The XRD spectra of the TCP‐TiO2‐MgF2

F) and Mg2

(PO4

F and the fluorapatite (Ca5

2Ca3

)F and Ca5

of the TCP‐38 wt% TiO2

‐4 wt% MgF2

(PO4 )3 F].

The diffractograms of the TCP‐40 wt% TiO2

(PO4

compounds. The interaction between TCP and MgF2

(PO4 )3

‐4 wt% MgF2

, the mechanical properties of the TCP‐TiO2

titania (**Figure 17**). The spectrum of the TCP‐40 wt% TiO2

close to those of the enamel (**Table 4**). In fact, the optimum values of the Shear modulus (G) and Young's modulus (E) of the composites reached 20 and 51 GPa, respectively. If we compare those

notice that the elastic corresponding modulus (E and G) of the composites sintered with 4 wt%

are obtained after the sintering process for 1 h. In fact, for longer sintering dwell time, these

composites sintered for 60 min are improved which is probably due to the formation of the

‐MgF2

mum values of the mechanical properties (E = 51 GPa, G = 20 GPa, H = 360 Hv and σr

properties are hindered. Indeed, we can retain that the performances (σr

)F was obtained only for 2.5, 4 and 4.5 wt% of MgF2

(PO4

+ CaTiO3(s) + Mg2

increase generally about 18 GPa for Young's modulus and 4.3 GPa for the Shear modulus.

composites with different length of the sintering time is not improved. Thus, the opti‐

composites sintered for 300 min or the TCP‐35 wt% TiO2

composites structure enhances their performances with 4 wt%MgF2

indicate that the majority of peaks are relative to the rutile phase of

(**Figure 17**).This disappearance is attributed to the incorporation of this

F) according to the following reaction (Eq. (7)):

(PO4

‐MgF2

)2(s) + 2MgF2(s) + TiO2(s) + H2

composites sintered at 1200°C for 1 h. In fact, the continuous phases formed and the

Several previous analyses (XRD and MAS‐NMR) confirm the presence of these compounds

SEM analyses prove that excellent performances are attained with the TCP‐38 wt% TiO2

new compounds accompanied with the liquid phase appearance improved the performances

the XRD analysis. Thus, the new lamella form detected in the microstructure of the TCP‐38

solid reaction between MgF2 and TCP according to the previous reaction (Eq. (7)). Above 4

the grain growth presence and the bubbles form in the microstructure of composites. Those

composites is probably accorded to the Mg2

)F) into the calcium phosphate structure by a solid state reaction through the

)F which is in agreement with the 31P MAS‐NMR analysis. Thus,

‐4 wt% MgF2

http://dx.doi.org/10.5772/intechopen.68501

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

composites sintered without MgF2

) of the TCP‐38 wt% TiO2

composites. Consequently, the MgF2

and α‐TCP besides the rutile phase and the β‐TCP.

composites depicts the formation of both fluorapatite

composites sintered at 1200°C with different

promotes the formation of the Mg2

O ↔ Ca5

composites. These results confirm those obtained with

(PO4

composites sintered without MgF2

and disappears for higher

(PO4 )3 F(s)

)F(s) + 2HF(g) (7)

)F derived from the

composites were hindered by

composites are

, we

69

‐4 wt%

added

(PO4 )

‐4

= 27 MPa)

‐10 wt% MgF2

.

and H) of the TCP‐38

The densification discrepancy of the samples prepared by a mixture of the TCP‐40 wt% TiO2 composites with adding 4 wt% MgF2 (ρ = 94%) and without MgF2 (ρ = 89%) may be due to different phenomena. Thus, the formation of the new compounds and the liquid phase gives rise to the performances of the composites. In general, the very important conditions for the densification by liquid phase sintering are the low viscosity of liquid for fast diffusion of solid through liquid and the solubility of solid to liquid [72] which is responsible for the significant improvement of these properties. Furthermore, several studies have shown the formation of the liquid phase during the sintering process of the tricalcium phosphate/titania system [12, 22]. Thus, Caroff et al. proved that a binary eutectic between the tricalcium phosphate and the titania is relative to this liquid phase [22]. So, in this study, there are two rate‐controlling processes in the liquid phase sintering: the dissolution‐precipitation phenomena which are dominant sintering mechanism in enhancing the densification by smoothing the contact inter‐ face and dissolution of the finer particles and the grain rearrangement occurs due to capillary force between particles. These results are confirmed by literature [72]. So, liquid phase sinter‐ ing helps to fill the pore in the microstructure of the TCP‐TiO2 ‐MgF2 composites.

Several examples of the mechanical properties of the bone tissues are shown in **Table 4**. The optimum value of Vickers hardness of the TCP‐40 wt% TiO2 composites was obtained after the addition of 4 wt% MgF2 (H = 360 Hv). If we compare this result with that of the TCP‐40 wt% TiO2 composites without adding the magnesium fluoride, we notice an increase from 270 to 360 Hv. So, adding 4 wt% MgF2 enhances Vickers hardness of the TCP‐40 wt% TiO2 com‐ posites about 90 Hv. This value (H = 360 Hv) is so close to the value of the enamel (340–370 Hv) (**Table 4**).

**Table 4.** Literature examples of the mechanical properties of the bone tissues.

The values of the elastic modulus (E and G) of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites are close to those of the enamel (**Table 4**). In fact, the optimum values of the Shear modulus (G) and Young's modulus (E) of the composites reached 20 and 51 GPa, respectively. If we compare those values with the elastic modulus of the TCP‐40 wt% TiO2 composites sintered without MgF2 , we notice that the elastic corresponding modulus (E and G) of the composites sintered with 4 wt% MgF2 increase generally about 18 GPa for Young's modulus and 4.3 GPa for the Shear modulus.

TiO2

68 Sintering of Functional Materials

composites. Such similarity between the present work and the enamel behaviors allows

(ρ = 94%) and without MgF2

The densification discrepancy of the samples prepared by a mixture of the TCP‐40 wt% TiO2

different phenomena. Thus, the formation of the new compounds and the liquid phase gives rise to the performances of the composites. In general, the very important conditions for the densification by liquid phase sintering are the low viscosity of liquid for fast diffusion of solid through liquid and the solubility of solid to liquid [72] which is responsible for the significant improvement of these properties. Furthermore, several studies have shown the formation of the liquid phase during the sintering process of the tricalcium phosphate/titania system [12, 22]. Thus, Caroff et al. proved that a binary eutectic between the tricalcium phosphate and the titania is relative to this liquid phase [22]. So, in this study, there are two rate‐controlling processes in the liquid phase sintering: the dissolution‐precipitation phenomena which are dominant sintering mechanism in enhancing the densification by smoothing the contact inter‐ face and dissolution of the finer particles and the grain rearrangement occurs due to capillary force between particles. These results are confirmed by literature [72]. So, liquid phase sinter‐

Several examples of the mechanical properties of the bone tissues are shown in **Table 4**. The

posites about 90 Hv. This value (H = 360 Hv) is so close to the value of the enamel (340–370

 **(Hv) Eb (GPa) Gc**

Dentin 40–75 18 – [6, 8, 69] Enamel 340–370 50–82 – [8, 69, 70] Thigh bone – 20 – [6, 67, 68] Cortical bone – 7–25 – [6, 67, 68]

composites without adding the magnesium fluoride, we notice an increase from 270

270 33.1 15.7 Present work

360 51 20 Present work

‐4 wt% MgF2

‐MgF2

(H = 360 Hv). If we compare this result with that of the TCP‐40

enhances Vickers hardness of the TCP‐40 wt% TiO2

composites.

 **(GPa) References**

composites was obtained after

com‐

composites in biomedical

(ρ = 89%) may be due to

us to study the performances of the TCP‐38 wt% TiO2

ing helps to fill the pore in the microstructure of the TCP‐TiO2

optimum value of Vickers hardness of the TCP‐40 wt% TiO2

**Table 4.** Literature examples of the mechanical properties of the bone tissues.

applications and specifically as enamel.

composites with adding 4 wt% MgF2

the addition of 4 wt% MgF2

to 360 Hv. So, adding 4 wt% MgF2

wt% TiO2

Hv) (**Table 4**).

β‐TCP‐40 wt% TiO2 composites

‐4 wt% MgF2 composites

Vickers hardness.

Young's modulus.

Shear modulus.

β‐TCP‐38 wt%

TiO2

a

b

c

**Materials Ha**

The evolution of the mechanical properties (E, G, H and σr ) of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites with different length of the sintering time is not improved. Thus, the opti‐ mum values of the mechanical properties (E = 51 GPa, G = 20 GPa, H = 360 Hv and σr = 27 MPa) are obtained after the sintering process for 1 h. In fact, for longer sintering dwell time, these properties are hindered. Indeed, we can retain that the performances (σr and H) of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites sintered for 300 min or the TCP‐35 wt% TiO2 ‐10 wt% MgF2 composites sintered for 60 min are improved which is probably due to the formation of the FAp with more quantities in the TCP‐TiO2 ‐MgF2 composites. Consequently, the MgF2 added to the TCP‐TiO2 composites structure enhances their performances with 4 wt%MgF2 .

The diffractograms of the TCP‐40 wt% TiO2 composites sintered at 1200°C with different percentages of MgF2 indicate that the majority of peaks are relative to the rutile phase of titania (**Figure 17**). The spectrum of the TCP‐40 wt% TiO2 composites sintered without MgF2 proves the presence of traces of CaTiO3 and α‐TCP besides the rutile phase and the β‐TCP. The XRD spectra of the TCP‐TiO2‐MgF2 composites depicts the formation of both fluorapatite (Ca5 (PO4 )3 F) and Mg2 (PO4 )F which is in agreement with the 31P MAS‐NMR analysis. Thus, the Mg2 (PO4 )F was obtained only for 2.5, 4 and 4.5 wt% of MgF2 and disappears for higher amounts of MgF2 (**Figure 17**).This disappearance is attributed to the incorporation of this phase (Mg2 (PO4 )F) into the calcium phosphate structure by a solid state reaction through the compounds. The interaction between TCP and MgF2 promotes the formation of the Mg2 (PO4 ) F and the fluorapatite (Ca5 (PO4 )3 F) according to the following reaction (Eq. (7)):

$$2\text{Ca}\_3\text{(PO}\_4\text{)}\_{2(s)} + 2\text{MgF}\_{2(s)} + \text{TiO}\_{2(s)} + \text{H}\_2\text{O} \quad \leftrightarrow \text{Ca}\_5\text{(PO}\_4\text{)}\_3\text{F}\_{(s)}$$

$$+ \text{CaTiO}\_{3(s)} + \text{Mg}\_2\text{(PO}\_4\text{)F}\_{(s)} + 2\text{HF}\_{(g)}\tag{7}$$

Several previous analyses (XRD and MAS‐NMR) confirm the presence of these compounds [Mg2 (PO4 )F and Ca5 (PO4 )3 F].

