alloy (Fig. 13b) (Minić, et al. 2011c). It can be seen that average crystal size has a significant influence on the change in microhardness. When the formed nanocrystals are relatively small (below 50 nm), microhardness remains high, while appearance of larger nanocrystals (over 100 nm) leads to a sharp decrease in microhardness. After thermal treatment at temperatures from 200 to 700oC, the alloy structure gradually transformed from a relatively homogeneous to a granulated structure with larger crystalline domains. This is consistent with creation, at the beginning of the crystallization, of a nanocomposite of small nanocrystals dispersed in amorphous matrix, which significantly increases microhardness, due to the fact that the dominant type of interface in the alloy is crystal/amorphous, as opposed to crystal/crystal in a completely crystalline alloy. As the average crystal size increased, crystal/crystal interfaces became dominant and microhardness decreased, due to increased interfacial energy and more successful propagation of shear bands and cracks along these interfaces. During the course of the observed structural transformations, the overall composition of the alloy did not change drastically, while microhardness showed significant fluctuation. This means that the changes of microhardness in our alloy were caused primarily by changes in its microstructure, rather than changes in its composition, with the main factor being the change in the average crystal size and creation of a granulated structure, as opposed to nanocrystals embedded in amorphous matrix (Table 2). The granulated structure, in addition to containing larger crystals and more crystal/crystal interfaces, is also much more porous than the original nanocomposite crystal/amorphous structure of the as-prepared alloy.

#### **2.7 Kinetics of crystallization**

All kinetic studies assume that the isothermal rate of conversion dα/d*t* is a linear function of the temperature-dependent reaction rate constant, *k* (*T*), and a temperature-independent function of the conversion, *f* (α) (Vyazovkin, 2000)

$$\frac{da}{dt} = k(T)f(a) \,. \tag{2}$$

where *α* is the fractional extent of reaction (conversion degree), *t* is time and the function *f*(α) depends on the particular crystallization mechanism.

The temperature dependence of the rate conversion is introduced by replacing *k*(*T*) with the Arrhenius equation, which gives the relation

$$\frac{\mathrm{d}a}{\mathrm{d}t} = A \exp\left(-\frac{E}{RT}\right) f(a) \tag{3}$$

where *A* (pre-exponential factor) and *E* (activation energy) are the Arrhenius parameters and *R* is the gas constant.

Kinetic description of solid state transformations usually includes a kinetic triplet, involving Arrhenius parameters (activation energy, *E* and pre-exponential factor, *A*) as well as an algebraic expression of the conversion function, *f*(*α*) (presented in Table 3), which describes the dependence of the reaction rate on the conversion degree, *α*.

In solid state reactions, the constant value of activation energy can be expected only for a single-step reaction, therefore *E* in equation (3) becomes an apparent quantity (*Ea*), based on a quasi-single-step reaction. In non-isothermal measurements at constant heating rate, *β,* the equation (3) transforms to:

$$
\beta \frac{\mathrm{d}a}{\mathrm{d}T} = A \exp\left(-\frac{E\_a}{RT}\right) f(a) \tag{4}
$$

where dα/d*t* ≡ *β* (dα/d*T*).

256 Crystallization – Science and Technology

Fig. 13. Microhardness shown with DSC scan (a) and average crystal size on matte side (b).

These changes in microhardness can also be correlated with change in microstructure, induced by thermal treatment, through evolution of average crystal size of the crystalline phases in the alloy (Fig. 13b) (Minić, et al. 2011c). It can be seen that average crystal size has a significant influence on the change in microhardness. When the formed nanocrystals are relatively small (below 50 nm), microhardness remains high, while appearance of larger nanocrystals (over 100 nm) leads to a sharp decrease in microhardness. After thermal treatment at temperatures from 200 to 700oC, the alloy structure gradually transformed from a relatively homogeneous to a granulated structure with larger crystalline domains. This is consistent with creation, at the beginning of the crystallization, of a nanocomposite of small nanocrystals dispersed in amorphous matrix, which significantly increases microhardness, due to the fact that the dominant type of interface in the alloy is crystal/amorphous, as opposed to crystal/crystal in a completely crystalline alloy. As the average crystal size increased, crystal/crystal interfaces became dominant and microhardness decreased, due to increased interfacial energy and more successful propagation of shear bands and cracks along these interfaces. During the course of the observed structural transformations, the overall composition of the alloy did not change drastically, while microhardness showed significant fluctuation. This means that the changes of microhardness in our alloy were caused primarily by changes in its microstructure, rather than changes in its composition, with the main factor being the change in the average crystal size and creation of a granulated structure, as opposed to nanocrystals embedded in amorphous matrix (Table 2). The granulated structure, in addition to containing larger crystals and more crystal/crystal interfaces, is also much more porous than the original nanocomposite crystal/amorphous

All kinetic studies assume that the isothermal rate of conversion dα/d*t* is a linear function of the temperature-dependent reaction rate constant, *k* (*T*), and a temperature-independent

*<sup>d</sup> kT f dt*

, (2)

structure of the as-prepared alloy.

function of the conversion, *f* (α) (Vyazovkin, 2000)

**2.7 Kinetics of crystallization** 

The integral form of the reaction model can be obtained by integration of equation (4)

$$\log(\alpha) = \int\_0^a \frac{\mathbf{d}\,\alpha}{f(\alpha)} = \frac{AE\_a}{R\beta} \, p(\mathbf{x}) \tag{5}$$

where *p(x)* is the temperature integral for *x=Ea/RT* which does not have analytical solution.

Fig. 14. DSC curves at different heating rates

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 259

of specific surface area of nuclei. The saturation part that follows is the consequence of

Fig. 15. Fractional conversion (α) as a function of temperature (*T*) for the crystallization alloy

The overall activation energy of crystallization of an amorphous alloy under linear heating condition can be determined by the Kissinger as well as by the Ozawa peak method relating

Kissinger (Kissinger, 1957) proposed that the activation energy can be determined according

<sup>2</sup> ln ln *<sup>a</sup> p a*

For the determination the activation energy in non-isothermal conditions Ozawa (Ozawa,

The values of peak temperatures together with values of kinetic parameters (*Ea* and ln*A*)

ln ln 1.0516

calculated by both methods and the symmetry factors (SF) are given in the Table 4.

*T E RT*

*AE E a a C R RT*

*p*

(7)

*AR E*

(6)

nuclei merging together causing a decrease in their surface area.

at different heating rates (5, 10, 20 and 30oCmin-1).

the equation

1970) proposed the equation:

**2.7.1 The overall apparent activation energy of crystallization** 

to the dependence of exothermic peak temperature *Tp* on heating rate *β*.

The different algebraic expressions of conversion functions for solid state transformations are given in Table 3.

Fig. 14 shows the continuous DSC curves of Fe81B13Si4C2 ribbon in temperature range 300- 650oC taken at four different heating rates. All curves contain a single well formed exothermic peak representing crystallization in the temperature range 500-560oC.


