Synthesis of nanocrystalline Mg0.6Cd0.4Fe2O4 ferrite by glycine-nitrate auto-combustion method and investigation of its microstructure and magnetic properties

Author’s Accepted Manuscript Synthesis of nanocrystalline Mg0.6Cd0.4Fe2O4 ferrite by glycine-nitrate auto-combustion method and investigation of its microstructure and magnetic properties Fatemeh Zamani, Amir Hossein Taghvaei www.elsevier.com/locate/ceri

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S0272-8842(17)31993-4 http://dx.doi.org/10.1016/j.ceramint.2017.09.060 CERI16231

To appear in: Ceramics International Received date: 31 July 2017 Revised date: 28 August 2017 Accepted date: 9 September 2017 Cite this article as: Fatemeh Zamani and Amir Hossein Taghvaei, Synthesis of nanocrystalline Mg0.6Cd0.4Fe2O4 ferrite by glycine-nitrate auto-combustion method and investigation of its microstructure and magnetic properties, Ceramics International, http://dx.doi.org/10.1016/j.ceramint.2017.09.060 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Synthesis of nanocrystalline Mg0.6Cd0.4Fe2O4 ferrite by glycine-nitrate auto-combustion method and investigation of its microstructure and magnetic properties Fatemeh Zamani, Amir Hossein Taghvaei1 Department of Materials Science and Engineering, Shiraz University of Technology, Shiraz, Iran

Abstract Nanocrystalline Mg0.6Cd0.4Fe2O4 ferrite powders were produced by the glycine-nitrate autocombustion method for the first time. The influence of the different molar ratios of glycineto-nitrate (G.N-1) on the characteristics of the prepared powders was systematically investigated by X-ray diffraction (XRD), inductively coupled plasma optical emission spectroscopy (ICP-OES), differential scanning calorimetry (DSC), Fourier transform infrared spectroscopy (FTIR), field-emission scanning electron microscopy (FE-SEM) and vibrating sample magnetometry (VSM). Thermodynamic calculations revealed that the adiabatic flame temperature changes from 598.79 K to 1757.97 K by increasing the G.N-1 ratios from 0.30 to 0.85. The results confirmed that under fuel-lean combustion (G.N-1 = 0.30), Mg0.6Cd0.4Fe2O4 nanoparticles can be obtained at a significantly lower temperature and shorter synthesis time, compared to other preparation methods like standard ceramic and co-precipitation. The XRD and ICP results showed that the crystallite size of the powders changes in the range of 8 nm to 43 nm, and their Cd content notably decreases with increasing the G.N-1 values. The FESEM results proved that the porosity and size of the as-prepared ferrite nanoparticles drastically change with variations in the G.N-1 ratio. The evolution of phase, crystallite/particle size, and magnetic properties after annealing was discussed in detail. At the optimized annealing condition, the synthesized Mg0.6Cd0.4Fe2O4 ferrite offered a high saturation magnetization of 41.70 Am2.kg-1 and a coercivity of 1.92 kA.m-1, indicating noticeably better soft magnetic properties compared to the same ferrite produced by the other wet chemical methods. Keywords: Nanocrystalline, Glycine-nitrate auto-combustion, Magnetic properties. 1. Introduction Ferrites, as a typical category of magnetic materials, have attracted attention in different fields, due to their promising and tunable electrical, optical and magnetic properties [1, 2]. These materials can be used in a wide range of applications like telecommunications, 1

Corresponding author Tel : +98 917 7038948; E-mail addresses: [email protected] 1

electronics, microwave devices, magnetic drug delivery, ferrofluids, high-density storage, gas and humidity sensors, catalysts and adsorbent materials [3-12]. It is widely accepted that the physical properties of the ferrites strongly depend on many parameters, such as their cationic distribution, particle size, crystallites dimensions and their purity, which are significantly correlated with the synthesis method of the ferrites. Nowadays, diverse wet-chemical routs like co-precipitation, sol-gel, combustion synthesis, reverse micelle and hydrothermal are extensively used in order to prepare the different types of nanoparticles [13-15], including ferrites [16, 17]. Such process can offer numerous advantages compared to the conventional solid-state standard method like a lower synthesis temperature, faster production rate, higher purity/homogeneity of the products, and finer particle size [18]. On the other hand, the aforementioned processes may exhibit some limitations, such as a low efficiency of production and the existence of various parameters determining the properties of the final products [3]. For example, despite the fact that co-precipitation is a simple and suitable process for the mass production of ferrite nanoparticles, it requires a precise adjustment of parameters like pH, concentration, temperature and stirring speed [19, 20]. In addition, a long synthesis time, needing careful drying and calcination steps, and in some cases, high cost of precursor, like metal alkoxides are the main limitations of the sol-gel process [21]. Recently, the sol-gel auto-combustion process has been extensively employed for the production of the different types of ferrites [3, 21, 22]. This is a simple, cost effective, low external energy consuming and fast method, where chemical sol-gel and combustion process are combined [3]. Indeed, a homogeneous mixing of chelated cations in a molecular scale, followed by an exothermic redox reaction in the prepared xerogel can yield ultrafine or nanosized ferrite powders with a high-degree of crystallinity and low impurity content [9, 23]. It is well-known that the microstructure, composition, morphology, surface area and particle size of the final product are attributed to the flame temperature and rate of combustion, that can be controlled by the type of fuel (complexant), fuel-to-oxidizer (nitrates) ratio, and concentration of the additional oxidizer [3]. Among the mentioned parameters, the fuel-to-oxidizer ratio as a simply controllable parameter can play a leading role in the final properties of the prepared ferrite [24].