SEM analyses prove that excellent performances are attained with the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites sintered at 1200°C for 1 h. In fact, the continuous phases formed and the new compounds accompanied with the liquid phase appearance improved the performances of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites. These results confirm those obtained with the XRD analysis. Thus, the new lamella form detected in the microstructure of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites is probably accorded to the Mg2 (PO4 )F derived from the solid reaction between MgF2 and TCP according to the previous reaction (Eq. (7)). Above 4 wt% MgF2 , the mechanical properties of the TCP‐TiO2 ‐MgF2 composites were hindered by the grain growth presence and the bubbles form in the microstructure of composites. Those bubbles resulted from the HF gas evaporation according to the reaction (Eq. (7)). The micro‐ graph of the TCP‐35 wt% TiO2 ‐10 wt% MgF2 composites indicates the hexagonal grain forms relative to the FAp formation. This phenomenon was accelerated by the dissolution‐precipita‐ tion phenomena in the liquid phase. Therefore, the liquid phase sintering is used for homog‐ enization and consolidation which presents the advantages of this production method. Thus, the densification of those composites resulted from three processes: the dissolution‐precipi‐ tation, the rearrangement and the coalescence. This fact is previously confirmed by litera‐ ture [72–74]. Generally, the microstructure of sintered body by liquid phase sintering mode includes secondary phase, which influences the mechanical, chemical and thermal properties.

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Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

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## **5. Conclusion**

The effect of the MgF2 addition on the performances of the TCP‐40 wt% TiO2 composites was studied during the sintering process at 1200°C for different length of the sintering time. The addition of 4 wt% MgF2 to the TCP‐40 wt% TiO2 composites led to a maximum of both the densification and the mechanical properties after the sintering process at 1200°C for 1 h. A notable amelioration in the elastic modulus and Vickers hardness of the TCP‐40 wt% TiO2 composites was attained with 4 wt% MgF2 (E = 51 GPa, G = 20 GPa and H = 360 Hv). The enhancement of the composites properties resulted from the formation of new compounds and the liquid phase which permits to fill the pores in the microstructure. Thus, the micro‐ structure indicates the presence of new lamella form relative to the Mg2 (PO4 )F and the hexag‐ onal grains relative to FAp. Above 4 wt% MgF2 , the properties of the composites are hindered by the grain growth formation and the presence of the bubbles form as a result of the dissolu‐ tion‐precipitation phenomena.

## **Author details**

Ibticem Ayadi and Foued Ben Ayed\*

\*Address all correspondence to<EMAIL_ADDRESS>
Laboratory of Industrial Chemistry, National Engineering School, Sfax University, Sfax, Tunisia

## **References**

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bubbles resulted from the HF gas evaporation according to the reaction (Eq. (7)). The micro‐

relative to the FAp formation. This phenomenon was accelerated by the dissolution‐precipita‐ tion phenomena in the liquid phase. Therefore, the liquid phase sintering is used for homog‐ enization and consolidation which presents the advantages of this production method. Thus, the densification of those composites resulted from three processes: the dissolution‐precipi‐ tation, the rearrangement and the coalescence. This fact is previously confirmed by litera‐ ture [72–74]. Generally, the microstructure of sintered body by liquid phase sintering mode includes secondary phase, which influences the mechanical, chemical and thermal properties.

addition on the performances of the TCP‐40 wt% TiO2

studied during the sintering process at 1200°C for different length of the sintering time. The

densification and the mechanical properties after the sintering process at 1200°C for 1 h. A notable amelioration in the elastic modulus and Vickers hardness of the TCP‐40 wt% TiO2

enhancement of the composites properties resulted from the formation of new compounds and the liquid phase which permits to fill the pores in the microstructure. Thus, the micro‐

by the grain growth formation and the presence of the bubbles form as a result of the dissolu‐

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composites indicates the hexagonal grain forms

composites led to a maximum of both the

(PO4

, the properties of the composites are hindered

(E = 51 GPa, G = 20 GPa and H = 360 Hv). The

composites was

)F and the hexag‐

‐10 wt% MgF2

to the TCP‐40 wt% TiO2

structure indicates the presence of new lamella form relative to the Mg2

graph of the TCP‐35 wt% TiO2

70 Sintering of Functional Materials

**5. Conclusion**

The effect of the MgF2

addition of 4 wt% MgF2

tion‐precipitation phenomena.

Ibticem Ayadi and Foued Ben Ayed\*

**Author details**

**References**

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onal grains relative to FAp. Above 4 wt% MgF2

\*Address all correspondence to<EMAIL_ADDRESS>

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**Chapter 4**

**Provisional chapter**

**Evolution of Magnetic Properties in Ferrites: Trends of**

Microstructure of magnetic materials greatly influences the performance of magnetic properties, and sintering has been used as an agent to tailor the microstructure of these magnetic materials especially ferrites. Nanostructured ferrites prepared by high-energy milling method are often inherently unstable owing to their small constituent sizes, nonequilibrium cation distribution, disordered spin configuration, and high chemical activity. Therefore, sintering of the milled ferrites recrystallizes the nanostructure and causes its transition from an excited metastable (activated) state into the low-energy crystalline state. A better understanding of the response of nanoscale ferrites with changes in temperature is crucial not only for basic science (the development of an atomistic and microscopic theory of the mechanochemical processes) but also because of the technological high-temperature applications in catalysis, ferrofluids and information storage. This chapter discusses on two different sintering schemes, which are a commonly applied multi-sample sintering and a rarely adopted single-sample sintering. Experimental results of single-sample and multi-sample sintering of NiZn ferrites and yttrium iron garnet (YIG) were highlighted, and their microstructural consequences on the magnetic

**Evolution of Magnetic Properties in Ferrites: Trends of**

DOI: 10.5772/intechopen.68500

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution,

© 2018 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

and reproduction in any medium, provided the original work is properly cited.

Microstructure of polycrystalline ceramics is usually complex, consisting of grains, grain boundaries, porosity and secondary phases. This kind of structure is not seen in single crystals. Variations in the microstructure with different kinds of shape, size, distribution and

**Keywords:** microstructural evolution, BH-hysteresis, ferrites, NiZn ferrites,

**Single-Sample and Multi-Sample Sintering**

**Single-Sample and Multi-Sample Sintering**

Ismayadi Ismail, Idza Riati Ibrahim and

Ismayadi Ismail, Idza Riati Ibrahim and

http://dx.doi.org/10.5772/intechopen.68500

properties were also discussed.

**1. Sintering as a microstructure tailoring agent**

yttrium iron garnet (YIG)

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Rodziah Nazlan

**Abstract**

Rodziah Nazlan

**Provisional chapter**

## **Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering Single-Sample and Multi-Sample Sintering**

**Evolution of Magnetic Properties in Ferrites: Trends of**

DOI: 10.5772/intechopen.68500

Ismayadi Ismail, Idza Riati Ibrahim and Rodziah Nazlan Rodziah Nazlan Additional information is available at the end of the chapter

Ismayadi Ismail, Idza Riati Ibrahim and

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.68500

#### **Abstract**

[62] Gozalian A, Behnamghader A, Daliri M, Moshkforoush A. Synthesis and thermal behav‐ ior of Mg‐doped calcium phosphate nanopowders via the sol gel method. Scientia Iranica Transactions F: Nanotechnology. 2011;**18**(6):1614‐1622. DOI: 10.1016/j.scient.2011.11.014

[63] Serre CM, Papillard M, Chavassieux P, Voegel JC, Boivin G. Influence of magnesium substitution on a collagen‐apatite biomaterial on the production of a calcifying matrix by human osteoblasts. Journal of Biomedical Materials Research. 1998;**42**:626‐633. DOI:

[64] Suchanek WL, Byrappa K, Shuk P, Riman RE, Janas VF, TenHuisen KS. Mechanochemical‐ hydrothermal synthesis of calcium phosphate powders with coupled magnesium and carbonate substitution. Journal of Solid State Chemistry. 2004;**177**:793‐799. DOI: http://

[65] LeGeros RZ. Calcium phosphates in oral biology and medicine. In: Myers HM, editor. Handbook of Monographs in Oral Sciences.1st ed. Basel: Karger; 1991. pp. 1‐200. PMID:

[66] Otsuka M, Oshinbe A, LeGeros RZ, Tokudome Y, Ito A, Otsuka K, Higuchi WI. Efficacy of the injectable calcium phosphate ceramics suspensions containing magnesium, zinc and fluoride on the bone mineral deficiency in ovariectomized rats. Journal of

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[71] Staines M, Robinson WH, Hood JAA. Spherical indentation of tooth enamel. Journal of

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[73] Ben Ayed F, Bouaziz J, Bouzouita K. Pressureless sintering of Fluorapatite under oxy‐ gen atmosphere. Journal of the European Ceramic Society. 2000;**20**(8):1069‐1076. DOI:

[74] Ben Ayed F, Bouaziz J, Bouzouita K. Calcination and sintering of fluorapatite under argon atmosphere. Journal of Alloys and Compounds. 2001;**322**(1‐2):238‐245. DOI: http:// dx.doi.org/10.1016/S0925‐8388(01)01200‐2 DOI:10.1016/S0925‐8388(01)01200‐2#doilink

dx.doi.org/10.1016/j.jssc.2003.09.012 DOI:10.1016/j.jssc.2003.09.012#doilink

10.1002/(SICI)1097‐4636(19981215)42:4<626::AID‐JBM20>3.0.CO;2‐S

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1870604

76 Sintering of Functional Materials

dx.doi.org/10.5402/2013/253761

p. ISBN:<PHONE_NUMBER>028

10.1016/S0955‐2219(99)00272‐1

Microstructure of magnetic materials greatly influences the performance of magnetic properties, and sintering has been used as an agent to tailor the microstructure of these magnetic materials especially ferrites. Nanostructured ferrites prepared by high-energy milling method are often inherently unstable owing to their small constituent sizes, nonequilibrium cation distribution, disordered spin configuration, and high chemical activity. Therefore, sintering of the milled ferrites recrystallizes the nanostructure and causes its transition from an excited metastable (activated) state into the low-energy crystalline state. A better understanding of the response of nanoscale ferrites with changes in temperature is crucial not only for basic science (the development of an atomistic and microscopic theory of the mechanochemical processes) but also because of the technological high-temperature applications in catalysis, ferrofluids and information storage. This chapter discusses on two different sintering schemes, which are a commonly applied multi-sample sintering and a rarely adopted single-sample sintering. Experimental results of single-sample and multi-sample sintering of NiZn ferrites and yttrium iron garnet (YIG) were highlighted, and their microstructural consequences on the magnetic properties were also discussed.

**Keywords:** microstructural evolution, BH-hysteresis, ferrites, NiZn ferrites, yttrium iron garnet (YIG)

### **1. Sintering as a microstructure tailoring agent**

Microstructure of polycrystalline ceramics is usually complex, consisting of grains, grain boundaries, porosity and secondary phases. This kind of structure is not seen in single crystals. Variations in the microstructure with different kinds of shape, size, distribution and

orientation of the grains play a key role in many of the macroscopic properties including magnetic, thermophysical, mechanical, electrical and many other properties. Essentially, these phenomena are familiar with the polycrystalline ceramics samples having micronic grain size, and the information on their relationships is well understood. Materials in the micrometer scale mostly exhibit physical properties the same as that of bulk form; however, materials in the nanometer scale may exhibit physical properties distinctively different from that of micrometer scale. Nanomaterials may have significantly lower melting point or phase transition temperature and appreciably reduced lattice constants, due to a huge fraction of surface atoms in the total amount of atoms [1]. Materials with an altered 'nano'-microstructure provide potential for new or improved applications [2]. Sintering has been known as an agent to alter the microstructure condition of a polycrystalline material. Through optimization of sintering conditions such as sintering temperature, sintering atmosphere, heating and cooling rates, sintering time and partial pressure of sintering atmosphere, the best materials properties could be achieved. Tailoring the microstructure to attain certain desirable materials properties is the main challenge and of interest in material science.