Table 3. Algebraic expressions of conversion functions for solid state transformations

All values of initial (*Ti*), maximal (*Tp*) and final temperature (*Tf*) for both exo-peaks are shifted to higher values with increasing heating rate indicating the presence of kinetic effects (Table 4). The shape of DSC curves depends on the heating rate, because, as the values of shape factor *S* show, asymmetry of peaks increases with the decrease of heating rates indicating that the heating rate has a strong influence on the crystallization process. The values of shape factor *S* is obtained as the ratio of half-widths (left and right) for individual crystallization peaks for particular heating rate. The fractional extent of the sample transformed into crystalline phase, α, has been obtained from the DSC curve as a function of temperature (*T*) (Minić & Adnađević, 2008). At any temperature *T,* α is defined as α = *ST*/*S*, where *S* is the total area of the exotherm, between the temperature *Ti*, of the onset of crystallization and the temperature *Tf*, of the end of crystallization, and *ST* is the area between the initial temperature *Ti* and a generic temperature, *T*, ranging between *Ti* and *Tf*  (Friedman, 1964). The plots of α versus *T* at different heating rates for the considered crystallization process are shown in Fig. 15.

The sigmoid shape of fractional conversion curves in Fig. 15 indicates, for all heating rates, a slow initial period corresponding to nucleation with increasing rate caused by the increase

The different algebraic expressions of conversion functions for solid state transformations

Fig. 14 shows the continuous DSC curves of Fe81B13Si4C2 ribbon in temperature range 300- 650oC taken at four different heating rates. All curves contain a single well formed

All values of initial (*Ti*), maximal (*Tp*) and final temperature (*Tf*) for both exo-peaks are shifted to higher values with increasing heating rate indicating the presence of kinetic effects (Table 4). The shape of DSC curves depends on the heating rate, because, as the values of shape factor *S* show, asymmetry of peaks increases with the decrease of heating rates indicating that the heating rate has a strong influence on the crystallization process. The values of shape factor *S* is obtained as the ratio of half-widths (left and right) for individual crystallization peaks for particular heating rate. The fractional extent of the sample transformed into crystalline phase, α, has been obtained from the DSC curve as a function of temperature (*T*) (Minić & Adnađević, 2008). At any temperature *T,* α is defined as α = *ST*/*S*, where *S* is the total area of the exotherm, between the temperature *Ti*, of the onset of crystallization and the temperature *Tf*, of the end of crystallization, and *ST* is the area between the initial temperature *Ti* and a generic temperature, *T*, ranging between *Ti* and *Tf*  (Friedman, 1964). The plots of α versus *T* at different heating rates for the considered

The sigmoid shape of fractional conversion curves in Fig. 15 indicates, for all heating rates, a slow initial period corresponding to nucleation with increasing rate caused by the increase

crystallization process are shown in Fig. 15.

exothermic peak representing crystallization in the temperature range 500-560oC.

Mechanism label *f*(α) *g*(α) Power law P4 4α3/4 α1/4 Power law P3 3α2/3 α1/3 Power law P2 2α1/2 α1/2 Power law P3/2 3/2α1/3 α2/3 Avrami-Erofeev A3/2 3/2(1-α)[-ln(1-α)]1/3 [-ln(1-α)]2/3 Avrami-Erofeev A2 2(1-α)[-ln(1-α)]1/2 [-ln(1-α)]1/2 Avrami-Erofeev A3 3(1-α)[-ln(1-α)]3/2 [-ln(1-α)]1/3 Avrami-Erofeev A4 4(1-α)[-ln(1-α)]3/4 [-ln(1-α)]1/4 Prout-Tompkins B1 α(1-α) ln[α(1-α)-1] Šesták-Berggren αN(1-α)M - One dimensional phase boundary R1 1 α Contracting cylinder R2 2(1-α)1/2 1-(1-α)1/2 Contracting sphere R3 3(1-α)2/3 1-(1-α)1/3 D1 One-dimensional diffusion 1/2α α<sup>2</sup> D2 Two-dimensional diffusion [-ln(1-α)]-1 (1-α)ln(1-α)+α D3 three-dimensional diffusion 3/2(1-α)2/3[1-(1-α)1/3]-1 [1-(1-α)1/3]2 D4 Ginstling-Brounshtein 3/2[(1-α)-1/3-1]-1 (1-2α/3)-(1-α)2/3 F1 First-order 1-α -ln(1-α) F2 Second-order (1-α)2 (1-α)-1 F3 Third-order 1/2(1-α)3 (1-α)-2 Table 3. Algebraic expressions of conversion functions for solid state transformations

are given in Table 3.

of specific surface area of nuclei. The saturation part that follows is the consequence of nuclei merging together causing a decrease in their surface area.

Fig. 15. Fractional conversion (α) as a function of temperature (*T*) for the crystallization alloy at different heating rates (5, 10, 20 and 30oCmin-1).

#### **2.7.1 The overall apparent activation energy of crystallization**

The overall activation energy of crystallization of an amorphous alloy under linear heating condition can be determined by the Kissinger as well as by the Ozawa peak method relating to the dependence of exothermic peak temperature *Tp* on heating rate *β*.

Kissinger (Kissinger, 1957) proposed that the activation energy can be determined according the equation

$$\ln\left(\frac{\mathcal{J}}{T\_p^2}\right) = \ln\left(\frac{AR}{E\_a}\right) - \frac{E\_a}{RT} \tag{6}$$

For the determination the activation energy in non-isothermal conditions Ozawa (Ozawa, 1970) proposed the equation:

$$
\ln \beta = \ln \frac{AE\_a}{R} C - 1.0516 \frac{E\_a}{RT\_p} \tag{7}
$$

The values of peak temperatures together with values of kinetic parameters (*Ea* and ln*A*) calculated by both methods and the symmetry factors (SF) are given in the Table 4.

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 261

On the basis of dynamic DSC measurements for different heating rates, isoconversional method of Kissinger-Akahira-Sunose method, also known as the "model-free method", enables the determination values of *Ea* over a wide range of α without the determination of the conversion function (Kissinger, 1957; Akahira, Sunose, 1971). This model involves measuring the temperatures *Tα* corresponding to fixed values of the crystallized volume

, , .

where a subscript α designates values related to a given conversion degree, and *i* is a

The left-hand side of Eq. (8) is linear with respect to the inverse temperature, 1/*Tα*, and enables the activation energy to be evaluated using a linear regression method. In case of a single step process, a constant value of *Ea*(*<sup>α</sup>*) is obtained. On the other hand, the dependence of *Ea*(*<sup>α</sup>*) on α indicates complex process involving more than one step with different activation energies. It can be observed that the apparent activation energy for the considered crystallization process (Fig. 17) is practically constant in the 0.05 ≤ α ≤ 0.7 range indicating the existence of a single-step reaction. (Vyazovkin, 2000; Opfermann & Flammersheim, 2003). The average value of apparent activation energy was determined as *Ea* = 356.5 5.5

Fig. 17. Apparent activation energies (*Ea*) and the intercepts as function of the crystallized

For the preliminary determination of kinetic model of the crystallization process,

**2.7.3 Preliminary determination of kinetic model** 

Dollimore's method was used (D. Dollimore, 1991, 1992; Lee, 1998).