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Mg ferrites are an important class of soft magnetic ferrites that represent very low magnetocrystalline anisotropy, low coercivity, moderate saturation magnetization, low eddy current loss and high electrical resistivity, which attracts a great deal of attention for high frequency applications, such as power transformers and telecommunication [9, 10, 25]. In addition, Mg ferrite nanoparticles can exhibit a superparamagnetic behavior at room temperature, which is beneficial for biomedical applications [26]. It has been shown that the partial replacement of Mg2+ with Cd2+ or Zn2+ can improve the magnetic properties of Mg ferrite, especially its saturation magnetization [9, 10, 18]. Mg-Cd ferrites with inhomogeneous and coarse particle size (micron range) were synthesized by the standard ceramic method [27]. It was shown that the complete reaction between the constituents and formation of Mg-Cd ferrites needs a long annealing time at high temperature (1273 K, 48 h) [27]. Gadkari et al., produced Mg-Cd ferrite with ultrafine particles size by oxalate coprecipitation method [10]. They showed that the ferrite formation occurs at a high temperature of 1323 K after 5 h annealing [10]. In addition, it was indicted that the oxalate co-precipitation method is a long process, that needs several steps including pre-sintering at 973 K for 6 h and subsequent ball milling, before final heat treatment [10]. Furthermore, synthesis of the Mg-Cd ferrites was also carried out by combustion method using Oxalyl dihydrazide fuel [8]. It was shown that ferrite formation requires post annealing treatment at 573 K for 3 h, after combustion. On the other hand, the reported soft magnetic properties was not promising, since the maximum saturation magnetization was measured as 23.5 Am2.kg-1 for Mg0.6Cd0.4Fe2O4 ferrite, which exhibited a very large coercivity of 13.34 kA.m-1 [8]. To the best of our knowledge, synthesis of the Mg-Cd ferrite by sol-gel auto-combustion process using other fuels like glycine has not been reported. In the present work, the Mg0.6Cd0.4Fe2O4 ferrite was produced for the first time by sol-gel auto-combustion method using glycine as fuel and metal nitrates as oxidizer, which is also called glycine-nitrate combustion process. Glycine was selected due to its low cost, good availability, high chelating power of the metal cations [28], and its large heat of combustion (-13.56 kJ.g-1), which is more than urea (-12.47 kJ.g-1) or citric acid (-11.55 kJ.g-1) [29, 30]. The influence of different glycine-to-nitrate ratio on combustion temperature, composition, microstructure, morphology and magnetic properties of the synthesized ferrites were systematically investigated. It was shown that the production of Mg-Cd ferrite is much faster and it forms at a significantly lower temperature compared to the standard ceramic and co-precipitation methods. In addition, the results demonstrated that the glycine-nitrate method can produce 3

Mg-Cd ferrite with better soft magnetic properties, compared to other wet chemical methods, used for synthesis of this ferrite. 2. Experimental procedure Nanostructured Mg0.6Cd0.4Fe2O4 ferrite was produced by sol-gel auto-combustion using glycine as a fuel. For the synthesis of 2 g ferrite, Fe(NO3)3.9H2O (>99%, Merck, Germany), Mg(NO3)2.6H2O (>99%, Merck, Germany) and Cd(NO3)2.4H2O (>99%, Sigma-Aldrich, United States) with the stoichiometry of Mg0.6Cd0.4Fe2O4 ferrite were weighed and then dissolved in 25 ml of distilled water. Then, the various molar ratios of glycine (NH2CH2COOH, >99%, Merck, Germany) was added to the solution obtained and the resulting solution mixed at room temperature on a magnetic stirrer for 1 h. Finally, the combustion synthesis was performed by heating of the produced sol on a hot plate until the prepared viscous gel was converted to a porous and fluffy product with a dark-brown color, which was subsequently milled gently in a mortar for 5 min to obtain the powders. The phase studies and calculation of main structural parameters were conducted by X-ray diffraction (XRD: Bruker X-ray powder diffractometer) with a step size of 0.03 degree and step time of 1 s using CuKα radiation. The composition of the prepared powders was checked by inductively coupled plasma optical emission spectroscopy (ICP-OES: Model 730-ES, Varian). For this purpose, each powder was digested in a concentrated nitric acid solution at room temperature. The obtained solutions were diluted to 100 ml by deionized water before ICP measurements. The morphology and size distribution of produced particles were studied by field-emission scanning electron microscopy (FE-SEM: Model Mira3-XMU, Tescan). Fourier transform infrared spectroscopy analysis (FTIR: IRAffinity-1 Shimadzu) was performed to investigate the formation of chemical bonds. Thermo gravimetric behavior of the synthesized gel and the nature of combustion reaction were studied by a differential scanning calorimeter (DSC: Model c1600, Mettler Toledo) under flow of air, with a constant heating rate of 10 K.min-1. The magnetic properties of the powders were measured at ambient temperature by a vibrating sample magnetometer (VSM: Meghnatis Daghigh Kavir Co, Iran) under a maximum applied field of 800 kA.m-1.

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3. Results and discussion 3.1. Thermodynamic studies and the nature of combustion reaction It is highly accepted that the nature of combustion synthesis and the characteristics of the produced powders significantly depend on the G.N-1 ratio. The combustion phenomena, which is an exothermic reaction, proceeds through a self-progressive manner, by reaction between oxidizing (metal nitrates) and reducing reactants (glycine) [5]. It was also observed that even the color of synthesized powder is different for various G.N-1 ratios, and the color is changed from light-brown to dark-brown with increasing G.N-1 from 0.30 to 0.85. The color differences are correlated probably to the remaining extra carbon, resulting from the inappropriate oxidizing [9]. The combustion was associated with the release of a large volume of gases, changing based on the G.N-1 ratio as[20]: 0.6Mg(NO3)2.6H2O+0.4Cd(NO3)2.4H2O+2Fe(NO3)3.9H2O+φNH2CH2COOH+(2.25φ-10)O2 (g) → Mg0.6Cd0.4Fe2O4+2φCO2 (g)+(24+2.4φ)H2O (g)+(4+0.5φ)N2 (g)

(1)

The principle of propellant chemistry reveals that when the ratio between the total valences of metal salts and those of complexant agent (fuel) is unity, the combustion occurs through a stoichiometric redox reaction [31].

According to above relation, different contents of

glycine, φ, as fuel-lean values (2.4 and 3.36), stoichiometric amount (4.44) and fuel-rich values (5.6 and 6.8) were selected, which are equivalent to G.N-1 ratio of 0.30, 0.42, 0.55, 0.70 and 0.85, respectively. As can be observed, with increasing G.N-1 ratio, the amount of produced gas enhances. In the case that total produced heat is spent for heating up the product and there is no heat exchange with surrounding, the theoretical adiabatic flame temperature (Tad) can be correlated with reaction enthalpy as [5, 20]: ∫

∑

(2)

where T0 is 298 K, n denotes the number of moles, cp is the heat capacity of products at constant temperature and ∑ where

is the enthalpy reaction that expressed as follows [5, 20]: ∑

(3)

is the enthalpy of formation. By using the thermodynamic data for the reactants

and products (Table 1), the combustion enthalpy and Tad can be calculated as a function of 5