## **2. Single-sample and multi-sample sintering**

There are two different sintering schemes in producing polycrystalline materials, which are commonly applied multi-sample sintering scheme and rarely adopted single-sample sintering scheme. Generally, reported studies involving sintering and materials properties employed multi-sample sintering scheme [3–9]. The multi-sample sintering has as many starting compacts as the number of the intended sintering temperatures where each sample sintered only once at different temperatures. All compacts are assumed to have identical morphologies, for example, particle size distribution. However, a rarely adopted singlesample sintering scheme has only one single compact with definite starting point and one particular particle size distribution where only one sample sintered at different temperatures. Therefore, multi-sample sintering is subjected to possible statistical errors since the particle size distributions for all the samples may not be as identical as assumed as compared to that of single-sample sintering. Thus, more convincing data could be obtained for the scientific interpretation of the evolution study. Besides, it is more economical with respect to raw materials and sample preparation time. The schematic of the different sintering schemes is been shown in **Figure 1**. Ceramic is defined as the art and science of making and using solid articles [10]. The statement made a clear view that it is not an easy task to produce the same ceramic with almost the same properties because ceramic is composed of a complex system. The question is whether a material scientist is able to make such ceramic by employing single-sample and multi-sample sintering, with almost the same or enhanced properties is of great interest.

**3. Magnetic properties evolution and the research gap**

**Figure 1.** Schematic of different sintering schemes: multi-sample and Single-sample sintering.

Evolution in magnetic properties is laterally correlated with the evolution of the microstructure, particularly from nanometric to micronic regime of grains as shown in **Figure 2**. However, the reported cause and effect sequences, in the magnetic properties research literature, are an experimental sequence focused mainly on yielding the final outcome, for example, the final microstructure-magnetic properties relationship at final sintering temperature. Therefore, microstructural dependence of magnetic properties for polycrystalline ferrite having micrometer grain size has been widely studied and greatly understood. However, research on ferrite from the nanometer scale has been a field of intense study, due to the novel properties shown by particles located in the transition region between the isolated atoms and bulk solids. Their novel properties make them attractive, both from the scientific knowledge of understanding

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

http://dx.doi.org/10.5772/intechopen.68500

79

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering http://dx.doi.org/10.5772/intechopen.68500 79

**Figure 1.** Schematic of different sintering schemes: multi-sample and Single-sample sintering.

#### **3. Magnetic properties evolution and the research gap**

orientation of the grains play a key role in many of the macroscopic properties including magnetic, thermophysical, mechanical, electrical and many other properties. Essentially, these phenomena are familiar with the polycrystalline ceramics samples having micronic grain size, and the information on their relationships is well understood. Materials in the micrometer scale mostly exhibit physical properties the same as that of bulk form; however, materials in the nanometer scale may exhibit physical properties distinctively different from that of micrometer scale. Nanomaterials may have significantly lower melting point or phase transition temperature and appreciably reduced lattice constants, due to a huge fraction of surface atoms in the total amount of atoms [1]. Materials with an altered 'nano'-microstructure provide potential for new or improved applications [2]. Sintering has been known as an agent to alter the microstructure condition of a polycrystalline material. Through optimization of sintering conditions such as sintering temperature, sintering atmosphere, heating and cooling rates, sintering time and partial pressure of sintering atmosphere, the best materials properties could be achieved. Tailoring the microstructure to attain certain desirable materials properties is the main challenge and of interest in

There are two different sintering schemes in producing polycrystalline materials, which are commonly applied multi-sample sintering scheme and rarely adopted single-sample sintering scheme. Generally, reported studies involving sintering and materials properties employed multi-sample sintering scheme [3–9]. The multi-sample sintering has as many starting compacts as the number of the intended sintering temperatures where each sample sintered only once at different temperatures. All compacts are assumed to have identical morphologies, for example, particle size distribution. However, a rarely adopted singlesample sintering scheme has only one single compact with definite starting point and one particular particle size distribution where only one sample sintered at different temperatures. Therefore, multi-sample sintering is subjected to possible statistical errors since the particle size distributions for all the samples may not be as identical as assumed as compared to that of single-sample sintering. Thus, more convincing data could be obtained for the scientific interpretation of the evolution study. Besides, it is more economical with respect to raw materials and sample preparation time. The schematic of the different sintering schemes is been shown in **Figure 1**. Ceramic is defined as the art and science of making and using solid articles [10]. The statement made a clear view that it is not an easy task to produce the same ceramic with almost the same properties because ceramic is composed of a complex system. The question is whether a material scientist is able to make such ceramic by employing single-sample and multi-sample sintering, with almost the same or enhanced

material science.

78 Sintering of Functional Materials

properties is of great interest.

**2. Single-sample and multi-sample sintering**

Evolution in magnetic properties is laterally correlated with the evolution of the microstructure, particularly from nanometric to micronic regime of grains as shown in **Figure 2**. However, the reported cause and effect sequences, in the magnetic properties research literature, are an experimental sequence focused mainly on yielding the final outcome, for example, the final microstructure-magnetic properties relationship at final sintering temperature. Therefore, microstructural dependence of magnetic properties for polycrystalline ferrite having micrometer grain size has been widely studied and greatly understood. However, research on ferrite from the nanometer scale has been a field of intense study, due to the novel properties shown by particles located in the transition region between the isolated atoms and bulk solids. Their novel properties make them attractive, both from the scientific knowledge of understanding

properties, particularly in magnetic properties. This absence information has leaving behind

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

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81

i. How microstructural properties evolve with the magnetic properties from several nanometers

ii. How the evolution of an amorphous-crystalline mixture state to complete polycrystalline

iii. Does sample with nanometer grain size demonstrates the similar magnetic properties with

iv. What is the relationship of evolving microstructure properties with magnetic properties

v. When is the remarkable transition of magnetic properties between the nanometer and

Hence, revealing the systematic development of grains having size from several nanometers up to micrometer is an important parameter of designing best materials properties. The evolution has yet to be established in wide variation of properties since the knowledge of parallel

i. Other electroceramics, for example, high-temperature heating elements and electrodes, voltage dependent resistors, thermally sensitive resistors, solid fast-ion conductors, hu-

v. Mechanical metals and metallic alloys (elemental), for example, pure iron, copper, chro-

Generally, a class of magnetic oxide, which contains iron oxide as a primary component, is commonly described as ferrites. Similar to most ceramic material, the physical properties of ferrite are also hard but brittle. The drastic progress in the development and growth of ferrites for technological application has once force the industry to leave behind the research and study in ferrite. The industrial importance of ferrites becomes apparent when one examines the diversity of their applications. Ferrite has been extensively used in various electronic devices. These applications include choke filters [11], transformers [12], antenna rods [13], microwave devices [14, 15], isolators [16], circulators [17], phase shifters [18] and many others. The frequencies of the applications range from direct current (DC) to the highest one at which

O3

Fe14B, SmCo17

evolution of microstructure and properties is absence in these materials:

iv. Transition metal-alloys based magnetic materials, for example, NiFe, NiCo

many research gaps and research questions that have to be solved in this study:

state affects the microstructure-magnetic properties?

to micrometer grains size?

of material?

micrometer grain size?

mium, nickel/carbon steels

**4. Introduction to ferrites**

any electronic device can function [19].

samples having micron grain size?

midity and gas sensors, superconductors

ii. Thermo-mechanical ceramics, for example, SiC, Al2

iii. Rare-earths-based magnetic materials, for example, Nd9

vi. Polycrystalline semiconductors, for example, zinc oxide

**Figure 2.** Grand evolution-data acquisition scheme for polycrystalline materials.

their properties, and the technological importance of enhancing the performance of the present materials. A fundamental line of scientific enquiry thus has been neglected, particularly by ferrite researchers for more than 70 years: What would be the magnetic-microstructure relationships at various intermediate sintering conditions during the parallel evolutions of the morphology and magnetic properties? Therefore, much possible essential development information has been neglected, thus reducing the capabilities of producing good fundamental scientific knowledge, which lies behind the parallel evolution of the microstructure-material properties, particularly in magnetic properties. This absence information has leaving behind many research gaps and research questions that have to be solved in this study:

Hence, revealing the systematic development of grains having size from several nanometers up to micrometer is an important parameter of designing best materials properties. The evolution has yet to be established in wide variation of properties since the knowledge of parallel evolution of microstructure and properties is absence in these materials:

## **4. Introduction to ferrites**

their properties, and the technological importance of enhancing the performance of the present materials. A fundamental line of scientific enquiry thus has been neglected, particularly by ferrite researchers for more than 70 years: What would be the magnetic-microstructure relationships at various intermediate sintering conditions during the parallel evolutions of the morphology and magnetic properties? Therefore, much possible essential development information has been neglected, thus reducing the capabilities of producing good fundamental scientific knowledge, which lies behind the parallel evolution of the microstructure-material

**Figure 2.** Grand evolution-data acquisition scheme for polycrystalline materials.

80 Sintering of Functional Materials

Generally, a class of magnetic oxide, which contains iron oxide as a primary component, is commonly described as ferrites. Similar to most ceramic material, the physical properties of ferrite are also hard but brittle. The drastic progress in the development and growth of ferrites for technological application has once force the industry to leave behind the research and study in ferrite. The industrial importance of ferrites becomes apparent when one examines the diversity of their applications. Ferrite has been extensively used in various electronic devices. These applications include choke filters [11], transformers [12], antenna rods [13], microwave devices [14, 15], isolators [16], circulators [17], phase shifters [18] and many others. The frequencies of the applications range from direct current (DC) to the highest one at which any electronic device can function [19].

Ferrite exhibits ferrimagnetic behavior which possesses unequal, anti-parallel ionic magnetic moments resulting in a net moment due to incomplete compensation. There are three classes of commercial ferrite in the industry, and each of the types has their own specific crystal structure. The three classes of the commercial ferrites are as follows:

## **5. Sintering of ferrite materials**

Ferrites are commonly fabricated via two major techniques: the conventional technique and the non-conventional technique. Through the conventional ceramic technique, the raw material powders are mixed and sintered at over 1000°C sintering temperature. This process allows interdiffusion of atoms in a pre-selected composition to form a mixed crystal. The other technique for preparing ferrites is the non-conventional technique. A non-conventional technique in a liquid medium may produce intermediate, finely divided mixed hydroxides or mixed organic salts, which would assist the subsequent diffusion process [20].

Various synthesis methods including the conventional and non-conventional techniques of ferrites preparation have been shown in **Table 1**. The table shows that the formation of ferrite through non-conventional technique could be produced by using the obtained fine powders at much lower sintering temperature. Yet most of the techniques still require sintering, although at relatively lower temperatures to produce a single phase material. The sintering temperature could be as low as 200°C [21], though displaying least performance of magnetic properties compare to much higher sintering temperatures. Highest sintering temperature is normally employed for synthesizing bulk ferrites via solid state reaction and has been shown to produce optimum magnetic properties (see **Table 1**). However, low sintering temperature is required for nano-sized materials as basic requirement since magnetic properties of the bulk materials differ drastically from the nano-sized materials.