*i a i*

*AR E F T E RT*

ln ln ( ) *<sup>i</sup> <sup>a</sup>*

,

(8)

**2.7.2 The dependence of apparent activation energy on range of conversion** 

fraction, *α*, for different heating rates, *β*, and application of the relation:

2

kJmol-1.

fraction α.

number of the non-isothermal experiment conducted at the heating rate *βi*.


Table 4. Values of *Ti*, *Tp*, *Tf* and *S* for crystallization of the amorphous Fe81B13Si4C2 alloy upon continuous heating at different heating rates.

Fig. 16. The dependence of maximum rate of crystallization on heating rate.

The dependence of maximum of rate on rate of heating (Fig. 16) shows the growth of maximum of rate of crystallization with respect to the rate of heating. It indicates that the rate of crystallization reaches a saturation point at high heating rates.

The activation energy of crystallization process involving formation of nuclei and their growth, according to some opinions, has no physical meaning, just empirical significance and only establishes the dependence of the rate of conversion on temperature. This energy can be spent, not just for overcoming the activation barrier but also for its downturn due to cooperative displacement of groups of atoms. Finally, the crystallization of amorphous alloys is a very complicated process accompanied by nucleation and growth of various crystal phases under continuously varied conditions in the conversion zone. With the multitude of possible ways of conversions, only those mechanisms and activated complexes of the crystallization process will be realized that have the highest probability at a given temperature. Any change in crystallization conditions, such as heating rate, can result in change of the mechanism and main activation complex of the crystallization process. Thus, high values of activation energy of crystallization of amorphous alloys, first of all, indicate that a large number of atoms participate in an elementary act of structure reorganization, as well as high complexity of the transformation processes.

#### **2.7.2 The dependence of apparent activation energy on range of conversion**

260 Crystallization – Science and Technology

Ea kJmol-1

338.0 +1.8 1.10 ± 2.3 351.2±1.8 3.06 ±2.3 10 501 520 546 0.59

Table 4. Values of *Ti*, *Tp*, *Tf* and *S* for crystallization of the amorphous Fe81B13Si4C2 alloy upon

Fig. 16. The dependence of maximum rate of crystallization on heating rate.

rate of crystallization reaches a saturation point at high heating rates.

well as high complexity of the transformation processes.

The dependence of maximum of rate on rate of heating (Fig. 16) shows the growth of maximum of rate of crystallization with respect to the rate of heating. It indicates that the

The activation energy of crystallization process involving formation of nuclei and their growth, according to some opinions, has no physical meaning, just empirical significance and only establishes the dependence of the rate of conversion on temperature. This energy can be spent, not just for overcoming the activation barrier but also for its downturn due to cooperative displacement of groups of atoms. Finally, the crystallization of amorphous alloys is a very complicated process accompanied by nucleation and growth of various crystal phases under continuously varied conditions in the conversion zone. With the multitude of possible ways of conversions, only those mechanisms and activated complexes of the crystallization process will be realized that have the highest probability at a given temperature. Any change in crystallization conditions, such as heating rate, can result in change of the mechanism and main activation complex of the crystallization process. Thus, high values of activation energy of crystallization of amorphous alloys, first of all, indicate that a large number of atoms participate in an elementary act of structure reorganization, as

Ozawa Kissinger

Ea kJmol-1 A × 1021 min-1

A × 1022 min-1

 oCmin-1

Ti oC

Tp oC

5 492 512 542 0.59

20 509 531 554 0.70 30 513 538 560 0.75

continuous heating at different heating rates.

Tf oC <sup>S</sup>

On the basis of dynamic DSC measurements for different heating rates, isoconversional method of Kissinger-Akahira-Sunose method, also known as the "model-free method", enables the determination values of *Ea* over a wide range of α without the determination of the conversion function (Kissinger, 1957; Akahira, Sunose, 1971). This model involves measuring the temperatures *Tα* corresponding to fixed values of the crystallized volume fraction, *α*, for different heating rates, *β*, and application of the relation:

$$\ln\left(\frac{\beta\_i}{T\_{a,i}^2}\right) = \ln\left(\frac{AR}{E\_{a,\alpha}}F(\alpha)\right) - \frac{E\_{a,\alpha}}{RT\_{a,i}}\tag{8}$$

where a subscript α designates values related to a given conversion degree, and *i* is a number of the non-isothermal experiment conducted at the heating rate *βi*.

The left-hand side of Eq. (8) is linear with respect to the inverse temperature, 1/*Tα*, and enables the activation energy to be evaluated using a linear regression method. In case of a single step process, a constant value of *Ea*(*<sup>α</sup>*) is obtained. On the other hand, the dependence of *Ea*(*<sup>α</sup>*) on α indicates complex process involving more than one step with different activation energies. It can be observed that the apparent activation energy for the considered crystallization process (Fig. 17) is practically constant in the 0.05 ≤ α ≤ 0.7 range indicating the existence of a single-step reaction. (Vyazovkin, 2000; Opfermann & Flammersheim, 2003). The average value of apparent activation energy was determined as *Ea* = 356.5 5.5 kJmol-1.

Fig. 17. Apparent activation energies (*Ea*) and the intercepts as function of the crystallized fraction α.

#### **2.7.3 Preliminary determination of kinetic model**

For the preliminary determination of kinetic model of the crystallization process, Dollimore's method was used (D. Dollimore, 1991, 1992; Lee, 1998).

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 263

Table 5. It is clearly seen that the position of the broadening exotherm, which is connected with the crystallization, was shifted toward higher temperature with the increase of the heating rate as well as asymmetry of peaks. This suggests that the crystallization process should not be characterized by a definite critical temperature independent of the heating rate. The determined values of αmax for different heating rates were in the range from 0.51 to 0.55. These results indicate that the non-isothermal crystallization mechanism of α-Fe in amorphous Fe81B13Si4C2 alloy can not be fully described within the JMA (Johnson-Mehl-

General equation, enabling the analysis of conversion kinetics involving nucleation and

( ) 1 exp[ ( ) ] *<sup>n</sup>*

Differentiation of this equation with respect to time gives the rate equation, known as the

<sup>d</sup> 1 1 (1 )[ ln(1 )] <sup>d</sup> *<sup>n</sup> kn*

The JMA equation is based on assumptions of isothermal crystallization, homogenous or heterogeneous nucleation at randomly dispersed particles of the second phase. The growth rate of new phase is independent of time and controlled by temperature and low anisotropy of growing crystals. However if the entire nucleation process takes place during the early stage of transformation and becomes negligible afterwards, JMA equation can also be

The validity of listed assumptions is not given a priori and simple and reliable testing methods were developed (Málek, 1992, 1995; Gotor et al, 2000; Criado et al, 2003). Once the apparent activation energy has been determined, it is possible to find the kinetic model which best describes the measured set of thermoanalytical data. It can be shown that, for this purpose, it is useful to define two special functions y(α) and z(α), which can easily be obtained by simple transformation of the experimental data. The conversions, in which the

respectively. Under non-isothermal conditions these functions can be expressed as follows

 exp *Ea <sup>d</sup> <sup>y</sup> Af dt RT* 

 *<sup>d</sup>* <sup>2</sup> *z T dt* 

The maximum of the y(α) function for the JMA model depends on the value of the kinetic

y(α) and z(α) functions exhibit the maximum values are designated as

  

*t kt* , (9)

. (10)

 *y* and

(11)

(12)

 *z* ,

Avrami) models (A2, A3 and A4, Group A).