the G.N-1 ratio, as shown in Table 2. As can be seen, the number of moles of gases, and both and Tad increase by increasing the amount of used fuel for combustion. It is worth noting that the actual temperatures of flame is usually lower than the theoretical calculated values, because of incomplete combustion or consuming part of the evolved or dissipation of heat through the radiation [9]. At G.N-1 ratios of 0.30 and 0.42, which the amount of nitrate ions is more than the reductant (fuel), the released excess oxygen should be heated to the temperature of the products [21]. Hence, sufficient temperature rise is not obtained and the rate of the combustion reaction is slow. It should be noted that no intense flame was observed for these fuel-lean condition in our experiment. In contrast, at stoichiometric condition (0.55), the most intense flame was detected, demonstrating a complete combustion. For the fuel-rich condition, e.g, G.N-1 ratios of 0.70 and 0.85, the combustion reaction needs to supply external oxygen, thus, the reaction rate decreases and the required time to reach the combustion temperature and subsequent ignition increases. Generally, it can be said that at fuel-rich or fuel-deficient conditions, the reaction rate is supressed due to the lack of the oxidizing or reducing agents [3]. 3.2. Structural analysis To investigate the formation of the desired phase and structural studies, XRD analysis was performed on the synthesized powders and the measured patterns were evaluated with the X'Pert HighScore software (Fig. 1). Table 3 lists the crystallite size and lattice parameter of the synthesized powders. The crystallite size of the powders was measured by the DebyeScherer`s method [32]. Fig. 1 clearly shows that the content of glycine has a notable effect on the phase evolution and the patterns intensity. As the figure shows, except for the powder at produced at G.N-1 ratio of 0.30, which exhibits no sharp diffraction peaks, all samples indicate formation of the spinel structure with high degree of crystallinity. Furthermore, it can be inferred that the G.N1

ratio affects the formation of an impurity phase like CdO, besides the main spinel

compound, which also was observed in other works [9]. In addition, the intensity of the patterns initially increases and reaches a maximum value at stoichiometric (G.N-1 of 0.55) ratio, and subsequently decreases slightly with using more fuel. According to Fig. 1, the intensity of CdO compound declines with increasing the fuel content, and the obtained powders prepared at G.N-1 ratio above 0.55 show only the diffraction peaks of cubic spinel structure. This result demonstrates that the G.N-1 ratio in the 6

sol-gel auto-combustion process plays a leading role in final structure and composition of the synthesized powders. It should be considered that glycine not only acts as a fuel or reducing agent, but it can effectively chelate different cations, due to the presence of carboxylic acid at one side and amine group at its other side. As a result, glycine, as an important complexant molecule prevents the selective deposition of cations and helps to reach a homogeneous composition [3]. According to Fig. 1 and Table 3, the first powders synthesized at the lowest G.N-1 ratio (0.30) exhibit the smallest crystallite size of about 8 nm and the largest lattice parameter of 0.8477 nm. Such a small crystallite size is correlated with a low combustion temperature (see Table 2) that minimizes the crystallite growth. It is worth to note that the formation of CdO compound at G.N-1 = 0.30 is attributed probably to insufficient amount of glycine for chelating all cations in the prepared sol. On the other hand, a low combustion temperature at this condition may avoid a complete reaction among the cations, and consequently prevent the incorporation of all Cd2+ with a large ionic radius of 0.097 nm [8] into the spinel structure, resulting in formation of CdO. It should be noted that the lattice constant of the obtained ferrite at this G.N-1 (0.8477 nm) is slightly lower than that of Mg0.6Cd0.4Fe2O4 bulk ferrite (0.8512 nm) obtained by standard ceramic method [27], due to a lower content of Cd2+ cations. According to Fig. 1, by incresing the G.N-1 ratio to 0.42, both crystallite size and diffraction intensities increase due to the enhanced combustion enthalpy and the corresponding reaction temperature. Further complexant addition up to the stoichiometric ratio (0.55), that is in accordance with the complete combustion is accompanied by the larger crystallite sizes (43 nm) and peaks intensities. On the other hand, the XRD patterns show the decreased fraction of CdO by increase G.N-1 ratio. According to Table 3, the lattice constants of the powders synthesized at G.N-1 above 0.30, particularly near the stoichiometric ratio are notably smaller than that of the ferrite prepared at G.N-1 = 0.30, implying the substantial changes in the stoichiometry of the prepared ferrites, as can be confirmed by the ICP analysis. Table 4 lists the molar ratio between the constituent elements of the prepared ferrites at different G.N-1 ratio, measured by the ICP results. The obtained molar ratios determined at G.N-1 = 0.30 are in good agreement with nominal composition (Mg0.6Cd0.4Fe2O4). In contrast, a significant cadmium loss is observed for other cases, which is responsible for the observed lattice contraction for powders produced at large G.N-1 ratios. The cadmium loss is probably

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attributed to a noticeable temperature rise during combustion at large G.N-1 ratios (Table 2), leading to evaporation of CdO, which has a high vapor pressure and low sublimation temperature (1832 K). Furthermore, a vigorous gas evolution upon combustion, particularly for larger G.N-1 ratios can have a synergic effect on cadmium loss from the precursor [33]. A similar result was observed during synthesis of Mn-Zn ferrites, where the temperature rise upon combustion led to a notable zinc loss, despite the fact that ZnO has even larger sublimation temperature (2248 K) compared to CdO [33]. According to Table 3, the crystallite size of the powders is reduced in fuel-rich condition, due to the enhanced volume of gases increasing the cooling rate of products, and consequently impedes the growth of crystallites. Moreover, a finer crystallite size is not correlated only with the volume of evolved gases, but it depends on the chemical composition of the prepared ferrites. The above results demonstrate that while the ferrite produced at G.N-1 = 0.55 exhibits the improved crystallinity, however, the composition is only preserved in the fuel-deficient condition. As a result, G.N-1 = 0.30 is a proper value for a low temperature preparation of Mg-Cd ferrite without any significant deviation in the composition. In addition, compared to the co-precipitation process, synthesis of the Mg0.6Cd0.4Fe2O4 ferrite using glycine-nitrate combustion is much faster, since it doesn’t need two heat treatment (973 K, 6 h and 1323 K, 5 h) and ball milling steps, which can introduce contamination and impurities [10]. Moreover, as mentioned before, formation of the Mg0.6Cd0.4Fe2O4 ferrite in this work occurs at a significantly lower temperature compared to the co-precipitation synthesis. The formation temperature of the Mg0.6Cd0.4Fe2O4 ferrite prepared by oxalyl dihydrazide-nitrate method was reported as 573 K, which is still larger than one prepared with glycine-nitrate process. Indeed, the mixing of chelated cations by the complexant in the molecular-level, as well as the exothermic reaction between the glycine and nitrates can notably decrease the time and temperature of the synthesis [3].