Sintering in certain condition of atmosphere would form pure phase as has been observed in reference [22]. Furthermore, sintering atmosphere is also responsible in altering the magnetic properties and the effects can be observed in cation redistribution and oxygen deficiency [21]. The sintering time displays in **Table 1**, which shows that the non-conventional technique requires much shorter sintering time as compare to conventional technique. The right selection of sintering times would result in high densification and homogeneous materials which are largely important in magnetic materials since less densified materials result in hindrance to the domain wall movement, thus reducing the total magnetization. In addition, the

**Material** Ni0.3Cu0.2Zn0.5Fe

O2 4

Citrate precursor method

Fe(NO3

Zn(NO3

Cu (NO3

Ni(NO3

Citric acid

)2.6H

O2

atmospheres

Sintering temperature:

200,400 and 600°C

Sintering time: 2 h

NiFe

O2 4 Ni0.5Zn0.5Fe

O2 4 Ni0.266Zn0.66Cu0.09Fe1.968

O4-δ

Solid-state reaction

NiO

Sintering atmosphere:

Air

Sintering temperature:

1050°C

Sintering time: 1, 3, 5

and 7 h

ZnO

CuO

Fe

O2 3

method

Co-precipitation method

NiCl2

Sintering temperature:

Average crystallite sizes

Ms: 89.5 emu/g for

[26]

sample sintered at

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

900°C

increased from 38.4 to

42.2 nm with increasing

sintering temperatures

Average crystallite sizes

Ms: 245.1 emu/cm3

[23]

for sample sintered

for 7 h

Initial permeability,

µi, at 1 kHz: 2305

for sample sintered

for 7 h

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83

increased from 3.9 to

12.9 µm with increasing

sintering times

Density increased from

5.06 to 5.21 g/cm3 with

increasing sintering times

800 and 900°C

Sintering time: 3 h

ZnCl2

FeCl3

NaOH

Sol-gel auto-combustion

Fe(NO3

Ni(NO3

Citric acid

)2

⋅6H

O2

)3

⋅9H

O2

Sintering temperature:

Average grain sizes

Ms: ~50.0 emu/g for

[25]

sample sintered at

1100°C

increased from 0.1 to

2.2 µm with increasing

sintering temperatures

Density increased from

2.93 to 4.30 g/cm3 with

increasing sintering

temperatures

900, 1100, 1200 and

1300°C

Sintering time: 2 h

method

)2.3H

O2

)2. 4H

O2

)3.9H

O2

Sintering atmosphere:

Crystallite sizes increased

with increased sintering

temperatures for all

sintering atmosphere

ranging from 22.7 to 28.1

nm for air atmosphere,

8.8–24. 7 nm for argon

atmosphere and 10.1–31.6

nm for carbon monoxideambient atmosphere

air-, argon-, carbon

monoxide-ambient

**Synthesis method**

**Starting materials**

**Sintering conditions**

**Microstructure features**

**Optimum magnetic**

**Reference**

**properties**

Saturation

[21]

magnetization,

Ms: 69.1 emu/g for

sample sintered at

600°C under carbon

monoxide-ambient

atmosphere

Ferrite exhibits ferrimagnetic behavior which possesses unequal, anti-parallel ionic magnetic moments resulting in a net moment due to incomplete compensation. There are three classes of commercial ferrite in the industry, and each of the types has their own specific crystal struc-

i. Soft ferrite with spinel cubic structure, for example, nickel zinc ferrite and manganese zinc

ii. Soft ferrite with garnet structure, for example, yttrium-based garnets that are used in

iii. Hard ferrite with magnetoplumbite structure, for example, barium hexaferrite and strontium hexaferrite. The hexagonal ferrites develop high coercivity and are an important

Ferrites are commonly fabricated via two major techniques: the conventional technique and the non-conventional technique. Through the conventional ceramic technique, the raw material powders are mixed and sintered at over 1000°C sintering temperature. This process allows interdiffusion of atoms in a pre-selected composition to form a mixed crystal. The other technique for preparing ferrites is the non-conventional technique. A non-conventional technique in a liquid medium may produce intermediate, finely divided mixed hydroxides or

Various synthesis methods including the conventional and non-conventional techniques of ferrites preparation have been shown in **Table 1**. The table shows that the formation of ferrite through non-conventional technique could be produced by using the obtained fine powders at much lower sintering temperature. Yet most of the techniques still require sintering, although at relatively lower temperatures to produce a single phase material. The sintering temperature could be as low as 200°C [21], though displaying least performance of magnetic properties compare to much higher sintering temperatures. Highest sintering temperature is normally employed for synthesizing bulk ferrites via solid state reaction and has been shown to produce optimum magnetic properties (see **Table 1**). However, low sintering temperature is required for nano-sized materials as basic requirement since magnetic properties of the

Sintering in certain condition of atmosphere would form pure phase as has been observed in reference [22]. Furthermore, sintering atmosphere is also responsible in altering the magnetic properties and the effects can be observed in cation redistribution and oxygen deficiency [21]. The sintering time displays in **Table 1**, which shows that the non-conventional technique requires much shorter sintering time as compare to conventional technique. The right selection of sintering times would result in high densification and homogeneous materials which are largely important in magnetic materials since less densified materials result in hindrance to the domain wall movement, thus reducing the total magnetization. In addition, the

mixed organic salts, which would assist the subsequent diffusion process [20].

bulk materials differ drastically from the nano-sized materials.

ture. The three classes of the commercial ferrites are as follows:

member of the permanent magnet family.

**5. Sintering of ferrite materials**

ferrite.

microwave devices.

82 Sintering of Functional Materials

Zn0.35Ni0.57Co0.03Fe2.05

BaFe12O19

High-energy ball milling

BaCO3

Sintering atmosphere:

The highest measured

density is 4.88 g/cm3 for

the sample mechanically

alloyed for 3 h and

sintered at 1150°C

Air atmosphere

Sintering temperature:

800, 900 and 1150°C

Sintering time: 1 h

Fe

O2 3

O4 route

Chemical combustion

Metal nitrates

Sintering temperature:

Average grain sizes

increased from 0.61 to

0.94 µm with increasing

sintering temperatures

1050 and 1150°C

Sintering time: 1 h

Citric acid

**Synthesis method**

**Starting materials**

**Sintering conditions**

**Microstructure features**

**Optimum magnetic**

**Reference**

**properties**

Real part of the

[28]

initial permeability

increases with

increasing sintering

temperature from

85.2 to 209.7 at 10

kHz and from 90.4

to 238.1 at 1 MHz

The highest Ms

[29]

value of 63.57 emu/g

was measured

for the sample

mechanically

alloyed for 3 h and

sintered at 1150°C

The highest

coercivity, HC, value

is 5.31 kOe obtained

for the sample

milled for 9 h and

sintered at 900°C

Highest Ms of 49.4

[24]

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

emu/g for sample

sintered at 1200°C

with the largest

crystallite size

Maximum Hc

of13.82 Oe was

observed in

sample sintered at

http://dx.doi.org/10.5772/intechopen.68500

800°C (crystallite

size:36 nm) due to

transition of the

magnetic singledomain to magnetic

multi-domain

structure

85

Ni0.3Zn0.7Fe

O2 4

Sol–gel auto-combustion

Fe(NO3

Zn(NO3

Ni(NO3

)2.6H

O2

)

2. 4H

O2

Air atmosphere

Sintering temperature:

350, 400, 500, 600, 800,

1000 and 1200°C

Sintering time: 1 h

)3.9H

O2

Sintering atmosphere:

Average crystallite sizes

increased from 13.0 to

58 nm with increasing

sintering temperatures

method

Y3Fe

O5 12

Solid-state reaction

Fe

Y

O2 3

O2 3

Two different sintering

Sample sintered at

T1:1350°C, T2:1300°C for

18 h attained the highest

relative density (99.1%)

curves

Curve 1

Sintering

temperatures,

T1:1200–1450°C

Sintering time: none

Heating rate:10°C/min

Cooling rate: 2°C/min

Curve 2

Sintering

temperatures,

T1:1350°C, T2:1200 °C

T1:1350°C, T2:1300 °C

T1:1450 °C, T2:1300 °C

Sintering time at T2: 6,

12, 18 and 24 h

Heating rate:10°C/min

Cooling rate: 25°C/min

Sintering temperature:

Density increased from 4.0

Initial permeability,

[3]

µi, at 5 MHz

increased with

increased sintering

temperature with

maximum value

of ~112 for sample

sintered at 1050°C

Ms increased from

[4]

53.7 emu/g to 74.5

850, 900, 950, 1000 and

to 4.3 g/cm3

Average grain sizes

increased from 3.0 to

7.5 µm with increasing

sintering temperatures

1050°C

Sintering time: 2 h

NiCuZn ferrite

Commercial purchased

Commercial

purchased NiCuZn

ferrite

NiCuZn ferrite powder

and sintering

CoFe

O2 4

Citrate precursor method

Co(NO3

Fe(NO3

C

H6 O8 7

)2.9H O

2

)2.6H O

Sintering temperature:

Average grain sizes

increased from 90.0 to

100.0 nm with increasing

emu/g

sintering temperatures

900, 1000 and 1100°C

2

method

**Synthesis method**

**Starting materials**

**Sintering conditions**

**Microstructure features**

**Optimum magnetic**

**Reference**

**properties**

Ms: 27.4 emu/g for

[27]

84 Sintering of Functional Materials

sample sintered

at T1 = 1350°C,

T2:1300°C for 18 h

Y3Fe

O5 12

Solid-state

Fe

Y

O2 3

O2 3

**Conventional**

**Conventional sintering**

**sintering (CS)**

**(CS)**

Grain size: 3–5 µm

Density: 98% T.D

**Microwave sintering**

Sintering temperature:

1300°C

Sintering time: 6 h

Heating rate:2°C/min

**(MS)**

Grain size: 1.5 µm (900°C),

5–10 µm (1000°C)

Density: 96% T.D (900°C),

98% T.D (1000°C)

Cooling rate: 2°C/min

**Microwave sintering**

**(MS)**

Sintering temperature:

900°C and 1000°C

Sintering time: 20 min

(for 900°C) and 30 min

(for 1000°C)

Heating rate: 8°C/min

Cooling rate: 30°C/min

Sintering atmosphere:

Average grain size: ~1–2

µi increased from

[33]

2100 to 2450 with

increasing sintering

temperatures.

Ms: no sign of Ms

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

http://dx.doi.org/10.5772/intechopen.68500

87

even at 10000 Oe

magnetic field

Air atmosphere

µm

Density increased with

increasing sintering

temperatures from 4.64 to

4.80 g/cm3

Sintering temperature:

700, 800 and 950°C

Sintering time: 5 h

Mn1–xZnxFe

0.2, 0.4)

**Table 1.**

Summary of synthesis methods, sintering conditions, microstructure features and optimum magnetic properties of ferrites.