**2.7.4 Determination of the kinetic model** 

where 

JMA equation:

(Málek, 1992, 1995, 2000):

exponent:

growth in solid phase was proposed by Avrami (Avrami, 1939):

*t* 

applied to non-isothermal conditions (Henderson, 1979).

*(t)* is conversion degree, *n* is kinetic exponent, *k=ko*exp(*-Ea /RT*).


Table 5. Parameters describing the asymmetric DSC peaks of crystallization α-Fe phase in amorphous Fe81B13Si4C2 alloy

This model is based on the "sharpness" of initial and final temperature of differential rate curves, Fig. 18, as well as on its asymmetry. The "sharpness" of the initial and final temperatures is influenced by kinetic factors, and especially by the mechanism of the process. Certain kinetic models lead to an asymptotic or diffuse departure from the base line in the differential form of the thermal curve, while others produce a very sharp approach to the final plateau. The investigation of these parameters that describe geometry and asymmetry of the differential rate curves can be an indicator of the probable kinetic mechanism. According to these parameters, different types of kinetic mechanisms have been listed (Lee, 1998). So, when the crystallization process is not complex the qualitative approach to its kinetics may be obtained using parameters such as α*max* or (dα/dt)*max*, the shape of the initial and final temperatures as well as peak temperature (*Tp*) or half-width from differential rate curves.

Fig. 18. Differential rate curves (dα/d*t* vs. *T*) for different heating rates (β = 5, 10, 20 and 30oCmin-1).

In this case, Dollimore's procedure was applied to the conversion and differential rate curves (Fig.18) whose slight asymmetry is observed between *Ti* and *Tf*. The other parameters such as the conversion at the rate of maximum crystallization, α*max*, peak temperature, *Tp*, at (dα/dt)*max*, and the ratio ∆LoT/∆HiT (shape factor), which is the ratio between the low and high temperature points at half-width of the differential rate curve peak are presented in Table 5. It is clearly seen that the position of the broadening exotherm, which is connected with the crystallization, was shifted toward higher temperature with the increase of the heating rate as well as asymmetry of peaks. This suggests that the crystallization process should not be characterized by a definite critical temperature independent of the heating rate. The determined values of αmax for different heating rates were in the range from 0.51 to 0.55. These results indicate that the non-isothermal crystallization mechanism of α-Fe in amorphous Fe81B13Si4C2 alloy can not be fully described within the JMA (Johnson-Mehl-Avrami) models (A2, A3 and A4, Group A).

#### **2.7.4 Determination of the kinetic model**

262 Crystallization – Science and Technology

5 0.51 1.0 9.5 sharp sharp 10 0.53 1.0 9.5 sharp sharp 20 0.53 0.9 10.0 sharp sharp 30 0.55 0.8 11.0 sharp sharp Table 5. Parameters describing the asymmetric DSC peaks of crystallization α-Fe phase in

This model is based on the "sharpness" of initial and final temperature of differential rate curves, Fig. 18, as well as on its asymmetry. The "sharpness" of the initial and final temperatures is influenced by kinetic factors, and especially by the mechanism of the process. Certain kinetic models lead to an asymptotic or diffuse departure from the base line in the differential form of the thermal curve, while others produce a very sharp approach to the final plateau. The investigation of these parameters that describe geometry and asymmetry of the differential rate curves can be an indicator of the probable kinetic mechanism. According to these parameters, different types of kinetic mechanisms have been listed (Lee, 1998). So, when the crystallization process is not complex the qualitative approach to its kinetics may be obtained using parameters such as α*max* or (dα/dt)*max*, the shape of the initial and final temperatures as well as peak temperature (*Tp*) or half-width

Fig. 18. Differential rate curves (dα/d*t* vs. *T*) for different heating rates (β = 5, 10, 20 and

In this case, Dollimore's procedure was applied to the conversion and differential rate curves (Fig.18) whose slight asymmetry is observed between *Ti* and *Tf*. The other parameters such as the conversion at the rate of maximum crystallization, α*max*, peak temperature, *Tp*, at (dα/dt)*max*, and the ratio ∆LoT/∆HiT (shape factor), which is the ratio between the low and high temperature points at half-width of the differential rate curve peak are presented in

Half-width (oC)

Ti Tf

max *LoT*

*HiT* 

β (oC/min)

amorphous Fe81B13Si4C2 alloy

from differential rate curves.

30oCmin-1).

General equation, enabling the analysis of conversion kinetics involving nucleation and growth in solid phase was proposed by Avrami (Avrami, 1939):

$$a(t) = 1 - \exp[-(kt)^{\pi}] \, \text{ } \tag{9}$$

where *(t)* is conversion degree, *n* is kinetic exponent, *k=ko*exp(*-Ea /RT*).

Differentiation of this equation with respect to time gives the rate equation, known as the JMA equation:

$$
\left(\frac{\mathrm{d}\alpha}{\mathrm{d}t}\right) = kn(1-\alpha)[-\ln(1-\alpha)]^{1-1/\mu} \cdot \tag{10}
$$

The JMA equation is based on assumptions of isothermal crystallization, homogenous or heterogeneous nucleation at randomly dispersed particles of the second phase. The growth rate of new phase is independent of time and controlled by temperature and low anisotropy of growing crystals. However if the entire nucleation process takes place during the early stage of transformation and becomes negligible afterwards, JMA equation can also be applied to non-isothermal conditions (Henderson, 1979).