3.3. Thermal analysis The combustion reaction between the metal nitrates and glycine, and its subsequent transformation to ferrites powders were studied by the DSC-TG analysis. For this purpose, the synthesized gels at three G.N-1 ratios of 0.30, 0.55 and 0.85 were heated in air up to 1273 K, with a heating rate of 10 K.min-1 (Fig. 2). 8

According to Fig. 2, a weight loss of about 10% corresponding to endothermic event in the temperature range of 333 K to 423 K in all three G.N-1 ratios is observed, which is attributed to evaporation of the residual water in the gel structure [5, 19, 20]. Further increase of temperature is followed by a sharp exothermic peak in three plots, with an onset temperature of about 443 K, which is accompanied by a sharp weight loss of 70 to 75%. This exothermic event signifies that that combustion reaction is occurred very rapidly through a redox reaction among the metal nitrates and glycine, according to the reaction (1), with overall weight loss of about 80 to 85%. The onset temperature of combustion (443 K), which is comparable to that of other ferrite gels prepared by glycine-nitrate method, is lower than the ignition temperature of the synthesized precursor using other fuels like citric acid or urea, due to larger combustion heat of glycine [21]. According to Fig. 2 (a), a small weight loss of about 5 % occurs in a temperature range of 473 K to 673 K, which can be due to decomposition of residual organic material, not completely decomposed due to incomplete combustion at fuel-deficient condition. Further temperature increase is accompanied with a broad exothermic event in the temperature range of 773 K to 923 K, without any mass change, which is probably associated with the crystallite growth. As mentioned before, this sample has the lowest crystallite size of about 8 nm (Table 3), which is favorable for subsequent growth. It can be seen that an endothermic reaction occurs by further temperature rise in the range of 930 K to 1093 K, without any detectable mass variation. The nature of this reaction can be determined through the XRD analysis (Fig. 3) of the as-prepared powders after annealing (2 h) in the temperature range of the observed endothermic event. According to Fig. 3, the annealing significantly increases the intensity of the XRD patterns, due to the enhanced crystallinity and grain growth (see Table 3). On the other hand, the heat treatment noticeably decreases the intensity of CdO diffraction peaks and leads to formation of a single phase Mg0.6Cd0.4Fe2O4 ferrite with large crystallinity (see Fig. 3) Hence, the observed high-temperature endothermic event in the DSC plot (Fig. 2(a)) is attributed to incorporation of remained Cd2+ cations into the spinel lattice. This result is in agreement with slight increase of lattice constant, according to Table 3, after annealing, compared to that of as-prepared ferrite. According to Fig. 2(b), exothermic event, corresponding to the grain growth cannot be clearly detected for the powders synthesized at G.N-1 = 0.55, due to a larger crystallite size in this

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case (see Table 3), which exhibits a lower tendency for growth. Instead, the DSC plot exhibits an exothermic hump at higher temperature around 973 K, without a notable mass change. The XRD result of this powder after annealing at a temperature above the mentioned reaction (T = 1073 K) indicates the Bragg peaks of hematite (α-Fe2O3), besides those of spinel phase (Fig. 4), which was not observed in the XRD pattern of the as-prepared powder (Fig.1). Hence, the appeared exothermic reaction (Fig. 2(b)) indicates hematite formation at high temperature. Table 3 shows that for G.N-1 = 0.55, the lattice constant of spinel phase after annealing (a = 0.8388 nm) is comparable to that of as-prepared sample (a = 0.8366 nm), and both are notably lower than that of single phase Mg0.6Cd0.4Fe2O4 ferrite (a = 0.8512 nm), due to significant cadmium loss, discussed before. Based on Fig. 2(c), the TG curve of the samples prepared at fuel-rich condition (G.N-1 = 0.85) shows still a weight loss of 16 % after combustion reaction, up to 873 K, which is accompanied by an exothermic event in the DSC plot. Such effects are in line with decomposition of excess glycine used or burning of the remained carbonaceous material. Similar to previous case, annealing above the DSC peaks is followed by formation of Hematite (see Fig. 4), even at lower temperature (873 K vs 1073 K) Such formation of α-Fe2O3 upon heat treatment probably correlated with a large deviation in the stoichiometry of as-synthesized ferrites, confirmed by ICP analysis (Table 4). Indeed, the existence of excessive Fe3+ relative to other cations, due to Cd loss may result in formation of some iron oxides, such as Maghemite (γ-Fe2O3) or Magnetite (Fe3O4) upon combustion, which can transform to a more stable Hematite phase after annealing. As shown in Fig. 4, the annealed powders represent also the diffraction peaks of spinel phase, which is probably attributed to Cd- deficient Mg-Cd ferrite. In fact, due to the same crystal structure (spinel) and close lattice constants values of Maghemite (0.835 nm[34]), Magnetite (0.839 nm [34]) and Mg ferrite (0.837 nm[4]), it is very difficult to separate their diffraction peaks in the XRD results of the as-synthesized powders (Fig. 1). The evaporation of some oxides with high vapor pressure and low sublimation temperature like ZnO and CdO upon combustion or high temperature annealing (usually above 1273 K) of ferrites can result in a notable compositional deviation due to Zn or Cd loss. It has been reported that such volatilization leads to the excess unsaturated oxygen ions, which can reduce Fe3+ ions to Fe2+ and the same time, loss of some oxygen for preserving the charge neutrality [35, 36]. In this case, if a significant evaporation takes place, the amount of Fe2+

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ions increases and formation of Fe3O4 is expected. Upon temperature rise, Fe3O4 transforms initially to γ-Fe2O3 and then to α-Fe2O3. The former transition occurs usually in the temperature range of 423-573 K and the later can happen from 473 K even up to 1023 K, depending on chemical composition (existence of other cations besides Fe3+ ) and initial particle size of the sample [34, 37]. On the other hand, it was shown that γ-Fe2O3 is formed in the form of solid solution with stoichiometric Cd-Ni ferrite, due to drastic Cd loss observed upon annealing of Cd-Ni ferrites [38]. Thus, the detected high temperature exothermic events in Figs. 2(b) and (c) can be assigned to γ-Fe2O3 - α-Fe2O3 phase transformation.

3.4. FTIR analysis The synthesized gel and ferrite powders prepared at different G.N-1 content were further characterized using Infrared spectroscopy. Fig. 5 shows the FTIR plot in the frequency range of 400-3000 cm-1 of dried gel, obtained at G.N-1 = 0.55. As the figure shows, visible absorption bands at 800 to 850 cm-1 and 1350 to 1400 cm-1 are related to

ions, implying

that the oxidant ions are exist in the structure of dried gel [19, 29]. The detected bands around 825 cm-1 and 690 cm-1 are assigned to deformation mode of C-H group. The band observed at 591 cm-1 can imply the formation iron hydroxide (Fe(OH)2+) [29]. The two bands indicated at 1040 cm-1 and 1115 cm-1 are attributed to rocking vibration of NH2 group of glycine [29]. Moreover, the appeared band at 1640 cm-1 originates from the deformation vibration of NH2 group in glycine, which locates slightly at a higher frequency, compared to that of pure glycine (not shown here). This effect is probably caused by cations chelation in the dried gel through the amino group of glycine and formation of the monodentate ligand [29]. Furthermore, the absorption band near 1750 cm-1 shows the stretching vibration of carbonyl group in the carboxylic (COOH) functional group of glycine [29, 39]. It has been shown that the coordinated cations by carboxylic group usually form bands between 1357-1390 cm-1 and 1625 cm-1, corresponding to the asymmetric and symmetric COO- stretching vibrations, respectively [29, 40]. According to Fig. 5, both of these bands appear as weak shoulders, which suggests that the cations are mainly coordinated around amino group of glycine through the formation of monodentate ligand. Fig. 6 shows the FTIR spectra of ferrite powders prepared at different G.N-1 values. According to Fig. 6, appearance of two absorption bands between 400-600 cm-1, are corresponding to the metal-oxide stretching vibrations, characteristic of materials with spinel 11