O2

4 (x = 0, 0.1,

Sol-gel combustion

Fe

O2 3

SrO

method

reaction method

and microwave sintering

**Synthesis method**

**Starting materials**

**Sintering conditions**

**Microstructure features**

**Optimum magnetic**

**Reference**

**properties**

**Conventional**

[32]

**sintering (CS)**

Ms: 25.42 emu/g

Coercive force, Hc:

25.36 Oe

**Microwave**

**sintering (MS)**

Ms: 14.60 emu/g

Hc: 34.82 Oe

Y3Fe

O5 12

Low

Y(NO3

Fe(NO3

Citric acid

)3

⋅9H

O2

)3

⋅6H

O2

Sintering atmosphere:

Optimum sintering

temperature of 1280°C

with highest value of ~5.08

g/cm3 that was about

98% of the XRD density

(5.17 g/cm3

)

Air atmosphere

Sintering

temperatures: 1220,

1240, 1280 and 1320°C

Sintering times: 3 h

Heating rate:10°C/min

Sintering atmosphere:

Sintering in nitrogen

Largest Ms value

[22]

of 90.02 emu/g

was observed in

sample sintered at

880°C in nitrogen

atmosphere with

5°C/min heating

rate

The Ms decreased

with increasing

heating rate from

90.02 to 80.60

produced pure

Mn0.49Zn0.48Fe2.06

while sintered in air or

mixture of air and nitrogen

contained oxides such as

O4 ferrite

air, mixture of nitrogen

and air, and nitrogen

atmospheres

Sintering temperature:

850, 880, 900 and 950°C

Heating rate: 5, 6, 7

Fe

O2 3, Mn

O2

Highest sintering

density of 4.82 g/cm3 and

homogeneous grain size

were found in sample

sintered at 880°C in

nitrogen atmosphere with

emu/g

5°C/min heating rate

Largest grains and highest

µi at 10 kHz

[31]

showed maximum

values of ~2825 for

density of 5.28 and 4.95 g/

cm3 were observed in Ni0.

35Cu0.05Zn0.60Fe

O2

at 950°C and Mg0.35Cu0.05

Zn0.60Fe

O2

900°C, respectively

4 sintered at

4 sintered

both Ni0.35Cu0.05

n0.60Fe

O2

at 950°C and Mg0

.35Cu0.05Zn0.60Fe

sintered at 900°C,

respectively

Highest Ms of ~129

and 88 emu/g for Ni

0.35Cu0.05Zn0.60Fe

sintered at 950°C

and Mg0.35Cu0.05Zn

0.60Fe

O2

900°C, respectively

4 sintered at

O2 4

O2 4

4 sintered

Z

Ni0.35Cu0.05Zn0.60Fe

Mg0.35Cu0.05Zn0.60Fe

O2 4

O2 4

Conventional mixed oxide

NiO

Sintering temperature:

850, 875, 900, 925, 950,

975 and 1000°C

Sintering time: 30 min

MgO

CuO

ZnO

Fe

O2 3

method and microwave

sintering

3 and ZnO

and 8°C/min

temperature solid state

reaction

Mn0.49Zn0.48Fe2.06

O4

Co-precipitation method

**Synthesis method**

**Starting materials**

**Sintering conditions**

**Microstructure features**

**Optimum magnetic**

**Reference**

**properties**

Ms: 13.8 mT for

[30]

sample sintered at

86 Sintering of Functional Materials

1280°C

**Table 1.** Summary of synthesis methods, sintering conditions, microstructure features and optimum magnetic properties of ferrites. crystallite size is observed to be significantly increased with prolong sintering time up to 7 h [23], thus enhancing the saturation magnetization of the material. This may be as a result of the improved crystallinity, which implying a better exchange interaction.

However, the effects of sintering conditions, particularly sintering temperatures and times, on magnetic properties evolution are not necessarily increased with increase in sintering temperature or sintering time. This is attributed to the resulting microstructure features such as abnormal grains and pores which are related to the decrease of density, thus decreasing the magnetic properties, mainly the volume magnetization and the magnetic induction. The non-linear relationship is also due to the characteristic of the transition from single-domain to multi-domain grains. This phenomenon is largely observed in coercivity value against particle or grain size of the magnetic material [24].

## **6. Results from experimental works on single-sample sintering (SSS) and multi-sample sintering (MSS) of NiZn Ferrites and YIG**

#### **6.1. Comparative study of single-sample and multi-sample sintering of NiZn ferrites**

Sintering temperatures increments from 600 to 1400°C increase the average grain size in both MSS and SSS as has been shown in **Table 2**, resulting from several processes. Those processes involve particles rearrangement and formation of dumbbell-liked structure between the particles contact points or known as the necking process. The grains are formed when the particles move closer during intermediate sintering stage as the sintering temperature goes higher. Finally, pores near or on the grain boundaries are gradually removed through the diffusion of vacancies associated by the pores along the grain boundaries, having only slight densification of the sample. The average grain size between the two different schemes shows small but significant difference. The striking difference in the microstructure is seen in Ni0.3Zn0.7Fe2 O4 sintered at 1100°C as shown in **Figure 3(a)** and **(b)**. The striking difference arises from the

different surface reactivities prior to the 1100°C sintering: for the MSS, the surface reactivity

as-milled powders. However, the SSS is subjected to several times of repeated sintering, thus reducing the surface reactivity of the material prior to the 1100°C sintering. The pores which

in density. Intragranular pores are trapped pores in the grains due to rapid grain growth and also probably due to zinc loss. The pores are known to be bad inclusions because they would pin down the domain wall, thus reducing the magnetization. However, no significant pores

could provide more time for the trapped pores to be removed; consequently, no significant

The focal question of what factors that subjected to different hysteresis shapes characteristic in both sintering treatments is of great interest. The shapes of the hysteresis loop are largely correlated with the microstructural features of the material, particularly the grains in the sample [34–40]. Besides, the disparities of the shapes also arise from the various grain shapes,

O4 are compacted from originally high-reactivity

sintered at: (a) 1100°C (MSS), (b) 1100°C (SSS), (c) 1400°C (MSS) and

sintered at 1400°C as shown in **Figure 3(c)** causing the decrease

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

http://dx.doi.org/10.5772/intechopen.68500

89

sintered at 1400°C. This is due to repeated sintering which

is high because the green bodies Ni0.3Zn0.7Fe2

**Figure 3.** FESEM micrographs for Ni0.3Zn0.7Fe2

O4

amount of intragranular porosity was observed.

O4

O4

exist in MSS Ni0.3Zn0.7Fe2

(d) 1400°C (SSS).

are observed in SSS Ni0.3Zn0.7Fe2

**Table 2.** Average grain size, experimental density, saturation induction and coercivity of Ni0.3Zn0.7Fe2 O4 for different sintering temperatures.

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering http://dx.doi.org/10.5772/intechopen.68500 89

crystallite size is observed to be significantly increased with prolong sintering time up to 7 h [23], thus enhancing the saturation magnetization of the material. This may be as a result of the

However, the effects of sintering conditions, particularly sintering temperatures and times, on magnetic properties evolution are not necessarily increased with increase in sintering temperature or sintering time. This is attributed to the resulting microstructure features such as abnormal grains and pores which are related to the decrease of density, thus decreasing the magnetic properties, mainly the volume magnetization and the magnetic induction. The non-linear relationship is also due to the characteristic of the transition from single-domain to multi-domain grains. This phenomenon is largely observed in coercivity value against par-

**6. Results from experimental works on single-sample sintering (SSS) and**

Sintering temperatures increments from 600 to 1400°C increase the average grain size in both MSS and SSS as has been shown in **Table 2**, resulting from several processes. Those processes involve particles rearrangement and formation of dumbbell-liked structure between the particles contact points or known as the necking process. The grains are formed when the particles move closer during intermediate sintering stage as the sintering temperature goes higher. Finally, pores near or on the grain boundaries are gradually removed through the diffusion of vacancies associated by the pores along the grain boundaries, having only slight densification of the sample. The average grain size between the two different schemes shows small but significant difference. The striking difference in the microstructure is seen in Ni0.3Zn0.7Fe2

sintered at 1100°C as shown in **Figure 3(a)** and **(b)**.  Average grain size, experimental density, saturation induction and coercivity of Ni0.3Zn0.7Fe2

O4

O4

for different

**6.1. Comparative study of single-sample and multi-sample sintering of NiZn ferrites**

improved crystallinity, which implying a better exchange interaction.

**multi-sample sintering (MSS) of NiZn Ferrites and YIG**

ticle or grain size of the magnetic material [24].

88 Sintering of Functional Materials

Average grain size (µm)

Experimental density (g/cm3

Saturation induction, Bs (Gauss)

Coercivity, Hc (Oe)

sintering temperatures.

)

**Figure 3.** FESEM micrographs for Ni0.3Zn0.7Fe2 O4 sintered at: (a) 1100°C (MSS), (b) 1100°C (SSS), (c) 1400°C (MSS) and (d) 1400°C (SSS).

different surface reactivities prior to the 1100°C sintering: for the MSS, the surface reactivity is high because the green bodies Ni0.3Zn0.7Fe2 O4 are compacted from originally high-reactivity as-milled powders. However, the SSS is subjected to several times of repeated sintering, thus reducing the surface reactivity of the material prior to the 1100°C sintering. The pores which exist in MSS Ni0.3Zn0.7Fe2 O4 sintered at 1400°C as shown in **Figure 3(c)** causing the decrease in density. Intragranular pores are trapped pores in the grains due to rapid grain growth and also probably due to zinc loss. The pores are known to be bad inclusions because they would pin down the domain wall, thus reducing the magnetization. However, no significant pores are observed in SSS Ni0.3Zn0.7Fe2 O4 sintered at 1400°C. This is due to repeated sintering which could provide more time for the trapped pores to be removed; consequently, no significant amount of intragranular porosity was observed.

The focal question of what factors that subjected to different hysteresis shapes characteristic in both sintering treatments is of great interest. The shapes of the hysteresis loop are largely correlated with the microstructural features of the material, particularly the grains in the sample [34–40]. Besides, the disparities of the shapes also arise from the various grain shapes, grain sizes, compositions, strains, and imperfections present in the sample. Maximum magnetic induction, Bs , of nickel zinc ferrite could range from 1000 to 3000 G [20, 41]. The experimental values of Bs are shown in **Table 2** range from 23.9 to 1076.0 G for MSS and from 23.4 to 930.7 G for SSS. The various ranges of Bs are subjected to the influence of several reasons in which categorizing the B-H hysteresis loops into several groups. The noticeably different B-H hysteresis loops are seen as three different shapes in both MSS and SSS. The loops are divided into three groups based on their magnetic behavior: strongly, moderately and weakly ferromagnetic, which are known to be strongly influenced by microstructural properties, domain states, and crystallinity of the samples. Ni0.3Zn0.7Fe2 O4 sintered from 600 to 800°C for both MSS and SSS as shown in **Figure 4(a)** and **(b)**, respectively, is classified as the first group with weakly ferromagnetic behavior. The shape of the hysteresis loops is affected by mixture phases of ferromagnetic and paramagnetic phase and also most likely by some superparamagnetic phase [42]. The significant amount of amorphous grain boundary volumes has contributed to the paramagnetic phase, which arises from the fine grain size of the samples [36, 39, 42]. In addition, a superparamagnetic phase is contributed by the nanosized grains. The shapes show a little hysteresis with narrowly bulging but linear-looking loops and have a very low Bs, indicating a very small amount of ferromagnetic phase. Due to the lower sintering temperature than the other two groups, the crystalline-phase percentage is small, while the amorphous-phase percentage is still significant. The grouping for the moderately ferromagnetic second group is slightly different between MSS and SSS where

(see **Figures 4(c)** and **(d)**). The difference between the two sintering schemes is due to the

group, shows a slanted sigmoid shape which is recognized to demonstrate moderate ferromagnetic behavior with negligible paramagnetic behavior since there is still remained a significant amount of the amorphous phase. The B-H loops of the MSS for this group have

higher ferromagnetic phase crystallinity and starting dominance of multi-domain magneti-

domain grains as shown in **Figure 5**. Consequently, the magnetization of the SSS samples is largely exhibiting via spin rotation, thus lowering the magnetization values than that of MSS samples which already possessing multi-domain grains though sintered at similar sintering

MSS sample already behaving as strongly ferromagnetic (third group), the SSS sample is still belong to the second group. The third group displays strongly ferromagnetic behavior (Bs

tering temperature. The sintering temperatures range from 1100 to 1400°C for the MSS and from 1200 to 1400°C for the SSS. The well-known erect, narrower and well-defined sigmoid shape has been observed for the third group of hysteresis loops. This strongly ferromagnetic behavior is contributed by very high crystallinity, high density with a minute amount of microstructural defects, and large size of grains, resulting from high sintering temperature. Therefore, the combinations of these particular parameters would allow domain walls

movement to become easier in the magnetization and demagnetization process.