The validity of listed assumptions is not given a priori and simple and reliable testing methods were developed (Málek, 1992, 1995; Gotor et al, 2000; Criado et al, 2003). Once the apparent activation energy has been determined, it is possible to find the kinetic model which best describes the measured set of thermoanalytical data. It can be shown that, for this purpose, it is useful to define two special functions y(α) and z(α), which can easily be obtained by simple transformation of the experimental data. The conversions, in which the y(α) and z(α) functions exhibit the maximum values are designated as *y* and *z* , respectively. Under non-isothermal conditions these functions can be expressed as follows (Málek, 1992, 1995, 2000):

$$y(\alpha) = \left(\frac{d\alpha}{dt}\right) \exp\left(\frac{E\_a}{RT}\right) = A f\left(\alpha\right) \tag{11}$$

$$z(a) \approx \left(\frac{da}{dt}\right)T^2\tag{12}$$

The maximum of the y(α) function for the JMA model depends on the value of the kinetic exponent:

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 265

crystallization kinetics, where the crystallized phase further increases the rate of the crystallization. Such autocatalytic behavior can be well described using an empirical two parameter Šesták-Berggren's kinetic model (Šesták, Berggren, 1971; Málek, et al. 1989). This

1 *<sup>M</sup> <sup>N</sup> f*

In this case the expression for reaction rate of the investigated crystallization process can be

exp <sup>1</sup> *<sup>d</sup> Ea <sup>M</sup> <sup>N</sup> <sup>A</sup>*

For this model, the ratio of the kinetic exponents *p* = *M* / *N* can be calculated from the

 1 *y y*

 

. (16)

(14)

. (15)

 

), it is possible to obtain the kinetic exponent *<sup>N</sup>*

 

, (13)

 

*dt RT*

*M <sup>p</sup> <sup>N</sup>*

ln exp ln ln 1 *<sup>d</sup> Ea <sup>p</sup> A N*

This equation describes the processes of nucleation and growth in non-crystalline solids very well. The parameters *M* and *N* define relative contributions of acceleratory and decay regions of the kinetic process. From the linear dependence

and the pre-exponential factor, ln *A*. The value of kinetic exponent *M* can be obtained

oC/min *M N* ln*A (\*1022)*

Table 7. Kinetic exponents *M* and *N* at different heating rates (Adnađević et al., 2010)

Table 7 lists the values of kinetic exponents *M* and *N*, as well as the values of ln*A* obtained by the procedure described above, for the considered crystallization process at different

5 0.75 ± 0.03 1.08 ± 0.10 9.0±0.2 10 0.66 ± 0.05 0.92 ± 0.05 10.7±0.3 20 0.64 ± 0.05 0.89 ± 0.07 9.4±0.2 30 0.81 ± 0.10 1.17 ± 0.04 10.6±0.2 Average 0.72 ± 0.06 1.02 ± 0.07 9.9±0.2

*dt RT*

 

model is based on the equation:

given as:

where *M* and *N* represents the kinetic exponents.

maximum of the y (α) function (Málek, 2000):

=f( ln (1 ) *<sup>p</sup>*

ln / exp( / ) *<sup>a</sup> d dt E RT*

directly from eq. (15).

β

heating rates.

Introducing this equation in equation this (13) gives:

 *y* = 0 for n ≤ 1 *y* =1-exp (n-1-1) for n > 1

The value of *y* is always lower than the maximum of value for *z* . For JMA model, *z* =0.632. This value is characteristic "fingerprint" of the JMA model and it can be used as a simple test of the applicability of this model.

The obtained normalized functions y(α) and z(α), Fig 19, are independent on the heating rate (*β*) of the system, and the both functions exhibits the well-defined maxima which were located at exactly defined value of α ( *y* for the y(α) function and *z* for z(α) function, respectively), Table 6.


Table 6. The maximum of αy and αz for the different heating rates

Fig. 19. Normalized *y*(*α*) and *z*(*α*) functions at the different heating rates.

From Table 6, it can be seen that the values of *y* fall into the range *y* (0, *<sup>z</sup>* ) (0.41 ≤ *y* ≤ 0.42) and the values of *z* are less than 0.632 (0.51 ≤ *z* ≤ 0.55). From the obtained results, it follows that the conditions of validity of the JMA model are not fulfilled for crystallization of *α*-Fe. The displacement of *z* in lower value range indicates complexity of the process and can be caused by the influence of surface nucleation or the effect of released crystallization heat on temperature distribution within the sample. However, the relatively high value of *y* indicates an increasing effect of the crystallized phase to overall crystallization kinetics, where the crystallized phase further increases the rate of the crystallization. Such autocatalytic behavior can be well described using an empirical two parameter Šesták-Berggren's kinetic model (Šesták, Berggren, 1971; Málek, et al. 1989). This model is based on the equation:

$$f(a) = a^{\mathcal{M}} \left(1 - a\right)^{\mathcal{N}},\tag{13}$$

where *M* and *N* represents the kinetic exponents.

264 Crystallization – Science and Technology

= 0 for n ≤ 1

=1-exp (n-1-1) for n > 1

=0.632. This value is characteristic "fingerprint" of the JMA model and it can be used as a

The obtained normalized functions y(α) and z(α), Fig 19, are independent on the heating rate (*β*) of the system, and the both functions exhibits the well-defined maxima which were

for the y(α) function and

*z*

> *z*

. For JMA model,

for z(α) function,

is always lower than the maximum of value for

*y*

oC min-1 <sup>α</sup>y\* αz\*

5 0.41 ± 0.01 0.53 ± 0.01 10 0.42 ± 0.01 0.51 ± 0.01 20 0.42 ± 0.01 0.55 ± 0.03 30 0.41 ± 0.01 0.52 ± 0.01

*y*

*y*

Table 6. The maximum of αy and αz for the different heating rates

Fig. 19. Normalized *y*(*α*) and *z*(*α*) functions at the different heating rates.

are less than 0.632 (0.51 ≤

*y*

it follows that the conditions of validity of the JMA model are not fulfilled for crystallization

and can be caused by the influence of surface nucleation or the effect of released crystallization heat on temperature distribution within the sample. However, the relatively

fall into the range

in lower value range indicates complexity of the process

*z*

indicates an increasing effect of the crystallized phase to overall

 *y* (0,*<sup>z</sup>* ) (0.41 ≤

≤ 0.55). From the obtained results,

 *y* 

From Table 6, it can be seen that the values of

*z*

> *z*

≤ 0.42) and the values of

high value of

of *α*-Fe. The displacement of

*y*

The value of

respectively), Table 6.

*z* *y*

simple test of the applicability of this model.

located at exactly defined value of α (

In this case the expression for reaction rate of the investigated crystallization process can be given as:

$$\frac{da}{dt} = A \exp\left(-\frac{E\_a}{RT}\right) a^M \left(1 - a\right)^N \tag{14}$$

For this model, the ratio of the kinetic exponents *p* = *M* / *N* can be calculated from the maximum of the y (α) function (Málek, 2000):

$$p = \frac{M}{N} = \frac{a\_y^\*}{\left(1 - a\_y^\*\right)}\tag{15}$$

Introducing this equation in equation this (13) gives:

$$\ln\left[\left(\frac{d\alpha}{dt}\right)\exp\left(\frac{E\_a}{RT}\right)\right] = \ln A + N\ln\left[\alpha^p \left(1 - \alpha\right)\right].\tag{16}$$

This equation describes the processes of nucleation and growth in non-crystalline solids very well. The parameters *M* and *N* define relative contributions of acceleratory and decay regions of the kinetic process. From the linear dependence ln / exp( / ) *<sup>a</sup> d dt E RT* =f( ln (1 ) *<sup>p</sup>* ), it is possible to obtain the kinetic exponent *<sup>N</sup>* and the pre-exponential factor, ln *A*. The value of kinetic exponent *M* can be obtained directly from eq. (15).