structure [4, 10, 22]. The first band near 400 cm-1 indicates the stretching vibration of cations in the octahedral site, while the second band around 550 cm-1 describes the stretching vibration of cations in the tetrahedral site [4, 10, 22]. Formation of these two bands in all synthesized powders reveals that all of them have the spinel structure. According to Fig. 6, all samples show a shoulder after 550 cm-1, due to presence of different cations in the tetrahedral site. As can be seen, the intensities of above bands are lower for ferrite produced at G.N-1 = 0.30, due to its lower crystallinity. In addition, very weak absorptions, assigned to nitrates and residual organic groups are still present for this sample due to a lower temperature of combustion. The absence of such absorption peaks for other samples manifests a higher temperature attained upon combustion at larger G.N-1, which completely eliminates nitrate, carboxylic and amino groups. Based on Fig. 6, the intensity of characteristic bands of Mg0.6Cd0.4Fe2O4 ferrite significantly levels up after annealing at 1173 K, due to the increased crystallinity (see Fig. 3), while the bands corresponding to the residual organic materials are almost vanished. 3.5. Microstructural analysis To study the morphology and distribution of the particles produced at different G.N-1 values, FE-SEM experiment was conducted. According to Fig. 7, the images reveal some changes in porosity of powders with increase of G.N-1 ratio. It is inferred that all samples exhibit a spongy or foam-like structure due to release of gases evolved upon combustion reaction. At low combustion temperature (G.N-1 = 0.3), formation of a dense structure is prevented and owing to a lower content of generated gaseous products, the microstructure contains dominantly small pores. By increasing G.N-1 ratio, larger porosities are observed, resulting from the increased volume of generated gases (Fig. 7 (e)). The detected agglomerates in all three samples demonstrates that they include ultrafine or nanoparticles (Fig. 7 (b), (d) and (f)). Based on FE-SEM results, it is clear that the G.N-1 ratio considerably affects the size and distributions of the particles (Fig. 9), determined by image analysis, using MIP software. According to Figs. 7 and 9, it is concluded that the as-synthesized powders mainly contains nanoparticles with a relatively wide particle size distribution. In addition, according to Fig. 9, it is obvious that the size of particles becomes initially larger with increase of G.N-1 up to 0.55 and then, their size diminishes for fuel-rich condition. Such an observed trend, like crystallite size evolution (Table 3) is caused by different combustion temperature, and volumes of the escaping gases. The initial growth of the particles originates from the enhanced combustion temperature (Table 2), while the subsequent particle size reduction is 12

caused by a higher cooling rate, assisted by huge volume of evolved gases occurs in the fuelrich condition. Fig. 8 shows the FE-SEM results of the Mg0.6Cd0.4Fe2O4 ferrite produced at G.N-1 = 0.30, followed by heat treatment at 1173 K. Compared to Fig. 7 (a) and (b), the annealing notably reduces the size and content of the pores generated during combustion. In addition, the obtained Mg0.6Cd0.4Fe2O4 ferrite consists of ultrafine grains with a wide size distribution (Fig. 9 (d)), which experienced a considerable growth during annealing.

3.6. Magnetic properties Fig. 10 shows the M-H hysteresis loops of the powders prepared at different G.N-1 values. According to this figure, significant changes in the shape of loops and corresponding magnetic parameters are observed (Table 3). As can be seen, the saturation magnetization, Ms, of the powders initially levels up with increase of G.N-1 ratio up to 0.55 (stoichiometric ratio), and then declines with further addition of fuel content. The powders produced at G.N-1 = 0.30 exhibit a sigmoidal-shape hysteresis plot, showing a linear increase of magnetization with applied field up to 800 kA.m-1. The absence of saturation may result from the coexistence of some superparamagnetic particles in this sample that represents the smallest particle size (Fig. 9 (a)). In addition, the minimum values of coercivity, Hc, and remanence, Mr, measured for above sample can further support the presence of superparamagnetic nanoparticles in that. It is well-known that the particle size reduction in the nanometer regime can noticeably suppress the magnetization due to spin disorder on or near the particles surface [41]. Indeed, the broken exchange bonds on the surface of nanoparticles destroys the ferrimagnetically-aligned spins exist in the core of ferrites nanoparticles, and leads to formation of canted spins on their surface, which is an important source of Ms reduction [41]. Formation of such canted/disordered spin structure may prevent the saturation of magnetization of ferrite nanoparticles produced at G.N-1 = 0.30 (Fig. 10). The particles growth and improved crystallinity favored by increased G.N-1 contributes to an initial Ms enhancement, according to Table 3. In addition, the disappearance of non-magnetic CdO phase and possibility of Fe3O4/ γ-Fe2O3 formation for larger G.N-1 values can give rise to Ms increase. According to Table 3, the Ms decline observed in fuel-rich condition is mainly attributed to the particle size reduction, discussed before. It is worth mentioning that drastic changes in particles size and compositional deviations caused by Cd loss during synthesis, 13

can affect the cationic distributions, and consequently the magnetic properties of the synthesized ferrites According to Table 3, the G.N-1 ratio dramatically affects the Hc of the prepared powders. As the table shows, the Mg0.6Cd0.4Fe2O4 ferrite produced at G.N-1 =0.3 exhibit the lowest coercivity of 1.56 kA.m-1, indicating its soft magnetic behavior. It can be seen that the Hc doesn’t show a specific trend with increasing G.N-1 ratio. It is well-known that Hc, as a nonintrinsic magnetic property strongly depends on microstructural characteristics, such as grain/particle size, particles shape and other factors like impurity, surface roughness and demagnetizing field [32]. Fig. 11 depicts the evolution of Hc as a function of crystallite size (Table 3) of the synthesized powders. It can be found that Hc exhibits firstly an ascending trend and subsequently a decreasing tendency with crystallite size. The initial magnetic hardening can be a consequence of enhanced average magnetocrystalline anisotropy by increasing the crystallite size up to 38 nm. According to the random anisotropy model, when the crystallite size is smaller than magnetic exchange length, the crystallite growth can increase the Hc as [42]: (4)

where K1 is the crystal anisotropy constant and D is the crystallite size. From Fig. 11, it is concluded that the crystallite size below 38 nm are probably less than the magnetic exchange length of the synthesized ferrites. However, it should be taken into account that Ms and K1 are also important parameters in determining Hc values (Eq. (4)), and both can change with G.N1

, due to variations in chemical composition (Cd loss) and particle size. However, it seems

that the crystallite size is the most determinant factor due to its larger exponent (6), compared to Ms and K1 (Eq. (4)). It can be seen that the increase in crystallite size above 38 nm is followed by magnetic softening, according to Fig. 11. For crystallite size larger than magnetic exchange length, each crystallite can contain several domains, separated by domain walls. In this case, the grain boundaries can play a leading role in pinning of domain walls and then increasing the Hc. Hence, in this condition, crystallite growth and the reduced fraction of grain boundaries can favor magnetic softening.