**Figure 5.** Coercivity and grain size as a function of sintering temperatures for (a) MSS and (b) SSS of Ni0.3Zn0.7Fe2

) with a diminishing amorphous phase due to insignificant amount of amorphous grain

influence of microstructural properties (see **Figure 3**). Ni0.3Zn0.7Fe2

zation-demagnetization processes. However, in the SSS, Ni0.3Zn0.7Fe2

temperature. The result is clearly observed in Ni0.3Zn0.7Fe2

boundaries volume. This behavior is exhibited by Ni0.3Zn0.7Fe2

(Ms

sintered from 900 to 1000°C is for the MSS and from 900 to 1100°C is for SSS

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

) values but falling Hc values (see **Figure 5**) indicating, respectively,

O4

http://dx.doi.org/10.5772/intechopen.68500

O4

O4 sintered at 1100°C: while the

O4 sintered at much higher sin-

, which belongs to this

still exhibits single

O4 samples.

,

91

Ni0.3Zn0.7Fe2

Ms

O4

significantly higher Bs

**Figure 4.** B-H hysteresis loops of Ni0.3Zn0.7Fe2 O4 for first group of (a) MSS (b) SSS, second group of (c) MSS, (d) SSS, and third group of (e) MSS (f) SSS.

volumes has contributed to the paramagnetic phase, which arises from the fine grain size of the samples [36, 39, 42]. In addition, a superparamagnetic phase is contributed by the nanosized grains. The shapes show a little hysteresis with narrowly bulging but linear-looking loops and have a very low Bs, indicating a very small amount of ferromagnetic phase. Due to the lower sintering temperature than the other two groups, the crystalline-phase percentage is small, while the amorphous-phase percentage is still significant. The grouping for the moderately ferromagnetic second group is slightly different between MSS and SSS where Ni0.3Zn0.7Fe2 O4 sintered from 900 to 1000°C is for the MSS and from 900 to 1100°C is for SSS (see **Figures 4(c)** and **(d)**). The difference between the two sintering schemes is due to the influence of microstructural properties (see **Figure 3**). Ni0.3Zn0.7Fe2 O4 , which belongs to this group, shows a slanted sigmoid shape which is recognized to demonstrate moderate ferromagnetic behavior with negligible paramagnetic behavior since there is still remained a significant amount of the amorphous phase. The B-H loops of the MSS for this group have significantly higher Bs (Ms ) values but falling Hc values (see **Figure 5**) indicating, respectively, higher ferromagnetic phase crystallinity and starting dominance of multi-domain magnetization-demagnetization processes. However, in the SSS, Ni0.3Zn0.7Fe2 O4 still exhibits single domain grains as shown in **Figure 5**. Consequently, the magnetization of the SSS samples is largely exhibiting via spin rotation, thus lowering the magnetization values than that of MSS samples which already possessing multi-domain grains though sintered at similar sintering temperature. The result is clearly observed in Ni0.3Zn0.7Fe2 O4 sintered at 1100°C: while the MSS sample already behaving as strongly ferromagnetic (third group), the SSS sample is still belong to the second group. The third group displays strongly ferromagnetic behavior (Bs , Ms ) with a diminishing amorphous phase due to insignificant amount of amorphous grain boundaries volume. This behavior is exhibited by Ni0.3Zn0.7Fe2 O4 sintered at much higher sintering temperature. The sintering temperatures range from 1100 to 1400°C for the MSS and from 1200 to 1400°C for the SSS. The well-known erect, narrower and well-defined sigmoid shape has been observed for the third group of hysteresis loops. This strongly ferromagnetic behavior is contributed by very high crystallinity, high density with a minute amount of microstructural defects, and large size of grains, resulting from high sintering temperature. Therefore, the combinations of these particular parameters would allow domain walls movement to become easier in the magnetization and demagnetization process.

grain sizes, compositions, strains, and imperfections present in the sample. Maximum mag-

reasons in which categorizing the B-H hysteresis loops into several groups. The noticeably different B-H hysteresis loops are seen as three different shapes in both MSS and SSS. The loops are divided into three groups based on their magnetic behavior: strongly, moderately and weakly ferromagnetic, which are known to be strongly influenced by microstructural

to 800°C for both MSS and SSS as shown in **Figure 4(a)** and **(b)**, respectively, is classified as the first group with weakly ferromagnetic behavior. The shape of the hysteresis loops is affected by mixture phases of ferromagnetic and paramagnetic phase and also most likely by some superparamagnetic phase [42]. The significant amount of amorphous grain boundary

properties, domain states, and crystallinity of the samples. Ni0.3Zn0.7Fe2

, of nickel zinc ferrite could range from 1000 to 3000 G [20, 41]. The experi-

are shown in **Table 2** range from 23.9 to 1076.0 G for MSS and from

are subjected to the influence of several

O4 sintered from 600

netic induction, Bs

mental values of Bs

90 Sintering of Functional Materials

23.4 to 930.7 G for SSS. The various ranges of Bs

**Figure 4.** B-H hysteresis loops of Ni0.3Zn0.7Fe2

third group of (e) MSS (f) SSS.

O4

for first group of (a) MSS (b) SSS, second group of (c) MSS, (d) SSS, and

**Figure 5.** Coercivity and grain size as a function of sintering temperatures for (a) MSS and (b) SSS of Ni0.3Zn0.7Fe2 O4 samples.

The Hc values in **Figures 5(a)** and **(b)** are found to increase as the sintering temperature increased from 600 to 800°C, reaches a maximum value and decreased from 800 to 1400°C. Interestingly, the similar trend is showed in both MSS and SSS schemes, proving the transition of single-domain to multi-domain grains happened in the similar grain size range. The drop of Hc values in SSS has occurred earlier for Ni0.3Zn0.7Fe2 O4 having grain size of 0.19 µm as compared to MSS where it drops at 0.23 µm. Hc is probably the property most sensitive to porosity and grain size [20] nevertheless to the anisotropy field as well. Soft ferrites with nanometric grains exhibit a much higher Hc than samples having grain sizes of the order of few microns. An inversely proportional trend of Hc against grain size is observed for multidomain grains, which consisting of more domain walls. Therefore, the contribution of lower energy domain walls movement to demagnetization or demagnetization than that of domain rotation increases. Consequently, coarse grains are expected to display low Hc [43]. However, below a certain size, which the Hc reaches a maximum value, or known as the critical size, the grains are single-domain grains [44, 45]. The increasing values of coercivity for lower sintering (≤800°C) were due to size-shape anisotropy (necking phase in the microstructure) and magnetocrystalline anisotropy. For higher sintering temperatures (≥900°C), the grain size exceeded the critical grain size with the disappearing size-shape anisotropy but with remaining magnetocrystalline anisotropy. Magnetocrystalline anisotropy is reduced in larger grains by decreasing the internal stress and crystal anisotropy [46], helping in better domain walls movement, thus decreasing the Hc . Within this grain size range, the anisotropy and defects including pores govern the Hc values. **Figure 5(a)** and **(b)** greatly affirms the trend, giving a maximum Hc of 12.5 Oe at 0.23 µm and 11.5 Oe 0.19 µm for MSS and SSS, respectively. Therefore, the range of critical size for Ni0.3Zn0.7Fe2 O4 is approximately from 0.20 to 0.25 µm.

**Figures 6** and **7** present the real part of permeability with frequency dispersion from 1 MHz to 1.8 GHz for both sintering schemes. Generally, the permeability is related to two different magnetizing mechanisms which are spin rotational and domain wall movement. Normally, spin rotation occurs at higher frequency when domain is damped and could not follow the applied electromagnetic wave. According to Snoek's law [47], the relation between resonance frequency fR and the initial permeability *μ*i for Ni–Zn ferrites may be expressed as follows:

$$\mathbf{f\_n} = (1/\mu\_\mathrm{l}) \times \dots \times 10^s \,\mathrm{Hz} \tag{1}$$

**Figure 6.** Graph of real permeability, µ' against frequency for Ni0.3Zn0.7Fe2

**Figure 7.** Graph of real permeability, µ'against frequency for Ni0.3Zn0.7Fe2

O4 multi-sample sintering.

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

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93

O4 single-sample sintering.

This indicates that the lower the initial permeability values, the higher will be the frequency at which resonance phenomenon occurs. The value of real part of permeability for Ni0.3Zn0.7Fe2 O4 sintered at 600 and 700°C for both sintering schemes is independent of frequency in this measured frequency region (1 MHz–1.8 GHz), whereas Ni0.3Zn0.7Fe2 O4 sintered at 800°C only showed the dependency at about 100 MHz before reach the onset of resonance frequency. The resonance frequency is mostly observable in samples with lower sintering temperatures with the presence of single phase Ni0.3Zn0.7Fe2 O4 (800–1000°C) for both sintering schemes. The coarsened grains for Ni0.3Zn0.7Fe2 O4 sintered at 1100°C and above would lead to a ferromagnetic resonance at a lower frequency, in which, therefore, the resonance frequency could not be observed in the permeability spectra within the frequency region. At high frequencies, the domain walls cannot keep pace with the rapidly changing magnetic field, decreasing the value of real part of permeability. In powdered ferrites where each grain contains only a few domains, magnetization process occurs primarily by domain rotation and less by domain wall movement [48].

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering http://dx.doi.org/10.5772/intechopen.68500 93

**Figure 6.** Graph of real permeability, µ' against frequency for Ni0.3Zn0.7Fe2 O4 multi-sample sintering.

The Hc

92 Sintering of Functional Materials

drop of Hc

 values in **Figures 5(a)** and **(b)** are found to increase as the sintering temperature increased from 600 to 800°C, reaches a maximum value and decreased from 800 to 1400°C. Interestingly, the similar trend is showed in both MSS and SSS schemes, proving the transition of single-domain to multi-domain grains happened in the similar grain size range. The

to porosity and grain size [20] nevertheless to the anisotropy field as well. Soft ferrites with nanometric grains exhibit a much higher Hc than samples having grain sizes of the order of

domain grains, which consisting of more domain walls. Therefore, the contribution of lower energy domain walls movement to demagnetization or demagnetization than that of domain

below a certain size, which the Hc reaches a maximum value, or known as the critical size, the grains are single-domain grains [44, 45]. The increasing values of coercivity for lower sintering (≤800°C) were due to size-shape anisotropy (necking phase in the microstructure) and magnetocrystalline anisotropy. For higher sintering temperatures (≥900°C), the grain size exceeded the critical grain size with the disappearing size-shape anisotropy but with remaining magnetocrystalline anisotropy. Magnetocrystalline anisotropy is reduced in larger grains by decreasing the internal stress and crystal anisotropy [46], helping in better domain walls

including pores govern the Hc values. **Figure 5(a)** and **(b)** greatly affirms the trend, giving

**Figures 6** and **7** present the real part of permeability with frequency dispersion from 1 MHz to 1.8 GHz for both sintering schemes. Generally, the permeability is related to two different magnetizing mechanisms which are spin rotational and domain wall movement. Normally, spin rotation occurs at higher frequency when domain is damped and could not follow the applied electromagnetic wave. According to Snoek's law [47], the relation between resonance

This indicates that the lower the initial permeability values, the higher will be the frequency at which resonance phenomenon occurs. The value of real part of permeability for Ni0.3Zn0.7Fe2

sintered at 600 and 700°C for both sintering schemes is independent of frequency in this mea-

showed the dependency at about 100 MHz before reach the onset of resonance frequency. The resonance frequency is mostly observable in samples with lower sintering temperatures with

nance at a lower frequency, in which, therefore, the resonance frequency could not be observed in the permeability spectra within the frequency region. At high frequencies, the domain walls cannot keep pace with the rapidly changing magnetic field, decreasing the value of real part of permeability. In powdered ferrites where each grain contains only a few domains, magnetization process occurs primarily by domain rotation and less by domain wall movement [48].