Table 7. Kinetic exponents *M* and *N* at different heating rates (Adnađević et al., 2010)

Table 7 lists the values of kinetic exponents *M* and *N*, as well as the values of ln*A* obtained by the procedure described above, for the considered crystallization process at different heating rates.

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 267

already exist at *T* ≤ 500oC and at *T* ≥ 500oC these embryos are momentarily transformed into nuclei. The established acceleration of crystallization process is a consequence of significant

For non-isothermal crystallization, where the volume fraction of crystalline phase precipitated

1.052 ln[ ln(1 )] ln *mEa <sup>n</sup> const*

 

where *m* and *n* are constants with values between one and four depending on the

The values of *n* obtained from the slopes of linear plots ln[-ln (1-α)] versus –ln*β* at different temperatures for considered crystallization process are given in Table 9. For all considered temperatures, the value of *n* is ≈ 4.0, within the limits of experimental errors. It follows, then,

The crystallization exponent *n* is connected with the number of growth dimensions (*m*) and the number of nuclei forming stages (*s*) (Matusita, Sakka, 1979) by the following equation

where *m* is the number of growth dimensions as defined in Table 8, *s* is the number of the nuclei forming stages (*s* = 0 – at instantaneously nucleation; *s* = 1 – at constant nucleation

In order to describe the crystallization process in detail, the value of parameter *m* should be determined from the plot of ln [-ln (1-α)], because a function of reciprocal temperature is linear with a slope of 1.051× (*m* + 1)*Ea*/*R,* using the value of activation energy determined above.

Table 9. The values of *n* at three temperatures and values of *m* and *s* for four different

*n β*

518 3.92 5 2.84 1.16 520 4.08 10 3.08 0.92 522 4.07 20 3.22 0.78

oC/min

*RT*

is related with the activation energy *Ea*, Matusita et al.

, (18)

*nms* (19)

*m s* 

30 2.84 1.16

increase of strain in alloy, which arises on account of formation of α-Fe.

morphology and kinetics of the growth nuclei, Table 8.

rate and *s* > 1 at self-acceleratory nucleation rate).

Temperature oC

heating rates

proposed the following relation (Matusita & S. Sakka, 1979, 1980; Matusita, et al. 1984):

Mechanism *n m* 

Table 8. Values of constants *n* and *m* for different crystallization mechanisms

that the kinetics of crystallization process is independent from the temperature.

**2.7.5 Morphology of crystal growth** 

in glass heated at a uniform heating rate

 Three dimensional growth Two dimensional growth One dimensional growth

Bulk nucleation

Surface nucleation

The obtained values of kinetic exponents *M* and *N* vary a little with respect to heating rate *β.*  The values of M vary in the range of 0.64 ≤ *M* ≤ 0.81 with average value of *M*av = 0.72. The values of N vary in the range of 0.89 ≤ *N* ≤ 1.17 with average value of *N*av = 1.02. The values of the pre-exponential factor (ln*A*) are independent on the heating rate (*β*), within the limits of the experimental error. It was shown that this two parameter autocatalytic model has physical meaning only for M<1(Gotor, et al. 2000).

In order to check the established kinetic model we applied the "Master-plot" method (Criado, et. al 2003; Gotor, 2000). Using the value at α = 0.5 as a reference point, the following differential master equation is easily derived from eq. (4):

$$\frac{f\left(\alpha\right)}{f\left(0.5\right)} = \frac{d\alpha \;/\, dt}{\left(d\alpha \;/\, dt\right)\_{0.5}} \frac{\exp\left(E\_a \;/\, RT\right)}{\exp\left(E\_a \;/\, RT\_{0.5}\right)}\tag{17}$$

where (dα/d*t*)0.5, *T*0.5 and *f* (0.5) are the reaction rate, the temperature reaction and the differential conversion function, respectively at *α* = 0.5.

The left side of Eq. (17) is a reduced theoretical curve which is characteristic of each kinetic function. The right side of the equation is associated with the reduced rate and can be obtained from experimental data if the apparent activation energy is known and remains constant throughout the reaction. Comparison of both sides of Eq. (17) tells us which kinetic model describes an experimental reaction process. It can be seen in Fig. 20, using the average value for the apparent activation energy determined from Kissinger-Akahira-Sunose isoconversional method, the suggested kinetic model works very well in the entire conversion range. Therefore, the Šesták-Berggren autocatalytic model represents the best reaction model for describing the crystallization process of α-Fe in the amorphous Fe81B13Si4C2 alloy. The higher value of *N* exponent designates that the formed crystallized phase has the decisive influence on the kinetics of transformation and the rate of growth. In the propagation process, on account of overlapping of nuclei during growth, it acts to retard crystallization rate. Bearing this in mind, we can suppose that in the amorphous alloy, the α-Fe embryos

Fig. 20. The theoretical (solid line) and experimental differential master plots of *f*(α)/*f*(0.5) versus α for different heating rates: (□) 5 oCmin-1; (○) 10 oCmin-1; (▲) 20 oCmin-1 and (◊) 30 oCmin-1.

already exist at *T* ≤ 500oC and at *T* ≥ 500oC these embryos are momentarily transformed into nuclei. The established acceleration of crystallization process is a consequence of significant increase of strain in alloy, which arises on account of formation of α-Fe.

## **2.7.5 Morphology of crystal growth**

266 Crystallization – Science and Technology

The obtained values of kinetic exponents *M* and *N* vary a little with respect to heating rate *β.*  The values of M vary in the range of 0.64 ≤ *M* ≤ 0.81 with average value of *M*av = 0.72. The values of N vary in the range of 0.89 ≤ *N* ≤ 1.17 with average value of *N*av = 1.02. The values of the pre-exponential factor (ln*A*) are independent on the heating rate (*β*), within the limits of the experimental error. It was shown that this two parameter autocatalytic model has

In order to check the established kinetic model we applied the "Master-plot" method (Criado, et. al 2003; Gotor, 2000). Using the value at α = 0.5 as a reference point, the

0.5 / exp /

*f d dt E RT f d dt E RT*

where (dα/d*t*)0.5, *T*0.5 and *f* (0.5) are the reaction rate, the temperature reaction and the

The left side of Eq. (17) is a reduced theoretical curve which is characteristic of each kinetic function. The right side of the equation is associated with the reduced rate and can be obtained from experimental data if the apparent activation energy is known and remains constant throughout the reaction. Comparison of both sides of Eq. (17) tells us which kinetic model describes an experimental reaction process. It can be seen in Fig. 20, using the average value for the apparent activation energy determined from Kissinger-Akahira-Sunose isoconversional method, the suggested kinetic model works very well in the entire conversion range. Therefore, the Šesták-Berggren autocatalytic model represents the best reaction model for describing the crystallization process of α-Fe in the amorphous Fe81B13Si4C2 alloy. The higher value of *N* exponent designates that the formed crystallized phase has the decisive influence on the kinetics of transformation and the rate of growth. In the propagation process, on account of overlapping of nuclei during growth, it acts to retard crystallization rate. Bearing this in mind, we can suppose that in the amorphous alloy, the α-Fe embryos

Fig. 20. The theoretical (solid line) and experimental differential master plots of *f*(α)/*f*(0.5) versus α for different heating rates: (□) 5 oCmin-1; (○) 10 oCmin-1; (▲) 20 oCmin-1 and (◊) 30 oCmin-1.