14

The above results indicate that the Hc of the prepared powders with G.N-1 = 0.85, which has the crystallite size of 38 nm is maximum. Recently, Hajarpour et al., have found that the Mg0.6Zn0.4Fe2O4 powders synthesized by glycine-nitrate combustion exhibit the maximum Hc at crystallite size of about 32 nm, which is comparable to the current results [9]. It is suggested that the magnetic exchange length for most of magnetic particles lies between 2050 nm, which is in good agreement with our results [7, 9]. Fig. 12 shows the M-H hysteresis loops of prepared ferrite powders after heat treatment. The annealing temperatures were selected according to the DSC results, discussed before. Table 3 lists the magnetic parameters obtained from Fig. 12. The measurements reveal that for the Mg0.6Cd0.4Fe2O4 ferrite produced at G.N-1 = 0.30, Ms represents an increasing tendency with annealing temperature and it reaches a maximum value of 41.70 Am2.kg-1 after heat treatment at 1173 K. The saturation magnetic moment per formula unit, nB, (in µB) is calculated as [10]:

nB 

M wM s 5585

(5)

where Mw is the ferrite molecular weight. According to Table 3, nB is determined only for the as- prepared and annealed ferrites synthesized at G.N-1 = 0.3, due to significant Cd loss, and consequently unknown Mw values for other samples. The improve of Ms and nB values upon annealing can be attributed to the decreased contribution of the magnetic dead layers and disordered surface spins, as a result of enhanced crystallinity and particles size increase (see Fig. 8). Less canted spins for larger particles demonstrates that the spins are mostly collinear and the ferrite magnetization is governed by the Neel`s two sublattice model. Besides magnetization, the annealing affects the Hc of the Mg0.6Cd0.4Fe2O4 powders, as can be seen in Table 3. The evolution of Hc after annealing is mainly correlated with variations in the crystallite size (see Table 3) in the ranges below the magnetic exchange length (sample annealed at 973 K), or notably above it (samples annealed at 1123 and 1173 K). It is worth to note that the obtained Mg0.6Cd0.4Fe2O4 ferrite powders after annealing at 1173 K demonstrates a higher Ms and significantly lower Hc compared to the same composition produced by co-precipitation method (Ms = 30.44 Am2.kg-1, Hc = 7.79 kA.m-1) or combustion synthesis using oxalyl dihydrazide fuel (Ms = 23.5 Am2.kg-1, Hc = 13.34 kA.m-1), signifying considerably better soft magnetic properties [8, 10]. Interestingly, even the as-synthesized Mg0.6Cd0.4Fe2O4 ferrite in this study displays a markedly lower Hc (1.56 kA.m-1) with respect to the same ferrite synthesized by other methods, mentioned above. A superior magnetic

15

behavior can be a consequence of better crystallinity or higher homogeneity and purity of the ferrites produced by glycine-nitrate method. Fig. 12 and Table 3 reveal that the powders produced at G.N-1 = 0.55 and 0.85 exhibit drastically lower Ms and larger Hc values after heat treatment, manifesting their inferior soft magnetic behavior. According to Fig. 4, annealing of these samples leads to formation of non-magnetic α-Fe2O3 phase that can reduce the Ms and increase their Hc.

4. Conclusion In this study, nanostructured Mg0.6Cd0.4Fe2O4 ferrite was produced by the glycine-nitrate auto-combustion method and it was found that the molar ratio of glycine- to- nitrate (G.N-1) has a significant impact on the microstructure, morphology, particle size and magnetic properties of the obtained powders. It was shown that sol-gel autocombustion with glycine as a fuel is a simple, low-cost and rapid method for synthesis of the Mg0.6Cd0.4Fe2O4 ferrite nanoparticles at a significantly lower temperature compared to standard ceramic and coprecipitation methods. The results revealed that the Mg0.6Cd0.4Fe2O4 ferrite can be obtained only at fuel-lean condition (G.N-1 = 0.3) and larger glycine content resulted in significant Cd loss during combustion. The powders synthesized at G.N-1 = 0.3 represented a single phase Mg0.6Cd0.4Fe2O4 ferrite with a high- degree of crystallinity upon subsequent heat treatment up to 1173 K. On the other hand, those powders produced at larger fuel content showed a noticeable formation of α-Fe2O3 after annealing. The produced Mg0.6Cd0.4Fe2O4 ferrite showed promising soft magnetic properties, i.e., a notably larger saturation magnetization and a lower coercivity (in both as-prepared and annealed states) than the same ferrite synthesized by co-precipitation or oxalyl dihydrazide-nitrate combustion synthesis. Acknowledgments The authors are thankful to Shiraz University of Technology (Department of Materials Science and Engineering) due to support of this research.