R = (1/µi

O4

sured frequency region (1 MHz–1.8 GHz), whereas Ni0.3Zn0.7Fe2

O4

of 12.5 Oe at 0.23 µm and 11.5 Oe 0.19 µm for MSS and SSS, respectively.

O4

O4 having grain size of 0.19 µm

[43]. However,

O4

sintered at 800°C only

is probably the property most sensitive

against grain size is observed for multi-

is approximately from 0.20 to 0.25 µm.

. Within this grain size range, the anisotropy and defects

for Ni–Zn ferrites may be expressed as follows:

) × 3 × 109 Hz (1)

O4

(800–1000°C) for both sintering schemes. The coars-

sintered at 1100°C and above would lead to a ferromagnetic reso-

values in SSS has occurred earlier for Ni0.3Zn0.7Fe2

rotation increases. Consequently, coarse grains are expected to display low Hc

as compared to MSS where it drops at 0.23 µm. Hc

few microns. An inversely proportional trend of Hc

Therefore, the range of critical size for Ni0.3Zn0.7Fe2

frequency fR and the initial permeability *μ*i

f

the presence of single phase Ni0.3Zn0.7Fe2

ened grains for Ni0.3Zn0.7Fe2

movement, thus decreasing the Hc

a maximum Hc

**Figure 7.** Graph of real permeability, µ'against frequency for Ni0.3Zn0.7Fe2 O4 single-sample sintering.

The frequency stability for real permeability is varied from one group to another, which therefore varying the suitable applications for each group. The resonance frequency represents the high-frequency limit up to which the material can be used in a device. Ni0.3Zn0.7Fe2 O4 with strong ferromagnetic behavior is suitable for lower frequency application (less than 1 MHz) because of the frequency stability at lower frequency where resonance frequency is found to be lowered than measured frequency range. Some applications that are operating in the frequency range of 0.5–5 MHz are ferrite antennas for medium and long wave broadcast bands, power transformers, and cores for electromagnetic suppression. For moderate ferromagnetic behavior, Ni0.3Zn0.7Fe2 O4 sintered through MSS displays ferromagnetic resonance at frequency of 10.7 MHz (sintered at 900°C) and 3.9 MHz (sintered at 1000°C) with maximum real part of permeability value of 44.4 and 72.1, respectively, whereas for Ni0.3Zn0.7Fe2 O4 sintered through SSS, samples demonstrate ferromagnetic resonance at frequency of 20.1 MHz (sintered at 900°C), 6.31 MHz (sintered at 1000°C), 5.40 MHz (sintered at 1100°C) with maximum real part of permeability value of 37.4, 70.5 and 87.9, respectively. The materials could be used for the application of the solid core of inductors for resonant circuits or transformers operating in the approximate frequency range 2–20 MHz [41, 49], ferrite antennas for short wave broadcast bands, power transformers for the approximate frequency range 2–30 MHz and cores for electromagnetic interference suppression [41]. For weak ferromagnetic behavior, only Ni0.3Zn0.7Fe2 O4 sintered at 800°C for both sintering schemes displays the resonance phenomenon. The frequencies are stabled until 44.3 and 39.5 MHz for SSS and MSS, respectively, with maximum real part of permeability value of 8.64 and 4.42, respectively. Permeability with a value less than 12 is used for inductors and for resonant circuits operating at frequencies above 30 MHz and cores for electromagnetic interference suppression, whereas a much higher frequency than 1.8 GHz is needed to show resonance behavior in Ni0.3Zn0.7Fe2 O4 sintered at 600 and 700°C due to smaller grain size and lower magnetic mass in the Ni0.3Zn0.7Fe2 O4 .

The loss factor is observed to increase with a rise of the frequency from 1 MHz and attain the maximum value at a particular frequency and decreased with a further increase in frequency. The loss factor values increase with increasing sintering temperature in both MSS and SSS as shown in **Figures 8** and **9** for MSS and SSS, respectively. The frequency at which losses begin to increase due to the onset of resonance varies with the sintering temperatures from 2 to 100 MHz for both sintering schemes. As the sintering temperatures increase, the domain walls movement becomes easier in the larger grain, thus inducing larger eddy current. It is caused by the changing magnetic fields inside the sample which give rise to circulating currents inside the

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

not dominant due to its high electrical resistivity. The larger grain raises the number and size of magnetic domains which contribute to loss due to delay in domain wall motion. The losses in ferrites are associated with domain wall relaxation and rotational resonance. In higher frequency regions (>500 MHz), most of the domain walls are damped and become less important

O4

O4 multi-sample sintering.

O4 single-sample sintering.

, eddy current losses are

http://dx.doi.org/10.5772/intechopen.68500

95

sample, and hence to energy losses [50]. However, in Ni0.3Zn0.7Fe2

as spin rotational would continue to occur [20].

**Figure 8.** Graph of loss factor, µ" against frequency for Ni0.3Zn0.7Fe2

**Figure 9.** Graph of loss factor, µ" against frequency for Ni0.3Zn0.7Fe2

The complex permeability could also be classified into three different groups which are valuedifferentiated groups: strongly, moderately and weakly ferromagnetic behavior. By determining the critical size of single-domain to multi-domain grains through plot in **Figure 5**, it is found that Ni0.3Zn0.7Fe2 O4 with weakly ferromagnetic behavior contains less than 50% multi-domain grains (Ni0.3Zn0.7Fe2 O4 sintered from 600 to 800°C for MSS and SSS), whereas Ni0.3Zn0.7Fe2 O4 with the moderately ferromagnetic behavior possesses more than 50% multi-domain grains (Ni0.3Zn0.7Fe2 O4 sintered from 900 to 1000°C for MSS; Ni0.3Zn0.7Fe2 O4 sintered from 900 to 1100°C for SSS), and 100% of the grains are multi-domain grains (Ni0.3Zn0.7Fe2 O4 sintered from 1100 to 1400°C for MSS; Ni0.3Zn0.7Fe2 O4 sintered from 1200 to 1400°C for SSS) which are counted as strongly ferromagnetic behavior. Therefore, Ni0.3Zn0.7Fe2 O4 sintered below 800°C (with grain size less than 0.25 µm) is dominated by spin rotation, whereas Ni0.3Zn0.7Fe2 O4 sintered from 800°C upwards dominated by domain wall movement and spin rotation. The reason for the increase in permeability with sintering temperature is attributed to the increase of grain size and reduction of porosity, reducing the anisotropy arising from the demagnetizing fields outside of grains. Fewer number of the grain boundaries would be present in Ni0.3Zn0.7Fe2 O4 sintered at high temperatures, causing the existence of very mobile domain walls thus increasing the permeability value. Moreover, during grain growth, many pores would be removed, thus reducing the hindrance to the domain walls motion because pores provide stress concentration that may affect the magnetization's easy direction. However, the decrease in the real part of the permeability for Ni0.3Zn0.7Fe2 O4 sintered at 1400°C is attributed to zinc loss [47] and existence of pores (see **Figure 3**).

The loss factor is observed to increase with a rise of the frequency from 1 MHz and attain the maximum value at a particular frequency and decreased with a further increase in frequency. The loss factor values increase with increasing sintering temperature in both MSS and SSS as shown in **Figures 8** and **9** for MSS and SSS, respectively. The frequency at which losses begin to increase due to the onset of resonance varies with the sintering temperatures from 2 to 100 MHz for both sintering schemes. As the sintering temperatures increase, the domain walls movement becomes easier in the larger grain, thus inducing larger eddy current. It is caused by the changing magnetic fields inside the sample which give rise to circulating currents inside the sample, and hence to energy losses [50]. However, in Ni0.3Zn0.7Fe2 O4 , eddy current losses are not dominant due to its high electrical resistivity. The larger grain raises the number and size of magnetic domains which contribute to loss due to delay in domain wall motion. The losses in ferrites are associated with domain wall relaxation and rotational resonance. In higher frequency regions (>500 MHz), most of the domain walls are damped and become less important as spin rotational would continue to occur [20].

The frequency stability for real permeability is varied from one group to another, which therefore varying the suitable applications for each group. The resonance frequency represents the

strong ferromagnetic behavior is suitable for lower frequency application (less than 1 MHz) because of the frequency stability at lower frequency where resonance frequency is found to be lowered than measured frequency range. Some applications that are operating in the frequency range of 0.5–5 MHz are ferrite antennas for medium and long wave broadcast bands, power transformers, and cores for electromagnetic suppression. For moderate ferromagnetic

of 10.7 MHz (sintered at 900°C) and 3.9 MHz (sintered at 1000°C) with maximum real part of

SSS, samples demonstrate ferromagnetic resonance at frequency of 20.1 MHz (sintered at 900°C), 6.31 MHz (sintered at 1000°C), 5.40 MHz (sintered at 1100°C) with maximum real part of permeability value of 37.4, 70.5 and 87.9, respectively. The materials could be used for the application of the solid core of inductors for resonant circuits or transformers operating in the approximate frequency range 2–20 MHz [41, 49], ferrite antennas for short wave broadcast bands, power transformers for the approximate frequency range 2–30 MHz and cores for electromagnetic interference suppression [41]. For weak ferromagnetic behavior, only

enon. The frequencies are stabled until 44.3 and 39.5 MHz for SSS and MSS, respectively, with maximum real part of permeability value of 8.64 and 4.42, respectively. Permeability with a value less than 12 is used for inductors and for resonant circuits operating at frequencies above 30 MHz and cores for electromagnetic interference suppression, whereas a much higher fre-

The complex permeability could also be classified into three different groups which are valuedifferentiated groups: strongly, moderately and weakly ferromagnetic behavior. By determining the critical size of single-domain to multi-domain grains through plot in **Figure 5**, it is found that

sintered from 600 to 800°C for MSS and SSS), whereas Ni0.3Zn0.7Fe2

erately ferromagnetic behavior possesses more than 50% multi-domain grains (Ni0.3Zn0.7Fe2

O4

O4

by domain wall movement and spin rotation. The reason for the increase in permeability with sintering temperature is attributed to the increase of grain size and reduction of porosity, reducing the anisotropy arising from the demagnetizing fields outside of grains. Fewer number of the

existence of very mobile domain walls thus increasing the permeability value. Moreover, during grain growth, many pores would be removed, thus reducing the hindrance to the domain walls motion because pores provide stress concentration that may affect the magnetization's easy

direction. However, the decrease in the real part of the permeability for Ni0.3Zn0.7Fe2