(17)

0.5 0.5

*a a*

/ exp /

physical meaning only for M<1(Gotor, et al. 2000).

following differential master equation is easily derived from eq. (4):

differential conversion function, respectively at *α* = 0.5.

 For non-isothermal crystallization, where the volume fraction of crystalline phase precipitated in glass heated at a uniform heating rate is related with the activation energy *Ea*, Matusita et al. proposed the following relation (Matusita & S. Sakka, 1979, 1980; Matusita, et al. 1984):

$$
\ln[-\ln(1-\alpha)] = -n\ln\beta - \frac{1.052mE\_a}{RT} + \text{const.}\tag{18}
$$

where *m* and *n* are constants with values between one and four depending on the morphology and kinetics of the growth nuclei, Table 8.


Table 8. Values of constants *n* and *m* for different crystallization mechanisms

The values of *n* obtained from the slopes of linear plots ln[-ln (1-α)] versus –ln*β* at different temperatures for considered crystallization process are given in Table 9. For all considered temperatures, the value of *n* is ≈ 4.0, within the limits of experimental errors. It follows, then, that the kinetics of crystallization process is independent from the temperature.

The crystallization exponent *n* is connected with the number of growth dimensions (*m*) and the number of nuclei forming stages (*s*) (Matusita, Sakka, 1979) by the following equation

$$m = m + s \tag{19}$$

where *m* is the number of growth dimensions as defined in Table 8, *s* is the number of the nuclei forming stages (*s* = 0 – at instantaneously nucleation; *s* = 1 – at constant nucleation rate and *s* > 1 at self-acceleratory nucleation rate).

In order to describe the crystallization process in detail, the value of parameter *m* should be determined from the plot of ln [-ln (1-α)], because a function of reciprocal temperature is linear with a slope of 1.051× (*m* + 1)*Ea*/*R,* using the value of activation energy determined above.


Table 9. The values of *n* at three temperatures and values of *m* and *s* for four different heating rates

Fe-Based Nanocomposite Formed by Thermal Treatment of Rapid-Quenched Fe81B13Si4C2 Alloy 269

Dollimore, D.; Evans, T.A.; Lee, Y.F.; Pee, G.P. & Wilburn, F.W. (1992), The significance of

Friedman, H. L. (1964), Kinetics of thermal degradation of char-forming plastics from

Gotor, F. J.; Criado, J. M.; Málek, J. & Koga, N. (2000), Kinetic Analysis of Solid-State Reactions:

Experiments, Journal of Physical Chemistry A, vol. 104, No. 46, pp. 10777-10782 Henderson, D. W. (1979), Experimental analysis of non-isothermal transformations

Kissinger, H. E. (1957), Reaction Kinetics in Differential Thermal Analysis, Analytical

Klement, W.; Willens, R. H & Duwez, P.O.L. (1960), Non-crystalline Structure in Solidified

Lass, E. A.; Zhu, A.; Shiflet, G. J. & Poon, S. J. (2010), A short-range ordering description of

Lee, Y.F. & Dollimore, D. (1998), The identification of the reaction mechanism in rising

Libermann H. & Graham C. (1976), Production Of Amorphous Alloy Ribbons And Effects

Málek, J.; Criado, J. M.; Šesták, J. & Militky, J. (1989), The boundary conditions for kinetic

Málek, J. (1992), The kinetic analysis of non-isothermal data, Thermochimica Acta, vol. 200,

Málek, J. (1995), The applicability of Johnson-Mehl-Avrami model in the thermal analysis of the crystallization kinetics of glasses, Thermochimica Acta, vol. 267, pp. 61-73 Málek, J. (2000), Kinetic analysis of crystallization processes in amorphous materials,

Maričić, A.; Minić, D. M.; Blagojević, V. A.; Kalezić-Glišović, A.; Minić, D. M. (2012), *Effects of* 

Matusita, K. & Sakka, S. (1979), Kinetic study of the crystallisation of glass by differential scanning calorimetry, Physics and Chemistry of Glasses, vol. 20, pp. 81-85 Matusita, K. & Sakka, S. (1980), Kinetic study of crystallization of glass by differential

Matusita, K.; Konatsu, T. & Yokota, R. (1984), Kinetics of non-isothermal crystallization

Minić, D. M.; Maričić, A., Dimitrijevic, R. Z.; Ristić, M. M. (2007), Journal of Alloys and

*structural relaxation on functional properties of amorphous alloy Fe73.5Cu1Nb3Si15.5B7*,

thermal analysis—criterion on application of Kissinger plot, Journal of Non-

process and activation energy for crystal growth in amorphous materials, Journal

Gold-Silicon Alloys, Nature, vol. 187, No. 4740, pp. 869–870

models, Thermochimica Acta, vol. 153, pp. 429-432

Thermochimica Acta, vol. 355, No. 1-2, pp. 239-253

Crystalline Solids, vol. 38-39, No. 2, pp. 741-746

of Materials Science, vol. 19, No. 1, pp. 291-296

Compdounds, vol. 430, pp. 241-245

Thermochimica Acta, vol. 196, No. 2, pp. 255-265

Part C: Polymer Symposia, vol. 6, No. 1, pp. 183-195

vol. 15, No. 2, pp. 325-331

No. 16 pp. 5460-5470

vol. 12, No. 6, pp. 921

pp. 257-269

Acta, vol. 323, No. 1-2, pp. 75-81

Intermetallics vol. 21, pp. 45-49

Chemistry, vol. 29, No. 11, pp. 1702-1706

the onset and final temperatures in the kinetic analysis of TG curves,

thermogravimetry. Application to a phenolic plastic, Journal of Polymer Science

The Universality of Master Plots for Analyzing Isothermal and Nonisothermal

involving nucleation and growth, Journal of Thermal Analysis and Calorimetry,

amorphous metal alloys using the central atoms model, Acta Materialia, vol. 58,

temperature kinetic studies based on the shape of the DTG curve, Thermochimica

Of Apparatus Parameters On Ribbon Dimensions, IEEE Transactions on Magnetics,

The values of parameters *m* and *s* obtained at the different heating rates for the investigated crystallization process of α-Fe in Fe81B13Si4C2 amorphous alloy are given in Table 9. Based on the obtained values of parameters *m* and *s*, at the different heating rates, (Table 9), we asserted with high degree of reliability, that the nucleation process of α-Fe occurs within amorphous alloy with a constant rate, in three effective directions (three-dimensional growth) proceeding with constant nucleation rate.