16

References [1] P.J. van der Zaag, Ferrites, Reference Module in Materials Science and Materials Engineering, Elsevier2016. [2] K. Okutani, SOFT FERRITE MATERIALS AND THEIR APPLICATIONS, Magnetic Materials in Japan, Elsevier, Oxford, 1991, pp. 213-251. [3] A. Sutka, G. Mezinskis, Sol-gel auto-combustion synthesis of spinel-type ferrite nanomaterials, Frontiers of Materials Science, (2012) 1-14. [4] A. Pradeep, P. Priyadharsini, G. Chandrasekaran, Sol–gel route of synthesis of nanoparticles of MgFe 2 O 4 and XRD, FTIR and VSM study, Journal of Magnetism and Magnetic Materials, 320 (2008) 2774-2779. [5] M. Phadatare, A. Salunkhe, V. Khot, C. Sathish, D. Dhawale, S. Pawar, Thermodynamic, structural and magnetic studies of NiFe 2 O 4 nanoparticles prepared by combustion method: effect of fuel, Journal of Alloys and Compounds, 546 (2013) 314-319. [6] S. Masti, A. Sharma, P. Vasambekar, Study of Behavior of Dielectric Properties in Substituted MgCd ferrites, International Journal of Scientific and Research Publications, 214. [7] N. Kikukawa, M. Takemori, Y. Nagano, M. Sugasawa, S. Kobayashi, Synthesis and magnetic properties of nanostructured spinel ferrites using a glycine–nitrate process, Journal of Magnetism and Magnetic Materials, 284 (2004) 206-214. [8] M. Kaur, S. Rana, P. Tarsikka, Comparative analysis of cadmium doped magnesium ferrite Mg (1− x) Cd x Fe 2 O 4 (x= 0.0, 0.2, 0.4, 0.6) nanoparticles, Ceramics International, 38 (2012) 4319-4323. [9] S. Hajarpour, K. Gheisari, A.H. Raouf, Characterization of nanocrystalline Mg 0.6 Zn 0.4 Fe 2 O 4 soft ferrites synthesized by glycine-nitrate combustion process, Journal of Magnetism and Magnetic Materials, 329 (2013) 165-169. [10] A. Gadkari, T. Shinde, P. Vasambekar, Structural and magnetic properties of nanocrystalline Mg– Cd ferrites prepared by oxalate co-precipitation method, Journal of Materials Science: Materials in Electronics, 21 (2010) 96. [11] N. Deraz, M. Hessien, Structural and magnetic properties of pure and doped nanocrystalline cadmium ferrite, Journal of alloys and compounds, 475 (2009) 832-839. [12] J. Bhosale, S. Kulkarni, R. Sasmile, B. Chougule, Susceptibility and magnetization studies of Gd 3+ substituted Mg-Cd ferrites, Bulletin of Materials Science, 19 (1996) 767-774. [13] E. Salahinejad, M. Hadianfard, D. Macdonald, I. Karimi, D. Vashaee, L. Tayebi, Aqueous sol–gel synthesis of zirconium titanate (ZrTiO 4) nanoparticles using chloride precursors, Ceramics International, 38 (2012) 6145-6149. [14] S. Rastegari, O.S.M. Kani, E. Salahinejad, S. Fadavi, N. Eftekhari, A. Nozariasbmarz, L. Tayebi, D. Vashaee, Non-hydrolytic sol-gel processing of chloride precursors loaded at forsterite stoichiometry, Journal of Alloys and Compounds, 688 (2016) 235-241. [15] N. Namvar, E. Salahinejad, A. Saberi, M.J. Baghjeghaz, L. Tayebi, D. Vashaee, Toward reducing the formation temperature of diopside via wet-chemical synthesis routes using chloride precursors, Ceramics International, (2017). [16] H. Takebe, S. Kobayashi, H. Aono, S. Yamamuro, Chapter 4 - Fabrication and characterization of natural/synthesized, micro-, and nanostructured materials for biomedical applications A2 - Ficai, Denisa, in: A.M. Grumezescu (Ed.) Nanostructures for Novel Therapy, Elsevier2017, pp. 81-106. [17] R.K. Kotnala, J. Shah, Ferrite Materials, Handbook of Magnetic Materials, 23 (2015) 291-379. [18] R. Valenzuela, Magnetic ceramics, Cambridge University Press2005. [19] Z. Yue, L. Li, J. Zhou, H. Zhang, Z. Gui, Preparation and characterization of NiCuZn ferrite nanocrystalline powders by auto-combustion of nitrate–citrate gels, Materials Science and Engineering: B, 64 (1999) 68-72. [20] C. Ehi-Eromosele, B. Ita, E. Iweala, Low-temperature combustion synthesis of cobalt magnesium ferrite magnetic nanoparticles: effects of fuel-to-oxidizer ratio and sintering temperature, Journal of Sol-Gel Science and Technology, 76 (2015) 298-308. 17

[21] A. Salunkhe, V. Khot, M. Phadatare, S. Pawar, Combustion synthesis of cobalt ferrite nanoparticles—influence of fuel to oxidizer ratio, Journal of alloys and compounds, 514 (2012) 9196. [22] V. Khot, A. Salunkhe, M. Phadatare, S. Pawar, Formation, microstructure and magnetic properties of nanocrystalline MgFe 2 O 4, Materials Chemistry and Physics, 132 (2012) 782-787. [23] L. Junliang, Z. Wei, G. Cuijing, Z. Yanwei, Synthesis and magnetic properties of quasi-single domain M-type barium hexaferrite powders via sol–gel auto-combustion: Effects of pH and the ratio of citric acid to metal ions (CA/M), Journal of Alloys and Compounds, 479 (2009) 863-869. [24] K. Wua, T. Tinga, M. Li, W. Ho, Sol-gel autocombustion synthesis of SiO2-doped NiZn ferrite by using various fuels, Journal of Magnetism and Magnetic Materials, 298 (2006) 25-32. [25] K. Mohammed, A. Al-Rawas, A. Gismelseed, A. Sellai, H. Widatallah, A. Yousif, M. Elzain, M. Shongwe, Infrared and structural studies of Mg 1–x Zn x Fe 2 O 4 ferrites, Physica B: Condensed Matter, 407 (2012) 795-804. [26] M.G. Naseri, M.H.M. Ara, E.B. Saion, A.H. Shaari, Superparamagnetic magnesium ferrite nanoparticles fabricated by a simple, thermal-treatment method, Journal of Magnetism and Magnetic Materials, 350 (2014) 141-147. [27] B. Karche, B. Khasbardar, A. Vaingankar, X-ray, SEM and magnetic properties of MgCd ferrites, Journal of magnetism and magnetic materials, 168 (1997) 292-298. [28] H. Mohseni, H. Shokrollahi, I. Sharifi, K. Gheisari, Magnetic and structural studies of the Mndoped Mg–Zn ferrite nanoparticles synthesized by the glycine nitrate process, Journal of Magnetism and Magnetic Materials, 324 (2012) 3741-3747. [29] B. Pourgolmohammad, S. Masoudpanah, M. Aboutalebi, Effect of starting solution acidity on the characteristics of CoFe 2 O 4 powders prepared by solution combustion synthesis, Journal of Magnetism and Magnetic Materials, 424 (2017) 352-358. [30] C.-C. Hwang, J.-S. Tsai, T.-H. Huang, C.-H. Peng, S.-Y. Chen, Combustion synthesis of Ni–Zn ferrite powder—influence of oxygen balance value, Journal of Solid State Chemistry, 178 (2005) 382-389. [31] R. Purohit, B. Sharma, K. Pillai, A. Tyagi, Ultrafine ceria powders via glycine-nitrate combustion, Materials Research Bulletin, 36 (2001) 2711-2721. [32] B.D. Cullity, J.W. Weymouth, Elements of X-ray Diffraction, American Journal of Physics, 25 (1957) 394-395. [33] I. Szczygiel, K. Winiarska, Low-temperature synthesis and characterization of the Mn–Zn ferrite, Journal of thermal analysis and calorimetry, 104 (2011) 577-583. [34] G. Gnanaprakash, S. Ayyappan, T. Jayakumar, J. Philip, B. Raj, Magnetic nanoparticles with enhanced γ-Fe2O3 to α-Fe2O3 phase transition temperature, Nanotechnology, 17 (2006) 5851. [35] A. Verma, T. Goel, R. Mendiratta, R. Gupta, High-resistivity nickel–zinc ferrites by the citrate precursor method, Journal of magnetism and magnetic materials, 192 (1999) 271-276. [36] A. Verma, R. Chatterjee, Effect of zinc concentration on the structural, electrical and magnetic properties of mixed Mn–Zn and Ni–Zn ferrites synthesized by the citrate precursor technique, Journal of magnetism and magnetic materials, 306 (2006) 313-320. [37] X. Liang, Y. Zhong, W. Tan, J. Zhu, P. Yuan, H. He, Z. Jiang, The influence of substituting metals (Ti, V, Cr, Mn, Co and Ni) on the thermal stability of magnetite, Journal of thermal analysis and calorimetry, 111 (2013) 1317-1324. [38] E. Wolska, W. Wolski, J. Kaczmarek, E. Riedel, D. Prick, Defect structures in cadmium-nickel ferrites, Solid state ionics, 51 (1992) 231-237. [39] K. Nakamoto, Infrared Spectra of Inorganic and Coordination Compounds, 2 nd ed., John Wiley & Sons, (1970) 232–244. [40] M. Gharagozlou, R. Bayati, Low temperature processing and magnetic properties of zinc ferrite nanoparticles, Superlattices and Microstructures, 78 (2015) 190-200. [41] A.E.B. R. H. Kodama, E. J. McNiff, Jr., and S. Foner, Surface Spin Disorder in NiFe2O4 Nanoparticles, Phys Rev Lett., 77 ( 1996) 394-397. 18