1400°C is attributed to zinc loss [47] and existence of pores (see **Figure 3**).

sintered at 800°C for both sintering schemes displays the resonance phenom-

with weakly ferromagnetic behavior contains less than 50% multi-domain grains

O4

sintered below 800°C (with grain size less than 0.25 µm) is

sintered from 1200 to 1400°C for SSS) which are counted as strongly ferromagnetic

permeability value of 44.4 and 72.1, respectively, whereas for Ni0.3Zn0.7Fe2

quency than 1.8 GHz is needed to show resonance behavior in Ni0.3Zn0.7Fe2

and 700°C due to smaller grain size and lower magnetic mass in the Ni0.3Zn0.7Fe2

sintered through MSS displays ferromagnetic resonance at frequency

O4

O4

O4

O4

sintered at

sintered from 900 to 1100°C for SSS), and

sintered from 800°C upwards dominated

O4 sintered at high temperatures, causing the

sintered from 1100 to 1400°C for MSS;

O4 with

sintered through

sintered at 600

with the mod-

O4

O4 .

high-frequency limit up to which the material can be used in a device. Ni0.3Zn0.7Fe2

behavior, Ni0.3Zn0.7Fe2

94 Sintering of Functional Materials

Ni0.3Zn0.7Fe2

Ni0.3Zn0.7Fe2

(Ni0.3Zn0.7Fe2

Ni0.3Zn0.7Fe2

O4

O4

O4

behavior. Therefore, Ni0.3Zn0.7Fe2

O4

sintered from 900 to 1000°C for MSS; Ni0.3Zn0.7Fe2

dominated by spin rotation, whereas Ni0.3Zn0.7Fe2

grain boundaries would be present in Ni0.3Zn0.7Fe2

100% of the grains are multi-domain grains (Ni0.3Zn0.7Fe2

O4

O4

**Figure 8.** Graph of loss factor, µ" against frequency for Ni0.3Zn0.7Fe2 O4 multi-sample sintering.

**Figure 9.** Graph of loss factor, µ" against frequency for Ni0.3Zn0.7Fe2 O4 single-sample sintering.

#### **6.2. Comparative study of single-sample and multi-sample sintering of yttrium iron garnet**

A systematic track of microstructure-magnetic properties evolution of several polycrystalline Yttrium iron garnet (YIG) ferrite samples as a result of different sintering schemes was investigated in detail, focusing on the attendant occurrence of their dependency: an aspect seemingly ignored, hitherto in the garnet ferrite previous literatures for the past eight decades.

In order to prepare Yttrium iron garnet (YIG) ferrite sample, Fe2 O3 (Alfa Aesar, 99.945%) and Y2 O3 (Alfa Aesar, 99.99%) powders were weighed and mixed according to the stoichiometric proportions required in the final YIG samples based on the reaction:

$$\text{\textbulletY}\_2\text{O}\_3 + \text{5Fe}\_2\text{O}\_3 \to 2\text{Y}\_3\text{Fe}\_5\text{O}\_{12} \tag{2}$$

the saturation magnetization and magnetic induction with grain size, which was attributed to increase of crystallinity and demagnetizing field reduction in the grains. The variation in coercivity corresponded to the changes of anisotropy field within the samples due to grain size changes. Specifically, the starting appearance of room temperature ferromagnetic order suggested by the sigmoid-shaped B-H loops seems to be dependent on a sufficient number of large enough magnetic-domain containing grains formed in the microstructure. Viewed simultaneously, the B-H loops (appeared to be belonging to three groups with different magnetism-type dominance, respectively dependent on phase purity and distribution of grain size. The clearly tracked evolution of the hysteresis (**Figures 10** and **11**) and permeability component (**Figure 12**) strongly suggests that high reactivity grain surfaces and great-care human handling of the sample preparation process contributed to the startlingly clear microstructure-property evolution trends.

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

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97

**Figure 10.** B-H hysteresis loop for single-samples sintered at various temperatures. The circles inside the figure indicate

**Figure 11.** B-H hysteresis loop for multi-samples sintered at various temperatures. The circles inside the figure indicate

3 different groups of B-H curve evolution.

3 different groups of B-H curve evolution.

The powder then mechanically alloyed into nanosize via mechanical alloying technique. Two batches of samples were produced with different sintering scheme: SSS and MSS, each covering a range of low sintering temperature from 600°C up to high sintering temperature of 1400°C with 100°C increments. The samples were analyzed by using a LEO 912AB energy filter transmission electron microscope (TEM), Philips Expert PW3040 diffractometer operating at 40 kV/30 mA using Cu Kα radiation, scanning electron microscopy (SEM), MATS-2010S Static Hysteresis Graph at room temperature under applied magnetic fields 0–50 Oe (0–4000 A/m) and HP4291B Materials Impedance Analyzer at room temperature for their evolution stage in crystalline phases, microstructure, magnetic hysteresis-loop parameters, and magnetic permeability components, respectively.

With great experimental care, both the SSS and MSS batches yielded similar variation of microstructure-magnetic properties evolution (**Table 3**). The results showed an increasing tendency of

**Table 3.** Microstructural and magnetic parameters of single- and multi-sample sintering YIG samples with various sintering temperature variations.

the saturation magnetization and magnetic induction with grain size, which was attributed to increase of crystallinity and demagnetizing field reduction in the grains. The variation in coercivity corresponded to the changes of anisotropy field within the samples due to grain size changes. Specifically, the starting appearance of room temperature ferromagnetic order suggested by the sigmoid-shaped B-H loops seems to be dependent on a sufficient number of large enough magnetic-domain containing grains formed in the microstructure. Viewed simultaneously, the B-H loops (appeared to be belonging to three groups with different magnetism-type dominance, respectively dependent on phase purity and distribution of grain size. The clearly tracked evolution of the hysteresis (**Figures 10** and **11**) and permeability component (**Figure 12**) strongly suggests that high reactivity grain surfaces and great-care human handling of the sample preparation process contributed to the startlingly clear microstructure-property evolution trends.

**6.2. Comparative study of single-sample and multi-sample sintering of yttrium iron**

In order to prepare Yttrium iron garnet (YIG) ferrite sample, Fe2

3Y2

meability components, respectively.

**Grain size, (±0.01 μm)**

sintering temperature variations.

**Saturation induction, Bs (Gauss)**

proportions required in the final YIG samples based on the reaction:

O3 + 5Fe2

A systematic track of microstructure-magnetic properties evolution of several polycrystalline Yttrium iron garnet (YIG) ferrite samples as a result of different sintering schemes was investigated in detail, focusing on the attendant occurrence of their dependency: an aspect seemingly ignored, hitherto in the garnet ferrite previous literatures for the past eight decades.

(Alfa Aesar, 99.99%) powders were weighed and mixed according to the stoichiometric

O3 → 2Y3

The powder then mechanically alloyed into nanosize via mechanical alloying technique. Two batches of samples were produced with different sintering scheme: SSS and MSS, each covering a range of low sintering temperature from 600°C up to high sintering temperature of 1400°C with 100°C increments. The samples were analyzed by using a LEO 912AB energy filter transmission electron microscope (TEM), Philips Expert PW3040 diffractometer operating at 40 kV/30 mA using Cu Kα radiation, scanning electron microscopy (SEM), MATS-2010S Static Hysteresis Graph at room temperature under applied magnetic fields 0–50 Oe (0–4000 A/m) and HP4291B Materials Impedance Analyzer at room temperature for their evolution stage in crystalline phases, microstructure, magnetic hysteresis-loop parameters, and magnetic per-

With great experimental care, both the SSS and MSS batches yielded similar variation of microstructure-magnetic properties evolution (**Table 3**). The results showed an increasing tendency of

> **Coercivity, Hc (Oe)**

**Table 3.**  B-H hysteresis loop for single-samples sintered at various temperatures. The circles inside the figure indicate 3 different groups of B-H curve evolution.

**Figure 11.** B-H hysteresis loop for multi-samples sintered at various temperatures. The circles inside the figure indicate 3 different groups of B-H curve evolution.

**Author details**

Ismayadi Ismail1

**References**

\*, Idza Riati Ibrahim1

\*Address all correspondence to<EMAIL_ADDRESS>
Universiti Putra Malaysia, Serdang, Selangor, Malaysia

Universiti Malaysia Pahang, Kuantan, Pahang, Malaysia

2nd ed. Toh Tuck Lin: World Scientific. 2011

perature on magnetic properties of Y3

and magnetic properties of Ni0.55Zn0.45Fe2

http://dx.doi.org/10.1016/j.jmmm.2011.06.065

line Li0.5Fe2.5O4

jmmm.2007.10.016

Bi2 O3

Verlag Berlin Heidelberg. 2002

and Rodziah Nazlan2

Evolution of Magnetic Properties in Ferrites: Trends of Single-Sample and Multi-Sample Sintering

http://dx.doi.org/10.5772/intechopen.68500

99

1 Materials Synthesis and Characterisation Laboratory, Institute of Advanced Technology,

[1] Cao G, Wang Y. Nanostructures and Nanomaterials: Synthesis, Properties, and Applications.

[2] Winterer M. Nanocrystalline Ceramics Synthesis and Structure. New York: Springer-

[3] Yan YI, Ngo KDT, Hou D, Mu M, Mei Y, Lu G. Effect of sintering temperature on magnetic core-loss properties of a NiCuZn ferrite for high-frequency power converters. Journal of Electronic Materials. 2015;**44**(10):3788-3794. DOI: 10.1007/s11664-015-3836-z [4] Rana K, Thakur P, Sharma P, Tomar M, Gupta V, Thakur A. Improved structural and magnetic properties of cobalt nanoferrites : Influence of sintering temperature. Ceramics International. 2015;**41**:4492-4497. DOI: http://dx.doi.org/10.1016/j.ceramint.2014.11.143 [5] Huang CC, Hung YH, Huang JY, Kuo MF. Impact of stoichiometry and sintering tem-

Mnx

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ferrites. Materials Sciences and Applications.

2 Department of Material Technology, Faculty of Industrial Science and Technology,

**Figure 12.** Real permeability measured at room temperature in range 10 MHz to 1 GHz (a) single-sample sintering; (b) multi-samples sintering samples.

### **Acknowledgements**

We would like to dedicate this chapter and show our gratitude to the late Assoc. Prof. Dr. Mansor Hashim from Universiti Putra Malaysia, Malaysia for sharing his pearls of wisdom with us during the course of this research.

## **Author details**

Ismayadi Ismail1 \*, Idza Riati Ibrahim1 and Rodziah Nazlan2

\*Address all correspondence to<EMAIL_ADDRESS>
1 Materials Synthesis and Characterisation Laboratory, Institute of Advanced Technology, Universiti Putra Malaysia, Serdang, Selangor, Malaysia

2 Department of Material Technology, Faculty of Industrial Science and Technology, Universiti Malaysia Pahang, Kuantan, Pahang, Malaysia

## **References**

**Acknowledgements**

(b) multi-samples sintering samples.

98 Sintering of Functional Materials

with us during the course of this research.

We would like to dedicate this chapter and show our gratitude to the late Assoc. Prof. Dr. Mansor Hashim from Universiti Putra Malaysia, Malaysia for sharing his pearls of wisdom

**Figure 12.** Real permeability measured at room temperature in range 10 MHz to 1 GHz (a) single-sample sintering;

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