## **3. Conclusion**

Metallic amorphous alloy are a class of materials which has seen dramatic developments in recent times with design of materials stable enough to allow bulk production. While the functional properties of amorphous alloys allow for many possible fields of application, they can also be used as precursors in preparation of nanocomposite materials composed of nanocrystals in the amorphous matrix. These nanocomposites often exhibit superior mechanical, electrical and magnetic properties to both purely amorphous and purely crystalline materials. Thermal treatment of amorphous alloys can allow for controlled crystallization, leading to formation of nanocomposite materials with targeted properties.

Fe81B13Si4C2 alloy undergoes multi-step crystallization process as a result of thermal treatment. The changes in microstructure cause changes in electrical, magnetic and mechanical properties of the alloy as the alloy structure changes from predominantly amorphous to crystal/amorphous nanocomposite to nanocrystalline composite of α-Fe and Fe2B phases.

## **4. References**


The values of parameters *m* and *s* obtained at the different heating rates for the investigated crystallization process of α-Fe in Fe81B13Si4C2 amorphous alloy are given in Table 9. Based on the obtained values of parameters *m* and *s*, at the different heating rates, (Table 9), we asserted with high degree of reliability, that the nucleation process of α-Fe occurs within amorphous alloy with a constant rate, in three effective directions (three-dimensional

Metallic amorphous alloy are a class of materials which has seen dramatic developments in recent times with design of materials stable enough to allow bulk production. While the functional properties of amorphous alloys allow for many possible fields of application, they can also be used as precursors in preparation of nanocomposite materials composed of nanocrystals in the amorphous matrix. These nanocomposites often exhibit superior mechanical, electrical and magnetic properties to both purely amorphous and purely crystalline materials. Thermal treatment of amorphous alloys can allow for controlled crystallization, leading to formation of nanocomposite materials with targeted properties. Fe81B13Si4C2 alloy undergoes multi-step crystallization process as a result of thermal treatment. The changes in microstructure cause changes in electrical, magnetic and mechanical properties of the alloy as the alloy structure changes from predominantly amorphous to crystal/amorphous nanocomposite to nanocrystalline composite of α-Fe and Fe2B phases.

Adnađević, B.; Janković B. & Minić D. M.; (2010), Kinetics of the apparent isothermal and

Akahira, T. & Sunose, T. (1971) Trans. Joint Convention of Four Electrical Institutes, paper no. 246 (1969), Research Report, Chiba Institute of Technology, vol. 16, pp. 22-31 Avrami, M. (1939), Kinetics of Phase Change. I General Theory, Journal of Chemical Physics,

Blagojević, V. A.; Minić, D. M.; Žak, T.; Minić, D. M. (2011), Influence of thermal treatment

Balberg, I. & Helman, J. S. (1978), Critical behavior of the resistivity in magnetic systems. II.

Böhnke, G.; Kaul, S. N.; Kettler, W. & Rosenberg, M. (1983), Critical behaviour of the

Criado, J. M.; Pérez-Maqueda, L. A.; Gotor, F. J.; Málek, J. & Koga, N. (2003), A unified

Journal of Thermal Analysis and Calorimetry, vol. 72, No. 3, pp. 901-906 Dollimore, D.; Evans, T.A.; Lee, Y.F. & Wilburn, F.W. (1991), Calculation of activation

alloy, Journal of Physical Chemistry of Solids, vol. 71, pp. 927-934

non-isothermal crystallization of the α-Fe phase within the amorphous Fe81B13Si4C2

on structure and microhardness of Fe75Ni2Si8B13C2 amorphous alloy, Intermetallics

Below Tc and in the presence of a magnetic field, Physical Review B, vol. 18, No. 1,

electrical resistivity in amorphous ferromagnetic alloys, Solid State

theory for the kinetic analysis of solid state reactions under any thermal pathway,

energy and pre-exponential factors from rising temperature data and the generation of TG and DTG curves from A and E values, Thermochimica Acta, vol.

growth) proceeding with constant nucleation rate.

vol. 7, No. 12, pp. 1103-1113

Communications, vol. 48, No. 9, pp. 743-746

vol. 19, pp. 1780-1785

188, No. 1, pp. 77-85

pp. 303-318

**3. Conclusion** 

**4. References** 


**10** 

*Russia* 

**Crystallization of Iron-Containing** 

 *Institute of Metallurgy of the Ural Branch of the Russian Academy of Sciences* 

The processing of the sulphide raw materials (ores, concentrates and mattes) of non-ferrous metallurgy is related to the formation of a large amount of iron containing slags. The initial product of the oxidation of sulphides in real commercial plants is an oxide–sulphide melt, in which decomposition under the action of fluxes is accompanied by matte and slag formation (Selivanov et al., 2009a). The fraction of oxygen in a sulphide melt and the fraction of sulphur in an oxide melt are each controlled by the contents of silicon dioxide and iron oxides in a slag and the contents of non-ferrous metals in a matte. According to modern concepts, the heterogeneity of slags is caused by mechanical matte, magnetite and spinel inclusions, where the spinel inclusions form during oxidation processes (Selivanov et al., 2000; Spira & Themelis, 1969; Tokeda et al., 1983; Vanyukov & Zaitsev, 1969, 1973). The cooling (i.e., the crystallization) of a slag leads to the formation of new oxide and sulphide phases within it. Information on the available forms of the useful components is important for the reduction of

A number of works are devoted to the study of the kinds of copper existing in slags. Major results are generalized in the monographs of (Ruddle, 1953; Vanyukov et al., 1988; Vanyukov & Zaitsev, 1969; 1973). Phase equilibria in the systems relevant to copper pyrometallurgy have been discussed mostly for molten states (Elliott, 1976; Kopylov, 2001; Yazawa, 1974). It is considered that the loss of non-ferrous metals through slags is caused by their oxide, sulphide and metal solubility. It was discovered that a part of copper is presented in the crystallized slag by matte mechanical inclusions (Vanyukov & Zaitsev, 1969; 1973). Data on the copper sulphide solubility in a slag was reported by (Mohapatra, 1994; Nagamori, 1974; Vanyukov et al., 1988; Vaysburd, 1996). There is no valid confirmation of the presence of individual copper oxide inclusions or copper silicates and ferrites in a slag. Information on the existence of other metals (Zn, Pb, As, etc.) in a slag needs to be specified more exactly in each separate case. The bulk of the zinc is transferred into the slag during the smelting of sulphide copper-zinc concentrates in the Vanyukov furnace for a rich matte and crude metal (Vanyukov et al., 1988). It is assumed herein that zinc is present in a slag in the form of an oxide. Some questions concerning the constituent phases of crystallization during the rapid cooling of a non-ferrous metallurgy slag are partially disclosed by (Cardona et al., 2011). However, no task-oriented studies devoted to

metal loss through a slag and for the selection of their re-extraction methods.

**1. Introduction** 

**Oxide-Sulphide Melts** 

Evgeniy Selivanov and Roza Gulyaeva