[42] G. Herzer, Grain size dependence of coercivity and permeability in nanocrystalline ferromagnets, IEEE Transactions on Magnetics, 26 (1990) 1397-1402.

19

Figures captions Fig. 1. XRD patterns of as- synthesized ferrites at different G.N-1 ratio. Fig. 2. DSC-TG curves of the prepared gels at (a) G.N-1 = 0.30, (b) G.N-1 = 0.55 and (c) G.N-1 = 0.85. Fig. 3. XRD patterns of ferrites produced at G.N-1 = 0.30 after annealing for 2 h at different temperatures. Fig. 4. XRD patterns of ferrites synthesized at G.N-1 ratios of 0.55 and 0.85, after heat treatment for 2 h at different temperatures. Fig. 5. FTIR spectra of xerogel produced at G.N-1 = 0.55. Fig. 6. FTIR spectra of as-prepared and annealed ferrites prepared at different G.N-1 ratios. Fig. 7. FE-SEM images of as-synthesized ferrites obtained at (a) and (b) G.N-1 = 0.30, (c) and (d) G.N-1 = 0.55, and (e) and (f) G.N-1 = 0.85. Fig. 8. FE-SEM images of ferrite prepared at G.N-1 = 0.30, after heat treatment at 1173 K for 2 h. Fig. 9. Frequency distribution chart of particle size determined from the FE-SEM images of the powders produced at (a) G.N-1 = 0.30, (b) G.N-1 = 0.55, (c) G.N-1 = 0.85 and (d) G.N-1 = 0.30 (1173 K, 2 h). Fig. 10.Room temperature magnetic hysteresis curves of as- prepared ferrites at different G.N-1 ratios. Fig. 11. Evolution of coercivity, Hc, of the as-prepared ferrites as a function of their crystallite size. Fig. 12. Magnetic hysteresis loops, measured at ambient temperature, for annealed ferrites prepared at various G.N-1 values.

20

Tables captions Table 1. Thermodynamic data used for calculating the adiabatic flame temperature, Tad, upon combustion [5, 9]. Table 2. Variations of adiabatic flame temperature, Tad, reaction enthalpy,

, and number

of moles of gases evolved during combustion at different G.N-1 ratios. Table 3. Effects of G.N-1 ratios and annealing temperature on crystallite size, D, lattice parameter, a, saturation magnetization, Ms, and coercivity, Hc, of the prepared ferrites. Table 4. The ratios of different elements determined from the ICP-OES test of the assynthesized ferrites produced at different G.N-1 values.

Fig. 1

21

Fig. 2

22

Fig. 3

Fig. 4

23

Fig. 5

Fig. 6

24

Fig. 7

25

Fig. 8

26

Fig. 9

27

Fig. 10

Fig. 11 28

Fig. 12

Fig. 13

Compound Mg(NO3)2.6H2O Cd(NO3)2.4H2O* Fe(NO3)3.9H2O NH2CH2COOH Mg0.6Cd0.4Fe2O4** CO2 (g) H2O (g)

Enthalpy of formation (kJ.mol-1) -2612.8243 -483.9633 -3285.2768 -333.5066 -1179.1000 -393.5052 -241.7934 29

Heat capacity cp (J.mol-1K-1) 147.9291 43.26 + 0.01146T 30.12 + 0.0151T

N2 (g) O2 (g)

27.2 + 0.0042T 24.77 + 0.01536T

0 0

G.N-1 ratio

T (K)

D (nm)

a (nm)

Ms (Am2.kg-1)

nB (μB)

Hc (kA.m-1)

0.30 0.42 0.55 0.70 0.85 0.30 0.30

973 1123

8 32 43 40 38 29 81

0.8477 0.8386 0.8366 0.8359 0.8370 0.8502 0.8509

16.08 45.98 56.44 51.11 45.32 33.84 40.25

0.67 1.42 1.70

1.56 6.76 6.12 6.68 7.32 2.07 2.48

*

Due to a lack of thermodynamic data for Cd(NO3)2.4H2O, the formation enthalpy of Cd(NO3)2 was used. **

Due to the unavailability of thermodynamic data for Mg0.6Cd0.4Fe2O4, enthalpy of formation and heat capacity of ZnFe2O4 [9], are used to perform temperature calculations, because both Cd2 + and Zn2 + cations tend to form normal spinel structure.

G.N-1 ratio

Number of moles of gases released

Tad (K)

0.30 0.42 0.55 0.70 0.85

43.80 46.44 49.40 55.20 61.20

843.37 1263.12 1636.18 1904.31 2117.44

30

(kJ.mol-1) -996.0417 -2011.7093 -3154.3354 -4381.6005 -5651.1851

0.30 0.55 0.85 0.85

1173 1073 873 1173

149 49 41 48

0.8511 0.8388 0.8365 0.8399 Table 3

41.70 21.47 26.21 -

1.75 -

G.N-1 ratio

Cd/Mg

Cd/Fe

Mg/Fe

0.30

0.657

0.272

0.353

0.55

0.099

0.035

0.347

0.85

0.021

0.008

0.344

31

1.92 8.13 8.61 -